GaN epitaxial growth on sapphire (0 0 0 1): the role of the substrate nitridation

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Journal of Crystal Growth 178 (1997) 220-228

GaN epitaxial growth on sapphire (0 0 0 1): the role of the substrate nitridation N. Grandjean*, J. Massies, Y. Martinez, P. Venn6gu6s, M. Leroux, M. Lai.igt Centre de Recherche sur l 'H~t~ro-Epitaxie et ses Applications. Centre National de la Recherche Scientifique, Rue B. Grdgory-Sophia Antipolis, F-06560 Valbonne, France

Received 13 August 1996; accepted 12 November 1996

Abstract

The nitridation of sapphire substrates was monitored in situ by reflection high-energy electron diffraction. The evolution of the lattice-mismatch evidences the formation of an A1N relaxed layer when exposing the sapphire surface heated at 850'C to an ammonia flow. GaN thin films were grown by gas-source molecular beam epitaxy on variously nitridated sapphire substrates. High-resolution transmission electron microscopy study reveals the existence of two different epitaxial relationships between GaN and sapphire (0 0 0 1). The well-known orientation with the c-axis of the GaN crystal perpendicular to the A1203 surface is observed when starting the growth on a nitridated substrate. On the other hand, growing GaN directly on bare AI203 surfaces results in a different crystallographic orientation where the c-axis is tilted by about 19 with respect to the sapphire basal plane. Photoluminescence measurements show that both the intensity of the yellow-band emission (~ 2.2 eV) and the residual donor-acceptor pair recombinations are affected by the nitridation state of the starting surface. The dependence of the optical properties of GaN thin films versus the NH3 exposure time is then used to optimize the nitridation step. PACS.. 61.14.Hg; 61.16.Bg; 68.35.Ct; 68.55.Jk; 78.55.Cr; 81.15.Hi K e y w o r d s : III-V nitrides; Sapphire nitridation; Epitaxial relationship; Optical properties

1. Introduction

G a N and related nitride c o m p o u n d s are attracting a great deal of research due to their potential for blue-violet light-emitting devices [1]. This has been reinforced since the marketing of high-efficiency G a N - b a s e d light-emitting diodes (LEDs) [2],

* Corresponding author. E-mail: [email protected].

and the I I I - V nitride research is concerned now with the aim of current-injected laser diodes (LDs) in the blue-violet spectral region. Very recently, N a k a m u r a et al. have successfully achieved a L D operating at r o o m temperature at 417 nm under pulsed current injection [3]. This demonstrates the potentiality of nitrides for short wavelength lightemitting devices. However, t h o u g h N a k a m u r a ' s group has shown its high degree of mastery over the III V nitride material system, the growth of

0022-0248/97/$17.00 Copyright ~ 1997 Elsevier Science B.V. All rights reserved P l l S 0 0 2 2 - 0 2 4 8 ( 9 6 ) 0 1 191-8

N. Grand/ean et al, / Journal of" C~ystal Growth 178 (1997) 220 228

high-quality G a N and related compounds is far to be fully understood due to its complexity. One origin of this is the lack of well-suited substrates such as G a N single crystals having sufficiently large dimensions. Sapphire has then become the standard substrate in spite of its large lattice-mismatch with GaN. Indeed, it allows to get high-quality GaN material providing the growth of a buffer layer at low temperature [4]. As a matter of fact, the first LEDs and LDs mentioned previously have been grown on sapphire substrates [2, 3]. However, it has been recently evidenced that the use of a lowtemperature buffer layer must be joined to a nitridation of the sapphire substrate prior to the G a N growth. In fact, it has been reported that the nitridation step is essential for growing high-quality III V nitrides on A1203 substrates, whatever the growth technique [5 9]. Astonishingly, the exact function of the sapphire nitridation on the G a N growth is still unclear. In the present paper, the nitridation of c-plane sapphire substrates is investigated in situ by reflection high-energy electron diffraction (RHEED) in a gas-source molecular beam epitaxy (GSMBE) system. G a N overlayers have then been checked by high-resolution transmission electron microscopy (HRTEM), showing the influence of the nitridated sapphire surface on the epitaxial relationship between GaN and sapphire (0 0 0 1). A correlated photoluminescence (PL) study also shows that the optical properties of the G a N over|ayers are strongly affected too by the nitridation state of the starting sapphire surface. The tendencies given by PL are then used to optimize the nitridation procedure.

2. Experimental procedure The growth was performed in a G S M B E system RIBER 32P) equipped with a standard R H E E D facility. Group-III elements are supplied by solid sources and nitrogen is provided by the catalytic decomposition of ammonia at the G a N growing surface. NH3 is introduced through a 20 sccm mass flow controller into the growth chamber (the maximum pressure is 1 x 10 6 Torr). The growth temperature of G a N varies from 550°C to 820°C and

221

the deposition rates (determined in situ using the reftectivity oscillations of a 670nm laser diode) from 0.1 to 1.2 gm/h. The sapphire substrates have the c-plane orientation. H R T E M was carried out in cross section with a J E O L 2010 field emission gun microscope with a point resolution of 0.19 nm and an information limit better than 0.15 nm. TEM objects were prepared by a combination of mechanical polishing and ion milling. PL measurements were performed at 10 K with an H e - C d laser excitation, a 64 cm spectrometer and a cooled GaAs photomultiplier. A 100 W quartz lamp provided white light for the reflectivity spectra.

3. RHEED analysis of the sapphire nitridation The nitridation of c-plane sapphire substrates was carried out by exposing the surface to an ammonia flow of 20 sccm at a substrate temperature of 850c'C. The modification of the sapphire surface during this procedure has been followed by R H E E D through the variation of the in-plane lattice constant (Fig. l a). The lattice mismatch is measured from the R H E E D streak distance with the electron beam along the [1 T 0 0]A~:O, azimuth. Whereas the R H E E D pattern remains streaky-like,

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N. Grandjean et al. ,i Journal o]" Crystal Growth 178 (1997) 220 228

the in-plane surface parameter strongly varies when the sapphire substrate is exposed to NH3. The lattice-mismatch saturates at ,-~ 13% after 10 rain. This corresponds to the theoretical mismatch between A1N and A1203 (12.8%) [10]. This evidences that the sapphire surface is converted into an A1N layer when exposed to an ammonia flow. This is confirmed by growing G a N on substrate nitridated for a long time (60 rain) and for which the whole surface can be supposed to be an A1N relaxed layer. In this case, the relaxation of the G a N film with respect to the starting surface is in agreement with the expected lattice mismatch between GaN and A1N (Fig. lb). Finally, when G a N is grown directly on a bare sapphire surface (Fig. 2), the lattice mismatch is much larger and approaches that of G a N on sapphire (the slight discrepancy with the actual mismatch of GaN/A1203 (15.9%) will be discussed in the following part) [11]. All these R H E E D observations attest the formation of a crystalline A1N relaxed layer during the nitridation of the sapphire surface [8, 9].

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4. Effect of the sapphire nitridation on the GaN overlayer orientation Thin G a N films ( ~ 200 ,~) were deposited on two different starting sapphire surfaces in order to investigate the role of the nitridation step. Sample A corresponds to the standard procedure, i.e. GaN is grown on a nitridated substrate as checked by RHEED. It should be recalled that most of the GaN growth studies have been carried out following this experimental setup because it is wellknown to strongly improve the G a N layer quality [5-7]. In sample B, G a N was directly deposited on sapphire substrate, i.e. without nitridation. The structural properties of the two samples have been characterized by TEM. Fig. 3 exhibits a cross section image of sample A. The GaN crystallites have mainly a skyscraper-like shape. A typical G a N island is displayed in Fig. 4. Note the good crystal quality since only one stacking fault can be seen. The epitaxial relationship between G a N and the nitridated c-plane sapphire deduced from the H R T E M observation is in agreement with previous reports [12 17]: the c-axis of G a N is perpendicular

Fig. 3. Bright field transmission electron microscopy image of sample A which corresponds to 200 ,~, of G a N grown on a nitridated sapphire substrate.

to the A1203 surface and [1 i 0 0]~aNIl[l 1 2 0]AI~O3. It corresponds to an in-plane rotation of 30 ° of the G a N lattice with respect to the substrate one. Let us turn now to sample B for which no nitridation was performed. Compared to sample A, the G a N islands are much larger [18] and their stacking fault density is greater (Figs. 5 and 6 a). Surprisingly, the orientation of the c-axis is no more perpendicular to the sapphire basal plane but is in fact almost parallel. It should be noted that recent works by Raman spectroscopy or TEM report

N. Grandiean et al. / Journal o['Crystal Growth 178 (1997) 220 228

223

Fig. 5. Plane view transmission electron microscopy image of sample B. The dark areas are the GaN islands while the bright ones correspond to the bare sapphire surface.

Fig. 4. High-resolution lransmission electron micrograph of a typical GaN crystallite found in sample A. The sapphire and the GaN lattices are viewed in the [1 i 00J and [ 1 1 2 0 1 projections, respectively. The c-axis of GaN is perpendicular to the sapphire basal plane (Inset: selection area diffraction pattern).

similar disorientation of the GaN layers near the substrate [19 21]. In fact, diffraction (Fig. 6a) and high-resolution (Fig. 6b) TEM studies reveal that the c-axis of the G a N crystallite is tilted by about 19 ° with respect to the sapphire basal plane, the 30" in-plane rotation being conserved. The corresponding epitaxial relationship can be defined by [1 1 2 0 ] G a N I I [ l i 0 0 ] A I _ ~ O ` and [-li03]GaNI] [1 120]Al:O,. The GaN plane which is parallel to the sapphire basal plane is (3 3 0 2). It makes a

theoretical angle of 19.5 ° with the c-axis, in good agreement with the value of 19 ° measured on TEM micrographs. Let us come back to the R H E E D measurement of the lattice mismatch when GaN grows with this epitaxial relationship. In this case, two 180': symmetric orientations of the GaN c-axis are perpendicular to the [1 i 0 0]A~2o~direction corresponding to the electron beam azimuth. This results in the appearance of two 0002 Bragg positions on the R H E E D pattern which are close to the 1 0 i 0 and ]- 0 1 0 positions observed when GaN is deposited on a nitridated sapphire substrate. Actually, Fig. 2 displays an apparent lattice-mismatch deduced from the distance between these two 0002 Bragg positions. In this configuration, the expected mismatch should be 14.6% which is in rather good agreement with the measured value of 13-14% (Fig. 2). The main part of the GaN crystallites which are observed in sample B have the tilted orientation; a few other ones have the standard orientation with the c-axis perpendicular to the sapphire basal plane as for the nitridated A1203 surface [22]. We suggest that the presence of the standard orientation for some crystallites in sample B is due to the threedimensional growth of GaN on sapphire substrates (as shown in Fig. 5). Indeed, the part of the sapphire surface which is not covered during the G a N island growth is simultaneously nitridated by the ammonia flux. In such nitridated areas, the

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N. Grancl/ean el a/. / Journal o/" Crystal Growth 178 (1997) 220-228

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s u b s e q u e n t G a N growth can then occur following the s t a n d a r d o r i e n t a t i o n as in the case of sample A. An a r g u m e n t in favor of this hypothesis is that only the smaller crystallites shown in Fig. 5 have their c-axis p e r p e n d i c u l a r to the sapphire basal plane. The two previous epitaxial relationships are illustrated in Fig. 7. Sapphire a n d G a N lattices are viewed in the [1 ] 0 0] a n d [1 1 2 0] projections, respectively. The s t a n d a r d o r i e n t a t i o n with the caxis p e r p e n d i c u l a r to the A1203 (0 0 0 1) plane is

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displayed in Fig. 7a a n d the tilted one for which the c-axis a n d the [1 1 2 0]A~2o~direction form an angle of 19 ° is reported in Fig. 7b. Fig. 7a a n d Fig. 7b show that the lattice-matching c o n d i t i o n s between G a N a n d A1203 are similar for the two o r i e n t a t i o n s (the lattice matching positions are indicated by solid circles). It corresponds to a ratio between lattice distances close to 6/7. In fact, the exact inter-atomic

N. Grandjean et al. / Journal o[ Crystal Growth 178 (1997) 220-228

distances for the lattice matching conditions are 16.49A for the sapphire basal plane, 16.57 and 16.51 A for the (0 0 0 1) G a N plane (perpendicular orientation) and for the (3 3 0 2) G a N plane (tilted orientation), respectively. The closer value which exists for the (3 3 0 2) plane could explain the preference for a tilted orientation of G a N crystal on bare c-plane sapphire substrates. It is easier to understand why the perpendicular orientation of the c-axis occurs on a nitridated surface. When the A1203 (0001) surface is exposed to an ammonia flux, an A1N thin layer is formed, as deduced from the lattice-mismatch measurement. It is suggested that the formation of that A1N layer is due to the conversion of AI octahedral sites in tetrahedral ones by partial oxygen-nitrogen substitution. This should result in an A1N layer rotated by 30 ° from the sapphire lattice because of the arrangement of the Al atoms in the A120 3 crystal (they form an hexagonal sublattice rotated by 30 ° about the caxis with respect to the A1203 unit cell). In fact, we propose that the nitridation of the sapphire surface leads to the formation of a pseudo-A1N (0001) substrate. Consequently, G a N will grow with the same orientation that the A1N layer, i.e. with the c-axis perpendicular to the sapphire basal plane. To our knowledge, similar results have been reported only very recently by Christiansen et al. in GSMBE-grown G a N layers [21]. They observed by TEM some grainlets having the tilted orientation near the interface. This is surprising in view of the numerous studies which have been already devoted to the growth of G a N on c-plane sapphire substrate. A first reason for that can be the relatively little number of work addressed on GaN grown directly on sapphire due to the resulting poor material quality [-5-7]. Moreover, most of the studies do not concern the early stage of the growth, but rather thick layers for which a reorientation of the c-axis from tilted to perpendicular can occur as the thickness increases. Recent Raman experiments reinforce this argument. Indeed, Siegle et al. deduced from Raman selection rules that the G a N c-axis is parallel to the sapphire surface near the interface but turns by about 90" for a greater GaN thickness [20]. Actually, the fact that the tilted orientation is generally not observed comes more likely from the competition between the G a N deposition and the

225

A1203 surface nitridation occurring simultaneously at the beginning of the growth, as previously discussed. Indeed, at the beginning of the G a N growth, the part of the surface which is not immediately covered by GaN will be rapidly nitridated (in a few minutes, see Fig. la). As a consequence, G a N will grow on this nitridated areas with the perpendicular orientation. In the case of metal-organic chemical vapor deposition (MOCVD), which is the most commonly used G a N growth technique, the temperature is generally greater than 1000°C [-1], so that the nitridation is more efficient than in the case of GSMBE. Therefore, even if no intentional nitridation step is performed, AIN is formed on the sapphire surface simultaneously to the GaN nucleation, inducing the standard orientation for subsequent GaN deposition. In our case, the GaN growth temperature is lower (780°C), not allowing a nitridation to proceed faster than the GaN island growth itself. The tilted orientation can then be preponderant (sample B). It should be also emphasized that in the present experiments, the G a N growth is stopped before island coalescence, thus avoiding a possible reorientation of the crystallites which can occur for thicker layers, as suggested by recent observations [20, 21].

5. Optical properties of thin GaN layers versus sapphire nitridation The optical properties of samples A and B have been investigated by low-temperature (10K) photoluminescence (Fig. 8). The band-edge luminescence for the two samples is dominated by a band at 3.42 eV. Following Chung and Gershenzon [23], we tentatively assign it to free hole-oxygen donor transitions. A possible origin is the diffusion of oxygen coming from the A1203 substrate [9, 24]. The shoulder on the high-energy side (3.47 eV) of this PL peak is due to band-edge excitonic recombinations corresponding most presumably to residual donor bound excitons (it is not possible for such thin 3D film to get resolved reflectivity spectra that could allow a more precise identification). A feature of interest here is that the donor acceptor (D°A °) pair recombinations at 3.27 eV are clearly more pronounced when nitridation is not performed. The

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N. Grandjean et al. /Journal of Co'stal Growth 178 (1997) 220 228

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Wavelength (A) Fig. 8. Low-temperature (10 K) photoluminescence spectra of samples A and B. Both the donor acceptor pair recombination and the yellow-band emission intensities are strongly pronounced when GaN is grown directly on sapphire substrate (sample B). On the other hand, nitridation of sapphire prior to the growth (sample A) improves the optical properties of the GaN overlayer (note the disappearance of the yellow band).

lower optoelectronic quality of sample B is confirmed by the presence of a deep-level emission at 2.2 eV, the so-called yellow band, which is not detected in sample A. The apparent correlation between the yellow band and the G a N crystal orientation is in agreement with recent results showing that the yellow emission is located near the interface where in some places the c-axis is no more perpendicular to the sapphire basal plane [20]. Our observations suggest that the D°A ° and the yellow-band luminescence intensities are dependent on the G a N growing surface orientation. Of course, structural quality is strongly dependent on crystallite orientation, but it is also to be mentioned that in the case of GaAs, the relationship between the crystallographic orientation of the growing surface and the impurity incorporation is now well established [25].

6. Optimization of the sapphire nitridation procedure The sensitivity of the optical properties of thin G a N layers as a function of the nitridation state of the starting surface has been used to optimize the

nitridation procedure [9]. This is necessary because the nitridation must be long enough to convert entirely the surface into an A1N layer, whereas too long a nitridation time is known to result in a degradation of the A1203 surface morphology and optical properties [5 7]. Optical properties of G a N thin layers (--~600 A) were investigated as a function of the nitridation time. Two particular features have been considered: the band-edge emission characteristics and the yellow-band intensity. The low-temperature PL spectra of different samples versus the nitridation time are displayed in Fig. 9. When there is no nitridation step, a bandedge emission broadening is observed together with a strong yellow-band intensity. The sample nitridated for 1 min exhibits a sharper band-edge luminescence and a decrease of the yellow-band intensity. The improvement of the optical properties is more accented in the case of the 10 min nitridated sample. This nitridation time corresponds to the maximum of the lattice-mismatch variation as measured by R H E E D (Fig. la) indicating that the surface is entirely converted into A1N. Actually, for a longer nitridation time (30 min), the G a N optical properties are degraded

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N. Grandjean et al. / Journal of Crystal Growth 178 (1997) 220 228

again. The PL behavior as a function of the nitridation time of the sapphire substrate thus allows to determine an optimized nitridation process which consists in our experimental conditions to expose the sapphire substrate to 20 sccm of NH3 at 850°C for 10 rain.

7. Growth of high-quality GaN layers The optimized nitridation step was used to grow thick GaN layers. A two-step growth procedure [4] was carried out by growing at 550°C a 270 A thick G a N buffer layer followed by 0.8 ~tm of GaN grown at a substrate temperature of 820°C and a growth rate of 0.56 gm/h. The 10 K PL and reflectivity spectra of this sample are reported in Fig. 10. The three excitons (labeled A, B, C) associated to the wurtzite hexagonal structure are well resolved in the reftectivity spectrum together with interference fringes at larger wavelength (inset in Fig. 10). The PL is dominated by a peak at 3.471 eV. Its energy position, relative to the A exciton recorded in reflectivity, allows to assign it to residual donor bound excitons (I2 line) [26]. The full-width at

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In summary, we have demonstrated by in situ R H E E D measurements that the nitridation of the sapphire substrate involves the conversion of the A1203 surface into an A1N relaxed layer. H R T E M studies reveal that the role of the nitridation is to favor the well-known epitaxial relationship where the c-axis of G a N is perpendicular to the sapphire basal plane. On the other hand, growing GaN directly on a bare sapphire substrate involves a different crystallographic orientation in which the caxis of GaN is tilted by 19 ° with respect to the (0 0 0 1) sapphire plane. The PL properties of G a N thin layers depend on the crystal orientation, and therefore, on the nitridation state of the starting surface. This dependence is finally used to optimize the nitridation step.

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N. Grandjean et al. /Journal Of Ct3'stal Growth 178 (] 997) 220 228

Acknowledgements We gratefully acknowledge fruitful discussions with B. Beaumont. Special thanks to J.C. Supergillon for the use of the R H E E D data analysis facility (SG-RHEED). We also wish to thank E. Tourni6 for critical reading of the manuscript, and J.P. Faurie for strong support and interest. The MBE machine used in this work is a donation of the CNET.

[12] [13] [14] [15] [16] [17]

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