Growth of scandium doped GaN by MOVPE

Superlattices and Microstructures 60 (2013) 120–128

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Growth of scandium doped GaN by MOVPE C. Saidi a,⇑, N. Chaaben a, A. Bchetnia a, A. Fouzri b, N. Sakly c, B. El Jani a a Unité de Recherche sur les Hétéro-Epitaxies et Applications, Faculté des Sciences de Monastir, Avenue de l’Environnement, 5019 Monastir, Tunisia b Laboratoire Physico-Chimie des Matériaux, Unité de Service Commun de Recherche ‘‘High Resolution X-ray Diffractometer’’, Département de Physique, Université de Monastir, Faculté des Sciences de Monastir, Avenue de l’Environnement, 5019 Monastir, Tunisia c Laboratoire de Physique et Chimie des Interface, Faculté des Sciences de Monastir, Avenue de l’Environnement, 5019 Monastir, Tunisia

a r t i c l e

i n f o

Article history: Received 21 March 2013 Accepted 8 May 2013 Available online 16 May 2013 Keywords: MOVPE GaN Scandium Reflectivity HRXRD

a b s t r a c t Scandium (Sc) doped GaN layers (GaN:Sc) were grown on SiN treated sapphire substrate by atmospheric pressure metal–organic vapor phase epitaxy (AP-MOVPE) using Tris(cyclopentadienyl)scandium (Cp3Sc) as a scandium precursor. Standard growth conditions of GaN were used with varying Cp3Sc flow rate up to 1000 sccm. The growth was in situ monitored by laser reflectometry. GaN coalescence process is significantly influenced by the presence of Cp3Sc. Sc incorporation in GaN is confirmed by secondary ion mass spectrometry (SIMS) measurements. Electron concentration and mobility variations versus Cp3Sc flow rate were studied by room temperature Hall effect measurements. All samples showed n type conductivity. High resolution X-ray diffraction (HRXRD) results showed a slight broadening of rocking curves around symmetric (00.2) reflexions for Sc doped GaN layers with increasing Cp3Sc flow rate. Room temperature photoluminescence (PL) showed a quenching of ultraviolet near band edge emission (UV) by increasing Sc doping level. Ó 2013 Elsevier Ltd. All rights reserved.

1. Introduction III-nitrides have attracted extensive interest because of their excellent physical properties including large direct gaps and high thermal conductivity [1]. These properties make them useful in the fabrication of optoelectronic devices such as light emitting diodes and laser diodes. GaN is one of the most promising materials in the fabrication of these devices. By allowing GaN with AlN and/or InN it is possible to create materials with direct band gaps spanning the entire visible and ultraviolet spectrum. Low In content In⇑ Corresponding author. Tel.: +216 97 77 15 40. E-mail address: [email protected] (C. Saidi). 0749-6036/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.spmi.2013.05.010

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GaN based light emitting diodes and blue laser diodes with high internal quantum efficiencies (IQEs) are obtained [2]. However, the IQEs considerably drop for higher In content InGaN alloys. In addition, InN presents low melting temperature which limits the use of these devices at high temperatures. These problems have motivated the researchers to look for alternative materials with high thermal stability to replace InN. Recently, it has been suggested that transition metal nitrides could be good replacements of InN. In this regard, the alloys of GaN with ScN are of great interest [3–7]. ScN is stable in the rock-salt structure, with an indirect band gap near 1.3 eV and a direct band gap near 2.1 eV [3,8]. Theoretical and experimental studies suggested that low Sc content GaScN alloys might be stabilized in the wurtzite structure [9–11]. Epitaxial wurzite structure Ga1xScxN films with low Sc content (x 6 0.17) are successfully grown on sapphire substrates [11–15]. The greater part of Ga1xScxN films have been grown by sputter-deposition [12] and molecular beam epitaxy (MBE) [13–15], but they are never grown by MOVPE. The main reason is the lack of suitable organometallic precursor of Sc with a high vapor pressure at room temperature. In this paper we present preliminary results about the growth of low Sc incorporated GaN films on SiN treated sapphire substrate by MOVPE. The Tris(cyclopentadienyl)scandium (Cp3Sc) is used as Sc precursor. In situ Laser reflectometry, SIMS, Hall effect, HRXRD and PL were performed to study the Sc incorporation effects on GaN layers’ properties. 2. Experimental details GaN:Sc layers were grown by AP-MOVPE on SiN treated sapphire substrate at temperature of 1120 °C. Ammonia (NH3) and trimethylgallium (TMG) were used as nitrogen (N) and gallium (Ga) precursors respectively, while Cp3Sc was used as Sc precursor. The Cp3Sc is a solid source at room temperature with a melting point of about 240 °C. This solid precursor is stored inside a bubbler which is placed in a constant temperature bath. During samples growth, the Sc source bubbler was heated to 90 °C with downstream lines maintained at the same temperature to avoid source condensation. The growth process starts by nitridation of sapphire substrate under NH3 + H2 + N2 atmosphere during 10 min at 1080 °C. This step is followed by the so-called ‘‘SiN treatment’’ which consists in a deposition of SiN mask on sapphire substrate at the same temperature during 75 s. The SiN coating is obtained by introducing simultaneously SiH4 (50 sccm) and NH3 (2 slm) in vapor phase. The substrate temperature was then lowered to 600 °C for growing 25 nm thick GaN buffer layer. After that, the temperature was elevated to 1120 °C for growing Sc-doped GaN layers. The flow rates of NH3 and TMG were maintained at 3 slm and 10 sccm (40 lmol/min), respectively. A mixture of N2 (2 slm) and H2 (2 slm) was used as carrier gas. Using these optimized growth conditions [16,17], several samples of GaN:Sc having a thickness of about 1.2 lm were elaborated with different Cp3Sc flow rates varying from 0 to 1000 sccm. A summary of the growth parameters about all samples are listed in Table 1. All growth stages were in situ controlled by He–Ne laser-reflectometry under normal incident light (k = 632.8 nm) [18,19]. The incorporation of Sc in GaN films is investigated by SIMS. The Sc doping effects on structural, electrical and optical properties of GaN were investigated respectively by HRXRD, Hall effect and PL measurements. 3. Results and discussion Fig. 1 illustrates in situ reflectivity signals during the growth of undoped and Sc doped GaN epilayers with Cp3Sc flow rate varying from 300 to 1000 sccm. The presented reflectivities and their averages Table 1 A summary of growth parameters about all samples. Sample

Growth temperature (°C)

Pressure (atm)

NH3 (mmol/min)

TMG (lmol/min)

Cp3Sc (sccm)

1 2 3 4 5

1120 1120 1120 1120 1120

1 1 1 1 1

133 133 133 133 133

40 40 40 40 40

0 300 500 700 1000

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are normalized by the sapphire substrate reflectivity (Rs). For all samples, when starting GaN growth, we record a drastic decrease of the reflectivity signal which is caused by the formation of uncoalesced GaN islands. GaN starts to grow by a three-dimensional (3D) mode. Then, islands coalescence leads to a transition from three-dimensional (3D) to bi-dimensional (2D) growth mode that is announced by oscillations and an increase of average reflectivity. When average reflectivity reaches a constant value, a complete coalescence and a 2D growth mode were achieved leading to a good quality and smooth surface of GaN [20–23]. For the best growth conditions, reflectivity signal reaches the low temperature GaN buffer layer one (Rb). In our growth process, buffer thickness is controlled by the ratio r = Rb/Rs at the instant of sending TMG into the growth chamber to grow GaN buffer layer. In our case, the set point value of r is usually fixed at 2.1 corresponding to a 25 nm thick GaN buffer on sapphire substrate. So, the r value for each normalized reflectivity curve can be taken as an absolute maximum of the normalized average reflectivity. If we consider the growth starting instant of undoped GaN (TMG ON) and GaN:Sc (TMG + Cp3Sc ON) epilayers as zero time (see Fig. 1), then the instant when the average reflectivity reaches 75% of r value can be considered as the recovery time (Tr) [24]. This parameter quantifies the coalescence duration. From Fig. 1, one can note that the Cp3Sc presence modifies the coalescence phase (3D–2D). The time needed to reach the bi-dimensional growth mode (Tr) for the undoped sample is about 300 s. It increases to reach 950 s by increasing the Cp3Sc flow rate at about 1000 sccm. The same phenomena has been appeared in Mg doped GaN [25]. These authors suggested that this effect was due to the competing capture of Mg atoms into chemisorbed III site with Ga atoms. In our case, by introducing TMG and Cp3Sc fluxes a competition between Ga and Sc capture can occur. This mechanism may enhances the vertical growth rate in the initial stages of film growth increasing the recovery time as in indicated in the reflectivity spectra.

Reflectivity (a.u.)

-500 4 3 2 1 0 4 3 2 1 0 4 3 2 1 0 4 3 2 1 0 4 3 2 1 0

0

500

1000

1500

2000

1000

1500

2000

(TMG+Cp3Sc) ON

0 sccm

300 sccm

500 sccm

700 sccm

1000 sccm

-500

0

500

Time (s) Fig. 1. In situ reflectivity signals of GaN for different Cp3Sc flow rates. All reflectivity signals are normalized by the substrate reflectivity (Rs) and horizontally translated to have the same zero time corresponding to TMG ON that is indicated by the vertical dashed line. The dotted lines are the average reflectivity curves. The full circles indicate the recovery time which correspond to 75% of the buffer layer reflectivity.

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The average reflectivity is a description of morphological details of layers because an overall increase in roughness will tend to decrease the average reflectivity intensity of the oscillations [26]. In Fig. 1, the average reflectivity decrease with increasing Cp3Sc flow rate indicates that surface becomes progressively rougher. In contrast to the other average reflectivity signals, the one corresponding to 1000 sccm does not reach the r value indicating that surface degradation becomes relatively drastic for the highest value of Cp3Sc flow rate. Similar behavior was observed in vanadium (V) doped GaN [27]. The quality of GaN is likely to depend on the recovery speed and the crystallographic orientations between GaN and the sapphire substrate at coalescence phase (3D–2D). Song et al. have reported that the growth of GaN layers with the longer recovery time causes a surface degradation [28]. However it has been reported in Refs. [29–31] that the coalescence time variation has no influence on the average roughness of GaN layers. According to the aforementioned results, we suggested that the Sc presence during GaN growth slowdown surface diffusion and lateral growth, resulting in a rough surface, compared to that of undoped GaN layer. Fig. 2 presents the AFM (3  3 lm2) images of GaN layers together with the extracted AFM surface roughness and average height islands plots versus Sc doping. It is clearly shown that the GaN surface morphology is influenced by the presence of Sc atoms during the growth process. For undoped GaN, at a very small scale employed in AFM analysis, we can see a high density of nanoislands that are homogenously distributed on the surface. With the Sc doping increase, the lateral size of these nanoislands increases slightly to reduce their density. This is due to coalescence. Moreover, some of them have a larger lateral size indicating that nanoislands can grow independently. However, due to the vertical growth rate dominance, islands average height increases more than their lateral size. One can also clearly note in Fig. 2d that the surface roughness and the average height of nanoislands increase linearly with Sc doping. This is in good agreement with reflectivity traces showing the coalescence delay. These results suggest that Sc presence during GaN growth will slow down the lateral growth process. Consequently, GaN surface morphology becomes rougher with increasing Cp3Sc flow rate. It seems that Sc presence modifies diffusion kinetics on GaN growing surface. The same behavior was also observed in the growth of transition metal doped GaN such as V [27], Cr [32] and Fe [33]. Deterioration of surface layers’ quality was observed during V and Cr doped GaN growth. It was also found that Fe doping can play a surfactant role and lead to a 2D-layer-by-layer growth at the initial stage but after that it leads to a 3D mode [33]. In order to check Sc incorporation in GaN matrix, SIMS analysis was achieved. Unlike to others impurities (Si, Mg, V, Cr, etc.), SIMS data for Sc element in GaN are very poor in the literature. Fig. 3 shows Sc, Si, C and Ga profiles as a function of penetration depth. This analysis was made on Sc doped GaN sublayer grown with Cp3Sc flow rate of 1000 sccm on 300 nm thick undoped GaN layer. For GaN:Sc sublayer, a flat profile is observed indicating an homogenous Sc incorporation in GaN matrix. When depth exceeds GaN:Sc sublayer thickness, Sc level is also flat but lower than SIMS detection limit. At the interface GaN:Sc/GaN, a some tailing of Sc profile in undoped GaN layer. This broadening is may be associated to surface roughness and/or Sc diffusion in undoped GaN layer. On the other hand, according to the SIMS data of Si and C, one can note that they keep the same amount in GaN:Sc sublayer and in undoped GaN thick layer. This means that the use of Cp3Sc precursor dose not introduces unintentional doping as well as introducing Sc. The Si and C concentrations are in the same order of magnitude (1017 atoms/cm3) as previously reported in MOVPE GaN grown layers [34]. Fig. 4 shows the typical HRXRD patterns obtained for the undoped and the more doped GaN layers using 2h/x scan. In addition to the (00.6) reflection of Sapphire substrate, only two peaks located at 34.52° and 72.89° that are respectively associated to (00.2) and (00.4) reflections of GaN are shown. They keep the same positions in HRXRD patterns indicating that: firstly, GaN:Sc has an hexagonal structure growing toward c-axis and secondly, no secondary or ternary phase such as ScN or GaScN are detected. To get a reliable measure of the structural quality, both symmetric (00.2) and asymmetric (10.2) rocking curves (x scan) were performed for undoped and all doped samples (Fig. 5). It can be seen from the Table 2, compared to that of undoped GaN, the rocking curves of symmetric (00.2) reflection of doped GaN show a broadening that increases with Cp3Sc flow rates. However, the rocking curve broadening of asymmetric (10.2) reflection does not significantly changes with increasing Cp3Sc flow rate. Generally, broadening of rocking curves is due to the crystalline quality degradation such as

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22.66 nm

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(b) 56. 42 nm

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Average height (nm)

r.m.s. roughness

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400

600

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1000

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average height

0

Cp 3 Sc flow rate (sccm)

(d)

Fig. 2. 3  3 lm2 AFM images of undoped and GaN:Sc layers. (a) undoped GaN, (b) and (c) GaN:Sc with Cp3Sc flow rates of 500 and 1000 sccm, respectively. (d) Root mean square (r.m.s.) roughness and average height evolutions versus Cp3Sc flow rates.

presence of defects and/or stress in epitaxial films. It is also known that broadening of (00.2) plan rocking curve results from the increase of screw dislocations with Burgers vector b = h00.1i, while broadening of in plan rocking curve results from the increase of edge dislocations with Burgers vector b = 1/3h11.0i. The broadening of others planes (h0.l) rocking curves results from mixed dislocations with Burgers vector b = 1/3h11.3i [35]. There are usually less than 2% screw dislocations [36,37], but the ratio of mixed to edge dislocations is variable [38]. For the films containing mostly edge dislocations a very low broadening of (00.l) plan rocking curve are expected. However, in all Cp3Sc flow rates range, the FWHM of (00.2) reflection increases by a factors of 1.52 but the one of (10.2) does not significantly changes. This behavior can be explained by the bending of edge dislocations away from the [00.1] direction [39] that can occurs during the coalescence process of GaN:Sc layers. If g.(b  u) – 0 (where g is the reflection (00.l), b is the Burgers vectors and u is the dislocation line direction) then the dislocation will affect x-scans of the reflection g. For example, dislocations with b = 1/3h11.0i inclined away from [00.1] will broaden x-scans of (00.l) reflection, whereas pure edge dislocations running along [00.1] will not. The increase of (00.2) FWHM could be also attributed to the increase of Sc content in GaN approaching the characteristics of GaScN alloys [11].

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5

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Sc Ga C Si

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Concentration (at/cm )

Secondary Ions intensity (counts/s)

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1,0

Depth (µm) Fig. 3. SIMS depth profiles of Sc, Si, C and Ga in 1000 sccm Sc doped GaN grown on first undoped GaN epilayer on Al2O3.

(00.2) GaN

Intensity ( a.u.)

(00.6) Al2O3

(00.4) GaN

Sc doped GaN

32

36

40

44

48

52

56

60

64

68

72

(00.2) GaN

(00.4) GaN

(00.6) Al2O3

Undoped GaN

32

36

40

44

48

52

56

60

64

68

72

2θ (degree) Fig. 4. HRXRD 2h/x diffractograms for undoped and 1000 sccm Sc doped GaN layers.

Room temperature PL measurements were performed using a He–Cd laser emitting at ke = 325 nm. A special care was taken to meaningfully compare sample-to-sample PL intensity. Fig. 6a presents the room temperature PL spectra obtained for undoped and GaN:Sc layers with different Cp3Sc flow rates. For undoped GaN layer, PL spectrum displays two peaks namely the ultraviolet (UV) near band edge emission at 3.4 eV and the broad yellow luminescence (YL) at around 2.2 eV. For Cp3Sc flow rate higher than 300 sccm, UV emission disappears whereas YL emission persists. The UV emission extinction may be associated to the presence of Sc element rather than to structural defects acting as non-radiative recombination centers. The YL emission observed for undoped and GaN:Sc is mostly attributed to

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(10.2)

0

-3000

-2000

-1000

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1000 sccm

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Normalized Intensity (u.a)

Normalized Intensity (a.u)

(00.2)

3000

1000

-3000

-2000

-1000

0

1000

0 sccm

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3000

ω (arcsec)

ω (arcsec)

Fig. 5. (a) Symmetric (00.2) and (b) asymmetric (10.2) HRXRD rocking curves for undoped and GaN:Sc layers.

Table 2 Room temperature mobility (l) and carrier concentration (n) in GaN:Sc for different Cp3Sc flow rates and their corresponding FWHM of (00.2) and (10.2) rocking curves values. Cp3Sc flow rate (sccm) Carrier concentration n  1019 (cm3) Mobility l (cm2/Vs) FWHM (00.2) (arcsec) FWHM (10.2) (arcsec)

0 0.02 450 511.8 931.8

300 3.04 33 617.8 906.1

0.25

700 sccm

300 sccm

YL UV 0 sccm

YL PL Intensity (a.u)

PL intensituy (a.u)

1000 4.56 23 779.3 797.0

(b)

1000 sccm

2.0

700 6.12 19 670.7 803.6

0.30

(a)

1.8

500 3.06 13 548.0 880.0

0.20

0.15

0.10

0.05

2.2

2.4

2.6

2.8

Energy (eV)

3.0

3.2

3.4

0.00

0

200

400

600

800

1000

Cp3 Sc flow rate (sccm)

Fig. 6. (a) Room temperature PL spectra obtained for GaN:Sc layers with different Cp3Sc flow rates. The PL spectra are vertically shifted for clarity. (b) The YL band integrated intensity evolution with Sc doping.

Ga vacancies [33]. The Fig. 6b shows the decrease of YL band integrated PL intensity versus Sc doping. This behavior is similar to those observed for transition metal doped GaN layers. In fact, it has been reported that V and Fe incorporation in trivalent site of Ga reduces Ga vacancies which can suppress

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the YL band emission [27,33]. Thus, the reduction of Yellow band integrated PL intensity could be related to the Sc incorporation in Ga sites. In Table 2, room temperature Hall effect measurements of GaN:Sc are presented with their corresponding values of the FWHM of (00.2) and (10.2) rocking curves. We clearly show that in the used Cp3Sc flow rates range, GaN:Sc has an n-type conductivity. We also note the increase of carrier concentration (n) from 2  1017 to 6  1019 cm3 and the decrease of mobility (l) from 450 to 33 cm2/ Vs for undoped GaN and the less doped GaN, respectively. Then, they slightly vary with the same tendency having the same order of magnitude when the Cp3Sc flow rate is increased. The decrease in carrier mobility follows the crystalline quality degradation of GaN:Sc. Electron mobility is generally influenced by defects density [40]. Canstantin et al. have reported that Sc incorporation in Ga site produces a lattice distortion that induces defects [11]. Furthermore, according to the aforementioned PL results, we considered that the YB quenching is due to the decrease of Ga vacancies acting as an acceptor center in GaN [41]. Thus, this will contribute to the carrier concentration increase. On the other hand, the observed change in the coalescence time by the presence of Sc may be also considered as another reason for the carrier concentration increase. This phenomenon was previously reported in reference [42]. While the Sc addition perturbs the growth mechanism of GaN, the incorporation amount seems to be insufficient to attain GaScN alloys. Contrary to the MBE [15], the MOVPE growth process based on the vapor transport of the precursor molecules is very complex. Besides, the reactant parasitic reactions before reaching the substrate could lead to precursor losses by formation of adduct pathways which greatly influence the Sc incorporation. Thus, a development of suitable scandium precursor with a high vapor pressure will be needed to increase the Cp3Sc amount which must be delivered into the reactor. Accordingly, future work will focus on the adjustment of the MOVPE growth conditions such as growth temperature and pressure to expect better results. 4. Conclusion GaN:Sc layers were grown by MOVPE at 1120 °C using Tris(cyclopentadienyl)scandium (Cp3Sc) as scandium precursor. The Cp3Sc flow rate increase leads to a GaN coalescence delay. Sc incorporation in GaN matrix was confirmed by SIMS analysis. HRXRD analysis shows that GaN:Sc layers are perfectly oriented with c-axis on the sapphire substrate and presents degradations of their crystalline qualities when Cp3Sc flow rate is increased. All GaN:Sc layers have n type electrical conductivity with a decrease of carrier mobility following crystalline properties. Room temperature PL study shows that optical properties of GaN:Sc layers are affected by the presence of Sc element. The quenching of the principal UV emission is tentatively associated to the non radiative recombination centers and structure defects that are induced by the Sc incorporation in GaN. Acknowledgement This work is supported by an ARUBE program of DGRST. References [1] S. Strite, H. Morkoç, J. Vac. Sci. Technol. B 10 (1992) 1237. [2] H. Morkoc, Handbook of nitride semiconductors and devices, GaN-based Optical and Electronic Devices, vol. 3, WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim, 2009. [3] D. Gall, I. Petrov, N. Hellgren, L. Hulman, J.-E. Sundgren, J.E. Greene, J. Appl. Phys. 84 (1998) 6034. [4] D. Gall, I. Petrov, L.D. Madsen, J.-E. Sundgren, J.E. Greene, J. Vac. Sci. Technol. A 16 (1998) 2411. [5] M.A. Moram, Z.H. Barber, C.J. Humphreys, Thin Solid Films 516 (2008) 8569. [6] M.A. Moram, S.V. Novikov, A.J. Kent, C. Nörenberg, C.T. Foxon, C.J. Humphreys, J. Cryst. Growth 310 (2008) 2746. [7] J.L. Hall, M.A. Moram, A. Sanchez, S.V. Novikov, A.J. Kent, C.T. Foxon, C.J. Humphreys, R.P. Campion, J. Cryst. Growth 311 (2009) 2054. [8] W.R.L. Lambrecht, Phys. Rev. B 62 (2000) 13538. [9] V. Ranjan, L. Bellaiche, E.J. Walter, Phys. Rev. Lett. 90 (2003) 257602. [10] V. Ranjan, S. Bin-Omran, D. Sichuga, R.S. Nichols, L. Bellaiche, A. Alsaad, Phys. Rev. B 72 (2005) 085315. [11] C. Constantin, H. Al-Brithen, M.B. Haider, D. Ingram, A.R. Smith, Phys. Rev. B 70 (2004) 193309. [12] M.E. Little, M.E. Kordesch, Appl. Phys. Lett. 78 (2001) 2891.

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