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GaN thin film growth on GaAs (0 0 1) by CBE and plasma-assisted MBE
Journal of Crystal Growth 243 (2002) 456–462
GaN thin film growth on GaAs (0 0 1) by CBE and plasma-assisted MBE Esther Kim, I. Rusakova, I. Berishev, A. Tempez, A. Bensaoula* Texas Center for Superconductivity and Advanced Materials, University of Houston, 724 Science and Research one Bldg., Houston, TX 77204-5004, USA Received 15 February 2002; accepted 11 June 2002 Communicated by R.S. Feigelson
Abstract GaN thin films were deposited on GaAs (0 0 1) substrates using chemical beam epitaxy (CBE) and plasma-assisted molecular beam epitaxy (MBE) methods. The effect of in situ substrate cleaning and pre-treatment on the interfacial structure was investigated. Time-of-flight mass spectroscopy of recoiled ions was utilized to measure in situ the surface composition. The microstructure of GaN films was studied with selected area electron diffraction, and both conventional and high-resolution transmission electron microscopy. It was found that simultaneous electron cyclotron resonance (ECR) plasma treatment using an Ar/N2 mixture increases the yield of a cubic phase in GaN films during CBE growth. Resulting films are less defective than mixed phase GaN films grown after pure N2 plasma nitridation. In the case of plasma-assisted MBE growth, a low-temperature annealing followed by N2 ECR nitridation process yields cubic GaN thin films while N2 ECR nitridation without a sample annealing to 6001C yields hexagonal phase GaN films. r 2002 Elsevier Science B.V. All rights reserved. PACS: 81.15.H Keywords: A1. Nucleation; A3. Chemical beam epitaxy; A3. Molecular beam epitaxy; B1. Nitrides; B2. Semiconducting gallium arsenide; B2. Semiconducting gallium nitride
1. Introduction GaN materials are technologically important for a variety of device applications [1]. They are ideal candidates for fabricating high-temperature electronic devices as well as emitters and detectors operating in the visible to UV spectrum. While *Corresponding author. Tel.: +1-713-743-3621; fax: +1713-747-7724. E-mail address: [email protected] (A. Bensaoula).
hexagonal GaN thin films have been realized by a number of growth technologies, much less has been published concerning the realization of the cubic GaN phase [2–4]. Cubic GaN, which is favored on zinc blende substrates such as GaAs, is believed to be more amenable to p-type doping and has been predicted to have superior optical and electrical properties [5]. Also, the demonstration of GaN on GaAs opens up the possibility of monolithically integrating traditional III–V compounds with the new nitride alloys such as GaAsN
0022-0248/02/$ - see front matterr 2002 Elsevier Science B.V. All rights reserved. PII: S 0 0 2 2 - 0 2 4 8 ( 0 2 ) 0 1 5 3 9 - 7
E. Kim et al. / Journal of Crystal Growth 243 (2002) 456–462
and GaPN [6,7]. Similarly to h-GaN deposition on sapphire, the large lattice mismatch between GaAs and c-GaN (20%) suggests the necessity of employing preliminary steps (nitridation and buffer layer) prior to the epilayer growth. Nitridation conditions (nature of active species, time), N/Ga flux ratio and substrate temperature are the critical parameters to be controlled in order to obtain a pure c-GaN phase [8]. Using ammonia plasma, Shimaoka et al. showed that the c-GaN content increases with reducing the ammonia flux [9]. With an RF plasma source, a thicker GaN layer is formed during the nitridation using N2/H2 or N2/NH3 mixtures than with a pure N2 plasma [10]. However, some researchers have attempted direct growth. Georgakilas et al. reported oriented cubic GaN/GaAs (1 0 0) including hexagonal domains without nitridation and using near stoichiometric N/Ga conditions [11]. For direct growth, near stoichiometric conditions are thought to be beneficial as they provide a complete initial GaN layer that prevents N damage of GaAs (responsible for h-GaN nucleation) [12,13]. In contrast, Trampert et al. have evidenced epitaxial cube-on-cube relationship of c-GaN/GaAs (1 0 0) in N-rich conditions [14,15]. In this case, N-rich conditions allow for high nuclei density and are combined to high growth temperature favoring low nucleation rate. Finally, the substrate temperature influence on the phase of GaN layers on GaAs has also been studied in the presence of arsenic or phosphorus overpressure [16–18]. In this paper, we report our results of GaN thin film deposition on GaAs (0 0 1) by the molecular beam methods, including chemical beam epitaxy (CBE) and electron cyclotron resonance (ECR) plasma-assisted molecular beam epitaxy (MBE). We have investigated how the initial pre-growth conditions influenced the crystal symmetry of GaN thin films for each growth method.
2. Experimental procedure Growth of GaN thin films was performed in a custom made ultra-high vacuum chamber equipped with two sources of gallium: a standard metallic Ga effusion cell and a triethylgallium
457
(TEG) injector, and with two sources of nitrogen: an ASTeXs compact ECR plasma source and an ammonia (NH3) gas injector. For most experiments, Si-doped epiready GaAs (0 0 1) substrates were annealed in vacuum at 4501C, without arsenic overpressure. This procedure has been previously determined to minimize surface carbon while keeping most of the surface arsenic [19]. After annealing, substrates were nitridated by exposing the surface to a ECR plasma source flux of either nitrogen or nitrogen–argon mixture, with the ECR operated at the maximum power of 200 W. In the case of nitrogen, the flow was 4 sccm and in the case of nitrogen–argon mixture it was 2–2 sccm. GaN buffer layers and final films were grown by ECR plasma-assisted metal organic MBE (MOMBE), CBE, and plasma-assisted MBE. The substrates were prepared in two different ways prior to CBE growth of GaN thin films. In the first case (i), a mixed Ar/N2 ECR plasma beam was applied for surface cleaning, which consisted of a 10 min bombardment at a sample temperature of 4501C. In the second case (ii), a pure N2 ECR plasma was utilized instead, other conditions remaining the same. For ECR plasma-assisted MBE growth, the substrate nitridation was performed in two different ways. In the first case (iii), GaAs substrates were slowly annealed at up to 6001C in vacuum, without arsenic overpressure, to get rid of the oxide layer. Subsequently, substrate nitridation was performed for 10 min on the resulting Ga-rich surface after the sample temperature was lowered at 5001C. During nitridation, the substrate temperature was increased to 6001C under N2-ECR plasma. The epilayer was grown at 6001C without a buffer layer. For the second case (iv), the substrate was nitridated at a sample temperature of 4201C without a previous vacuum annealing. Sample temperature was increased during nitridation to 6001C under N2-ECR plasma, and GaN layer deposition was started at 6001C, also without buffer layer. Time-of-flight mass spectroscopy of recoiled ions (TOF-MSRI) [20] was utilized to monitor in situ the surface composition at each growth step,
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while reflection high-energy electron diffraction was used to identify surface structure. Conventional bright field transmission electron microscopy (TEM) and high-resolution transmission electron microscopy (HRTEM), performed using a JEOL 2000 FX operated at 200 kV, were used to investigate phase composition and defects. Crosssectional samples for TEM study were prepared by a standard procedure of cutting, gluing, mechanical polishing, dimpling, and ion milling that was carried out with 4 keV Ar+ ions at liquid nitrogen temperature.
3. Results and discussion The results described in this section are summarized in Table 1. 3.1. Chemical beam epitaxy After the nitridation processes described earlier in this paper, strong nitrogen peaks were detected by TOF-MSRI indicating successful GaN nucleation on the GaAs substrate for both cases. Similarly, the carbon level was seen to decrease while that of oxygen remained constant. Following either step (i) Ar/N2 or step (ii) N2 only, a buffer layer was grown by MOMBE at 4501C using N2 ECR and TEG. We used MOMBE
for the buffer layer deposition as a result of our previous work, where we found that GaN does not easily nucleate on the nitridated sapphire substrate using just ammonia and TEG [21]. TOF-MSRI measurements on MOMBE grown buffer layers showed a surface containing a major carbon peak and a small amount of oxygen as impurities. The surface oxygen peak intensity dramatically increased while carbon was reduced when the buffer layers were annealed under ECR exposure at 7501C. The increase in oxygen content is either due to surface diffusion from the interface or the breakage of C–O bond from the adsorbed alkoxides on the surface and subsequent desorption of the carbon containing species. This can be supported as well by the lower strength of the Ga–C bond relative to that with either N or O (reported bond strength of Ga–N 97 kcal/mol [22], Ga–O 84 kcal/mol [23], and Ga–C in TEG 50 kcal/ mol [24]). TOF-MSRI from GaN thin films grown by ECR-MBE or by gas source MBE with NH3 in the same growth system does not show a surface oxygen peak. Therefore, the oxygen most likely originates from the reaction between the residual water molecules from the reactor or precursor reservoir with the TEG precursor forming a stable alkoxide. Fig. 1 shows conventional bright field TEM images recorded from typical specimens grown in conditions (i) and (ii). The HRTEM images in
Table 1 Growth technique/conditions and crystal quality results of the four types of GaN thin films studied Growth technique
Type number
Substrate treatment
Buffer layer
Growth
Crystalline structure
Orientation
Interface quality
CBE
(i)
Ar/N2 ECR, 10 min, 4501C
7801C
Single crystal cubic
[011]f8[011]s
Atomically sharp
(ii)
N2 ECR, 10 min, 4501C
N2 ECR +TEG, 4501C N2 ECR +TEG, 4501C
7801C
Mainly hexagonal+some cubic
[2 1% 1% 0]f8[011]s
Sharp
(iii)
6001C annealing N2 ECR from 5001C to 6001C No annealing N2 ECR from 4201C to 6001C
None
6001C
Single crystal cubic
[011]f8[011]s
Rough
None
6001C
Single crystal hexagonal+some cubic
[2 1% 1% 0]f8[011]s
Atomically sharp
ECRMBE
(iv)
Subscripts f and s mean film and substrate, respectively.
E. Kim et al. / Journal of Crystal Growth 243 (2002) 456–462
459
Fig. 1. Conventional bright field TEM images of GaN thin films grown by CBE. Surface nitridation was performed on GaAs (0 0 1) by: (a) case (i)—N2/Ar ECR plasma, and (b) case (ii)—N2 ECR plasma before the epilayers were grown.
Fig. 2 provide detailed information on the atomic structure of films and interfaces between substrates and films. Inserts to Fig. 2 show the selected area electron diffraction (SAED) patterns taken from the corresponding area. Conventional TEM images reveal that GaN thin films in both cases have uniform thicknesses of about 250 nm over the entire specimen and smooth interfaces with the substrates. The films consist of aligned domains formed during the film growth. Differences between these two types of specimens are observed on SAED patterns. A typical type-(i) sample is a single crystal and has
an FCC structure like that of a GaAs substrate. The cubic GaN film is well aligned with the substrate lattice. SAED patterns recorded from the interface shows that both substrate and film have a [0 1 1] zone axis (Fig. 2a). In the case of type-(ii) samples, the SAED pattern presents mainly hexagonal symmetry but there are some reflections that could have resulted from an FCC type of structure as well (Fig. 2b). The crystal orientation of the substrate and the film is [0 1 1]substrate8[2 1% 1% 0]film. In addition, (ii) GaN films are slightly misaligned to the substrate (B2–31).
460
E. Kim et al. / Journal of Crystal Growth 243 (2002) 456–462
Fig. 2. HRTEM images of GaN films grown by CBE on GaAs (0 0 1) substrates. SAED patterns from thin films, interfaces and substrates are shown in insets. GaAs (0 0 1) substrates are ECR-plasma bombarded by an Ar/N2 mixture (a), and N2 only (b).
Conventional and high-resolution TEM images also show the strain contrast along the interfaces, which is attributed to the large lattice mismatch between substrate and thin film. The stress is better accommodated for the type-(i) sample (i) than for the type-(ii) sample where the thin film is a mixture of hexagonal and cubic phases. In the case of the cubic GaN film, the interface is atomically sharp (Fig. 2a). A high density of growth twins is observed on both conventional and high-resolution TEM images from the two types of films. Twinning is
caused by high lattice mismatch between GaN and GaAs that was elastically accommodated. There are also stacking faults and dislocations that are seen on HRTEM images. The existence of planar defects in GaN films results in streaks on the SAED patterns. 3.2. ECR plasma-assisted MBE A conventional TEM image and a SAED pattern of a type-(iii) sample film are shown in Fig. 3a. The GaN thin film is a single crystal with
E. Kim et al. / Journal of Crystal Growth 243 (2002) 456–462
461
Fig. 3. Conventional bright field TEM images of GaN thin films grown by ECR plasma-assisted MBE on GaAs (0 0 1). SAED patterns from the films and interfaces are shown in insets: (a) case (iii)—a cubic phase GaN layer was grown when the substrate was annealed in vacuum at 6001C then N2-ECR nitridated from 5001C, and (b) case (iv)—a hexagonal phase GaN layer was obtained by direct N2ECR nitridation at 5001C without a previous vacuum annealing.
an FCC structure, and it is well aligned to the substrate ([0 1 1]film8[0 1 1]substrate). The interface between thin film and substrate is not atomically sharp after this deposition process. TEM study of the type-(iv) sample reveals that here again, the GaN thin film is a single crystal and well aligned to the substrate, but has a hexagonal structure (Fig. 3b). The [2 1% 1% 0] direction of the film is parallel to the [0 1 1] direction of the substrate. This film–substrate interface is atomically sharp. It is well known that vacuum annealing of GaAs substrate without As over-pressure leads to metal Ga accumulation on the surface, which we believe is more severe in the (iii) case. The subsequent ECR plasma nitridation nucleates a very rough thin GaN layer, which, however, consists almost entirely of cubic phase GaN. The cubic GaN template leads to a cubic phase for the subsequent
GaN epilayer. A type-(iv) sample, which was not annealed in the vacuum, has less metal Ga and contains more oxygen on the surface before nitridation begins. The lack of cubic phase templates on the surface covered with amorphous oxides resulted in hexagonal phase GaN template and epilayer growth. As can also be seen from Fig. 3a and b, in both cases GaN films have a high density of planar defects (growth twins and stacking faults) and linear defects that are caused by lattice mismatch between substrate and film.
4. Conclusion We show that pure cubic GaN was successfully deposited by NH3-based CBE for the first time to our knowledge. The initial substrate treatment
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seems to greatly affect the crystal symmetry of following GaN films. For GaN growth by CBE, the use of an argon/nitrogen mixture instead of pure nitrogen for the nitridation step increases yield of cubic phase GaN thin films. These films have better structural quality than those deposited after plain N2-ECR nitridation, which yielded a mixed phase film. By carefully choosing the substrate nitridation condition and plasma composition, one can better control the GaN layer phase composition during its growth on the GaAs substrate. For the ECR MBE growth, cubic phase GaN layer was achieved only after high-temperature substrate annealing.
Acknowledgements This work is supported by fund from a NASA cooperative agreement, No. NCC8-127, to the Space Vacuum Epitaxy Center. I.R. would like to acknowledge support from the Texas Center for Superconductivity (TCSUH) at the University of Houston.
References [1] S.N. Mohammad, A.A. Salvador, H. Morkoc-, Proc. IEEE 83 (1995) 1306. [2] H. Okumura, K. Balakrishnan, G. Feuillet, K. Ohta, H. Hamaguchi, S. Chichibu, Y. Ishida, S. Yoshda, Mater. Res. Soc. Symp. Proc. 449 (1997) 435. [3] S. Ruvimov, Z. Liliental-Weber, J. Washburn, T.J. Drummond, M. Hafish, S.R. Lee, Appl. Phys. Lett. 71 (1997) 1931. [4] K.H. Ploog, O. Brandt, H. Yang, B. Yang, A. Trampert, J. Vac. Sci. Technol. B 16 (1998) 2229.
[5] M.E. Lin, G. Xue, G.L. Zhou, J.E. Greene, H. Morkoc, Appl. Phys. Lett. 63 (1993) 932. [6] G. Pozina, I. Ivanov, B. Monemar, J.V. Thordson, T.G. Andersson, J. Appl. Phys. 84 (1998) 3830. [7] L. Bellaiche, S.-H. Wei, A. Zunger, Appl. Phys. Lett. 70 (1997) 3558. [8] H. Chen, H. Liu, Z. Li, S. Liu, Q. Huang, J. Zhou, Y. Wang, J. Crystal Growth 201/202 (1999) 336. [9] G. Shimaoka, T. Aoki, Y. Nakanishi, Y. Hatanaka, T. Udagawa, Appl. Surf. Sci. 175–176 (2001) 436. [10] M. Losurdo, P. Capezzuto, G. Bruno, E.A. Irene, Phys. Rev. B 58 (1998) 15878. [11] A. Georgakilas, K. Amimer, P. Tzanetakis, Z. Hatzopoulos, M. Cengher, B. Pecz, Zs. Czigany, L. Toth, M.V. Baidakova, A.V. Sakharov, V.Yu. Davydov, J. Crystal Growth 227–228 (2001) 410. . J.V. Thordson, J.R. Gunnarson, Q.X. Zhao, L. [12] O. Zsebok, Ilver, T.G. Andersson, J. Appl. Phys. 89 (2001) 3662. . J.V. Thordson, T.G. Andersson, Jpn. J. Appl. [13] O. Zsebok, Phys. 40 (2001) 472. [14] A. Trampert, O. Brandt, H. Yang, K.H. Ploog, Appl. Phys. Lett. 70 (1997) 583. [15] O. Brandt, H. Yang, A. Trampert, M. Wassermeier, K.H. Ploog, Appl. Phys. Lett. 71 (1997) 473. [16] T.S. Cheng, C. Jemkins, S.E. Hooper, C.T. Foxon, J.W. Orton, D.E. Lacklison, Appl. Phys. Lett. 66 (1995) 1509. [17] Y. Zhao, C.W. Tu, I.-T. Bae, T.-Y. Seong, Appl. Phys. Lett. 74 (1999) 3182. [18] A. Hashimoto, T. Motizuki, H. Wada, A. Masuda, A. Yamamoto, J. Crystal Growth 201/202 (1999) 336. [19] A. Bensaoula, W.T. Taferner, E. Kim, A. Bousetta, J. Crystal Growth 164 (1996) 185. [20] A.R. Krauss, O. Auciello, J.A. Schultz, MRS Bull. 20 (5) (1995) 18. [21] E. Kim, I.E. Berishev, A. Bensaoula, J.A. Shultz, K. Waters, S. Lee, S.S. Perry, Appl. Phys. Lett. 71 (1997) 3072. [22] L.A. DeLouise, J. Vac. Sci. Technol. A 10 (1992) 1637. [23] A. Neubert, K.F. Zombov, J. Chem. Soc. Faraday Trans. 1 70 (1974) 2219. [24] D.F. McMillen, D.M. Golden, Ann. Rev. Phys. Chem. 33 (1982) 493.
GaN thin film growth on GaAs (0 0 1) by CBE and plasma-assisted MBE Esther Kim, I. Rusakova, I. Berishev, A. Tempez, A. Bensaoula* Texas Center for Superconductivity and Advanced Materials, University of Houston, 724 Science and Research one Bldg., Houston, TX 77204-5004, USA Received 15 February 2002; accepted 11 June 2002 Communicated by R.S. Feigelson
Abstract GaN thin films were deposited on GaAs (0 0 1) substrates using chemical beam epitaxy (CBE) and plasma-assisted molecular beam epitaxy (MBE) methods. The effect of in situ substrate cleaning and pre-treatment on the interfacial structure was investigated. Time-of-flight mass spectroscopy of recoiled ions was utilized to measure in situ the surface composition. The microstructure of GaN films was studied with selected area electron diffraction, and both conventional and high-resolution transmission electron microscopy. It was found that simultaneous electron cyclotron resonance (ECR) plasma treatment using an Ar/N2 mixture increases the yield of a cubic phase in GaN films during CBE growth. Resulting films are less defective than mixed phase GaN films grown after pure N2 plasma nitridation. In the case of plasma-assisted MBE growth, a low-temperature annealing followed by N2 ECR nitridation process yields cubic GaN thin films while N2 ECR nitridation without a sample annealing to 6001C yields hexagonal phase GaN films. r 2002 Elsevier Science B.V. All rights reserved. PACS: 81.15.H Keywords: A1. Nucleation; A3. Chemical beam epitaxy; A3. Molecular beam epitaxy; B1. Nitrides; B2. Semiconducting gallium arsenide; B2. Semiconducting gallium nitride
1. Introduction GaN materials are technologically important for a variety of device applications [1]. They are ideal candidates for fabricating high-temperature electronic devices as well as emitters and detectors operating in the visible to UV spectrum. While *Corresponding author. Tel.: +1-713-743-3621; fax: +1713-747-7724. E-mail address: [email protected] (A. Bensaoula).
hexagonal GaN thin films have been realized by a number of growth technologies, much less has been published concerning the realization of the cubic GaN phase [2–4]. Cubic GaN, which is favored on zinc blende substrates such as GaAs, is believed to be more amenable to p-type doping and has been predicted to have superior optical and electrical properties [5]. Also, the demonstration of GaN on GaAs opens up the possibility of monolithically integrating traditional III–V compounds with the new nitride alloys such as GaAsN
0022-0248/02/$ - see front matterr 2002 Elsevier Science B.V. All rights reserved. PII: S 0 0 2 2 - 0 2 4 8 ( 0 2 ) 0 1 5 3 9 - 7
E. Kim et al. / Journal of Crystal Growth 243 (2002) 456–462
and GaPN [6,7]. Similarly to h-GaN deposition on sapphire, the large lattice mismatch between GaAs and c-GaN (20%) suggests the necessity of employing preliminary steps (nitridation and buffer layer) prior to the epilayer growth. Nitridation conditions (nature of active species, time), N/Ga flux ratio and substrate temperature are the critical parameters to be controlled in order to obtain a pure c-GaN phase [8]. Using ammonia plasma, Shimaoka et al. showed that the c-GaN content increases with reducing the ammonia flux [9]. With an RF plasma source, a thicker GaN layer is formed during the nitridation using N2/H2 or N2/NH3 mixtures than with a pure N2 plasma [10]. However, some researchers have attempted direct growth. Georgakilas et al. reported oriented cubic GaN/GaAs (1 0 0) including hexagonal domains without nitridation and using near stoichiometric N/Ga conditions [11]. For direct growth, near stoichiometric conditions are thought to be beneficial as they provide a complete initial GaN layer that prevents N damage of GaAs (responsible for h-GaN nucleation) [12,13]. In contrast, Trampert et al. have evidenced epitaxial cube-on-cube relationship of c-GaN/GaAs (1 0 0) in N-rich conditions [14,15]. In this case, N-rich conditions allow for high nuclei density and are combined to high growth temperature favoring low nucleation rate. Finally, the substrate temperature influence on the phase of GaN layers on GaAs has also been studied in the presence of arsenic or phosphorus overpressure [16–18]. In this paper, we report our results of GaN thin film deposition on GaAs (0 0 1) by the molecular beam methods, including chemical beam epitaxy (CBE) and electron cyclotron resonance (ECR) plasma-assisted molecular beam epitaxy (MBE). We have investigated how the initial pre-growth conditions influenced the crystal symmetry of GaN thin films for each growth method.
2. Experimental procedure Growth of GaN thin films was performed in a custom made ultra-high vacuum chamber equipped with two sources of gallium: a standard metallic Ga effusion cell and a triethylgallium
457
(TEG) injector, and with two sources of nitrogen: an ASTeXs compact ECR plasma source and an ammonia (NH3) gas injector. For most experiments, Si-doped epiready GaAs (0 0 1) substrates were annealed in vacuum at 4501C, without arsenic overpressure. This procedure has been previously determined to minimize surface carbon while keeping most of the surface arsenic [19]. After annealing, substrates were nitridated by exposing the surface to a ECR plasma source flux of either nitrogen or nitrogen–argon mixture, with the ECR operated at the maximum power of 200 W. In the case of nitrogen, the flow was 4 sccm and in the case of nitrogen–argon mixture it was 2–2 sccm. GaN buffer layers and final films were grown by ECR plasma-assisted metal organic MBE (MOMBE), CBE, and plasma-assisted MBE. The substrates were prepared in two different ways prior to CBE growth of GaN thin films. In the first case (i), a mixed Ar/N2 ECR plasma beam was applied for surface cleaning, which consisted of a 10 min bombardment at a sample temperature of 4501C. In the second case (ii), a pure N2 ECR plasma was utilized instead, other conditions remaining the same. For ECR plasma-assisted MBE growth, the substrate nitridation was performed in two different ways. In the first case (iii), GaAs substrates were slowly annealed at up to 6001C in vacuum, without arsenic overpressure, to get rid of the oxide layer. Subsequently, substrate nitridation was performed for 10 min on the resulting Ga-rich surface after the sample temperature was lowered at 5001C. During nitridation, the substrate temperature was increased to 6001C under N2-ECR plasma. The epilayer was grown at 6001C without a buffer layer. For the second case (iv), the substrate was nitridated at a sample temperature of 4201C without a previous vacuum annealing. Sample temperature was increased during nitridation to 6001C under N2-ECR plasma, and GaN layer deposition was started at 6001C, also without buffer layer. Time-of-flight mass spectroscopy of recoiled ions (TOF-MSRI) [20] was utilized to monitor in situ the surface composition at each growth step,
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while reflection high-energy electron diffraction was used to identify surface structure. Conventional bright field transmission electron microscopy (TEM) and high-resolution transmission electron microscopy (HRTEM), performed using a JEOL 2000 FX operated at 200 kV, were used to investigate phase composition and defects. Crosssectional samples for TEM study were prepared by a standard procedure of cutting, gluing, mechanical polishing, dimpling, and ion milling that was carried out with 4 keV Ar+ ions at liquid nitrogen temperature.
3. Results and discussion The results described in this section are summarized in Table 1. 3.1. Chemical beam epitaxy After the nitridation processes described earlier in this paper, strong nitrogen peaks were detected by TOF-MSRI indicating successful GaN nucleation on the GaAs substrate for both cases. Similarly, the carbon level was seen to decrease while that of oxygen remained constant. Following either step (i) Ar/N2 or step (ii) N2 only, a buffer layer was grown by MOMBE at 4501C using N2 ECR and TEG. We used MOMBE
for the buffer layer deposition as a result of our previous work, where we found that GaN does not easily nucleate on the nitridated sapphire substrate using just ammonia and TEG [21]. TOF-MSRI measurements on MOMBE grown buffer layers showed a surface containing a major carbon peak and a small amount of oxygen as impurities. The surface oxygen peak intensity dramatically increased while carbon was reduced when the buffer layers were annealed under ECR exposure at 7501C. The increase in oxygen content is either due to surface diffusion from the interface or the breakage of C–O bond from the adsorbed alkoxides on the surface and subsequent desorption of the carbon containing species. This can be supported as well by the lower strength of the Ga–C bond relative to that with either N or O (reported bond strength of Ga–N 97 kcal/mol [22], Ga–O 84 kcal/mol [23], and Ga–C in TEG 50 kcal/ mol [24]). TOF-MSRI from GaN thin films grown by ECR-MBE or by gas source MBE with NH3 in the same growth system does not show a surface oxygen peak. Therefore, the oxygen most likely originates from the reaction between the residual water molecules from the reactor or precursor reservoir with the TEG precursor forming a stable alkoxide. Fig. 1 shows conventional bright field TEM images recorded from typical specimens grown in conditions (i) and (ii). The HRTEM images in
Table 1 Growth technique/conditions and crystal quality results of the four types of GaN thin films studied Growth technique
Type number
Substrate treatment
Buffer layer
Growth
Crystalline structure
Orientation
Interface quality
CBE
(i)
Ar/N2 ECR, 10 min, 4501C
7801C
Single crystal cubic
[011]f8[011]s
Atomically sharp
(ii)
N2 ECR, 10 min, 4501C
N2 ECR +TEG, 4501C N2 ECR +TEG, 4501C
7801C
Mainly hexagonal+some cubic
[2 1% 1% 0]f8[011]s
Sharp
(iii)
6001C annealing N2 ECR from 5001C to 6001C No annealing N2 ECR from 4201C to 6001C
None
6001C
Single crystal cubic
[011]f8[011]s
Rough
None
6001C
Single crystal hexagonal+some cubic
[2 1% 1% 0]f8[011]s
Atomically sharp
ECRMBE
(iv)
Subscripts f and s mean film and substrate, respectively.
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Fig. 1. Conventional bright field TEM images of GaN thin films grown by CBE. Surface nitridation was performed on GaAs (0 0 1) by: (a) case (i)—N2/Ar ECR plasma, and (b) case (ii)—N2 ECR plasma before the epilayers were grown.
Fig. 2 provide detailed information on the atomic structure of films and interfaces between substrates and films. Inserts to Fig. 2 show the selected area electron diffraction (SAED) patterns taken from the corresponding area. Conventional TEM images reveal that GaN thin films in both cases have uniform thicknesses of about 250 nm over the entire specimen and smooth interfaces with the substrates. The films consist of aligned domains formed during the film growth. Differences between these two types of specimens are observed on SAED patterns. A typical type-(i) sample is a single crystal and has
an FCC structure like that of a GaAs substrate. The cubic GaN film is well aligned with the substrate lattice. SAED patterns recorded from the interface shows that both substrate and film have a [0 1 1] zone axis (Fig. 2a). In the case of type-(ii) samples, the SAED pattern presents mainly hexagonal symmetry but there are some reflections that could have resulted from an FCC type of structure as well (Fig. 2b). The crystal orientation of the substrate and the film is [0 1 1]substrate8[2 1% 1% 0]film. In addition, (ii) GaN films are slightly misaligned to the substrate (B2–31).
460
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Fig. 2. HRTEM images of GaN films grown by CBE on GaAs (0 0 1) substrates. SAED patterns from thin films, interfaces and substrates are shown in insets. GaAs (0 0 1) substrates are ECR-plasma bombarded by an Ar/N2 mixture (a), and N2 only (b).
Conventional and high-resolution TEM images also show the strain contrast along the interfaces, which is attributed to the large lattice mismatch between substrate and thin film. The stress is better accommodated for the type-(i) sample (i) than for the type-(ii) sample where the thin film is a mixture of hexagonal and cubic phases. In the case of the cubic GaN film, the interface is atomically sharp (Fig. 2a). A high density of growth twins is observed on both conventional and high-resolution TEM images from the two types of films. Twinning is
caused by high lattice mismatch between GaN and GaAs that was elastically accommodated. There are also stacking faults and dislocations that are seen on HRTEM images. The existence of planar defects in GaN films results in streaks on the SAED patterns. 3.2. ECR plasma-assisted MBE A conventional TEM image and a SAED pattern of a type-(iii) sample film are shown in Fig. 3a. The GaN thin film is a single crystal with
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Fig. 3. Conventional bright field TEM images of GaN thin films grown by ECR plasma-assisted MBE on GaAs (0 0 1). SAED patterns from the films and interfaces are shown in insets: (a) case (iii)—a cubic phase GaN layer was grown when the substrate was annealed in vacuum at 6001C then N2-ECR nitridated from 5001C, and (b) case (iv)—a hexagonal phase GaN layer was obtained by direct N2ECR nitridation at 5001C without a previous vacuum annealing.
an FCC structure, and it is well aligned to the substrate ([0 1 1]film8[0 1 1]substrate). The interface between thin film and substrate is not atomically sharp after this deposition process. TEM study of the type-(iv) sample reveals that here again, the GaN thin film is a single crystal and well aligned to the substrate, but has a hexagonal structure (Fig. 3b). The [2 1% 1% 0] direction of the film is parallel to the [0 1 1] direction of the substrate. This film–substrate interface is atomically sharp. It is well known that vacuum annealing of GaAs substrate without As over-pressure leads to metal Ga accumulation on the surface, which we believe is more severe in the (iii) case. The subsequent ECR plasma nitridation nucleates a very rough thin GaN layer, which, however, consists almost entirely of cubic phase GaN. The cubic GaN template leads to a cubic phase for the subsequent
GaN epilayer. A type-(iv) sample, which was not annealed in the vacuum, has less metal Ga and contains more oxygen on the surface before nitridation begins. The lack of cubic phase templates on the surface covered with amorphous oxides resulted in hexagonal phase GaN template and epilayer growth. As can also be seen from Fig. 3a and b, in both cases GaN films have a high density of planar defects (growth twins and stacking faults) and linear defects that are caused by lattice mismatch between substrate and film.
4. Conclusion We show that pure cubic GaN was successfully deposited by NH3-based CBE for the first time to our knowledge. The initial substrate treatment
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seems to greatly affect the crystal symmetry of following GaN films. For GaN growth by CBE, the use of an argon/nitrogen mixture instead of pure nitrogen for the nitridation step increases yield of cubic phase GaN thin films. These films have better structural quality than those deposited after plain N2-ECR nitridation, which yielded a mixed phase film. By carefully choosing the substrate nitridation condition and plasma composition, one can better control the GaN layer phase composition during its growth on the GaAs substrate. For the ECR MBE growth, cubic phase GaN layer was achieved only after high-temperature substrate annealing.
Acknowledgements This work is supported by fund from a NASA cooperative agreement, No. NCC8-127, to the Space Vacuum Epitaxy Center. I.R. would like to acknowledge support from the Texas Center for Superconductivity (TCSUH) at the University of Houston.
References [1] S.N. Mohammad, A.A. Salvador, H. Morkoc-, Proc. IEEE 83 (1995) 1306. [2] H. Okumura, K. Balakrishnan, G. Feuillet, K. Ohta, H. Hamaguchi, S. Chichibu, Y. Ishida, S. Yoshda, Mater. Res. Soc. Symp. Proc. 449 (1997) 435. [3] S. Ruvimov, Z. Liliental-Weber, J. Washburn, T.J. Drummond, M. Hafish, S.R. Lee, Appl. Phys. Lett. 71 (1997) 1931. [4] K.H. Ploog, O. Brandt, H. Yang, B. Yang, A. Trampert, J. Vac. Sci. Technol. B 16 (1998) 2229.
[5] M.E. Lin, G. Xue, G.L. Zhou, J.E. Greene, H. Morkoc, Appl. Phys. Lett. 63 (1993) 932. [6] G. Pozina, I. Ivanov, B. Monemar, J.V. Thordson, T.G. Andersson, J. Appl. Phys. 84 (1998) 3830. [7] L. Bellaiche, S.-H. Wei, A. Zunger, Appl. Phys. Lett. 70 (1997) 3558. [8] H. Chen, H. Liu, Z. Li, S. Liu, Q. Huang, J. Zhou, Y. Wang, J. Crystal Growth 201/202 (1999) 336. [9] G. Shimaoka, T. Aoki, Y. Nakanishi, Y. Hatanaka, T. Udagawa, Appl. Surf. Sci. 175–176 (2001) 436. [10] M. Losurdo, P. Capezzuto, G. Bruno, E.A. Irene, Phys. Rev. B 58 (1998) 15878. [11] A. Georgakilas, K. Amimer, P. Tzanetakis, Z. Hatzopoulos, M. Cengher, B. Pecz, Zs. Czigany, L. Toth, M.V. Baidakova, A.V. Sakharov, V.Yu. Davydov, J. Crystal Growth 227–228 (2001) 410. . J.V. Thordson, J.R. Gunnarson, Q.X. Zhao, L. [12] O. Zsebok, Ilver, T.G. Andersson, J. Appl. Phys. 89 (2001) 3662. . J.V. Thordson, T.G. Andersson, Jpn. J. Appl. [13] O. Zsebok, Phys. 40 (2001) 472. [14] A. Trampert, O. Brandt, H. Yang, K.H. Ploog, Appl. Phys. Lett. 70 (1997) 583. [15] O. Brandt, H. Yang, A. Trampert, M. Wassermeier, K.H. Ploog, Appl. Phys. Lett. 71 (1997) 473. [16] T.S. Cheng, C. Jemkins, S.E. Hooper, C.T. Foxon, J.W. Orton, D.E. Lacklison, Appl. Phys. Lett. 66 (1995) 1509. [17] Y. Zhao, C.W. Tu, I.-T. Bae, T.-Y. Seong, Appl. Phys. Lett. 74 (1999) 3182. [18] A. Hashimoto, T. Motizuki, H. Wada, A. Masuda, A. Yamamoto, J. Crystal Growth 201/202 (1999) 336. [19] A. Bensaoula, W.T. Taferner, E. Kim, A. Bousetta, J. Crystal Growth 164 (1996) 185. [20] A.R. Krauss, O. Auciello, J.A. Schultz, MRS Bull. 20 (5) (1995) 18. [21] E. Kim, I.E. Berishev, A. Bensaoula, J.A. Shultz, K. Waters, S. Lee, S.S. Perry, Appl. Phys. Lett. 71 (1997) 3072. [22] L.A. DeLouise, J. Vac. Sci. Technol. A 10 (1992) 1637. [23] A. Neubert, K.F. Zombov, J. Chem. Soc. Faraday Trans. 1 70 (1974) 2219. [24] D.F. McMillen, D.M. Golden, Ann. Rev. Phys. Chem. 33 (1982) 493.