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Structural, optical and electrical study of undoped GaN layers obtained by metalorganic chemical vapor deposition on sapphire substrates
Thin Solid Films 519 (2011) 2255–2261
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Thin Solid Films j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / t s f
Structural, optical and electrical study of undoped GaN layers obtained by metalorganic chemical vapor deposition on sapphire substrates Victor-Tapio Rangel-Kuoppa a,⁎, Cesia Guarneros Aguilar b, Victor Sánchez-Reséndiz b,⁎ a
Institute of Semiconductor and Solid State Physics, Johannes Kepler Universität, A-4040 Linz, Austria Sección de Electrónica del Estado Sólido, Departamento de Ingeniería Eléctrica, Centro de Investigación y de Estudios Avanzados del Instituto Politécnico Nacional, A.P. 14740, C.P. 07360, México, Distrito Federal, Mexico b
a r t i c l e
i n f o
Article history: Received 24 March 2010 Received in revised form 15 October 2010 Accepted 28 October 2010 Available online 3 November 2010 Keywords: Gallium nitride Metal-organic chemical vapor deposition Scanning electron microscopy X-ray diffraction Photoluminescence Hall effect measurements Atomic force microscopy
a b s t r a c t We investigate optical, structural and electrical properties of undoped GaN grown on sapphire. The layers were prepared in a horizontal reactor by low pressure metal organic chemical vapor deposition at temperatures of 900 °C and 950 °C on a low temperature grown (520 °C) GaN buffer layer on (0001) sapphire substrate. The growth pressure was kept at 10,132 Pa. The photoluminescence study of such layers revealed a band-to-band emission around 366 nm and a yellow band around 550 nm. The yellow band intensity decreases with increasing deposition temperature. X-ray diffraction, atomic force microscopy and scanning electron microscopy studies show the formation of hexagonal GaN layers with a thickness of around 1 μm. The electrical study was performed using temperature dependent Hall measurements between 35 and 373 K. Two activation energies are obtained from the temperature dependent conductivity, one smaller than 1 meV and the other one around 20 meV. For the samples grown at 900 °C the mobilities are constant around 10 and 20 cm2 V−1 s− 1, while for the sample grown at 950 °C the mobility shows a thermally activated behavior with an activation energy of 2.15 meV. © 2010 Elsevier B.V. All rights reserved.
1. Introduction The III–V–N material system has entered the focus of research due to their interesting properties, which makes it suitable for light emitting devices in the UV and visible range [1], metal–semiconductor field effect transistors [2], sun blind solar detectors [3], and more. In particular, GaN has a large direct band gap and high saturation velocity, which makes it a candidate for high temperature and high power electron devices, such as the high electron mobility transistors [4]. Despite these promising applications, very few studies have been done on its electrical transport properties, and in particular, only some articles can be found on temperature dependent Hall (T–Hall) measurements on undoped and n-type material. Among the first electrical studies of undoped material, Maruska et al. [5] reported room temperature (RT) Hall measurements, while the Hall results of Ilegems et al. [6] were performed between 10 K and 1125 K. In both cases the material was grown by vapor deposition on (0001) sapphire. Later on, Hall studies were done mainly at RT [7–17]. Other researchers extended their experiments also to 77 K [18,19]. Some research was performed on the temperature dependent Hall (T-Hall) measurements on material grown on the (0001) plane (the c plane). They were Si-doped grown by either molecular beam epitaxy (MBE) ⁎ Corresponding authors. E-mail addresses: [email protected] (V.-T. Rangel-Kuoppa), [email protected] (V. Sánchez-Reséndiz). 0040-6090/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2010.10.053
or metal organic chemical vapor deposition (MOCVD) [20,21] or undoped or highly resistive samples grown by MBE [22,23]. Also, some results on Hydride Vapor Phase Epitaxy and on Low Pressure Chemical Vapor Deposition grown material can be found for the temperature range of 13 K to 600 K [24–26]. T-Hall results were obtained also on GaN grown on the (1120) plane (the r plane). These samples were Si-doped by MOCVD [27] and investigated between 70 K and 400 K, or undoped by MBE [28] from 10 K to 300 K. Recently, electrical measurements have been done on c plane GaN material grown on r plane, by MOCVD. In one case, the material was undoped and only the RT Hall results were given [29], while on the other case the material was Si-doped and the Hall measurements were done between 77 K and 300 K [30]. Also, some theoretical studies have appeared: Weimann et al. [16] calculated the Hall coefficient for 300 K and Shrestha et al. [31] published theoretical results for the temperature dependent Hall effect. The intention of this article is to complement the temperature dependent electrical transport study of undoped GaN grown by MOCVD on the c plane of sapphire substrates, under different trimethylgallium (TMGa) flow rates and growth temperatures. T-Hall measurements were done between 35 K and 373 K. We also report the structural and optical properties of our GaN samples. An important difference between our study and the one done by Ilegems et al. [6] is that they use the vapor phase reaction between GaCl and NH3 and He as the carrier gas. Their growth temperature is around 1050 °C. In our study we use TMGa, NH3, H2 and lower temperatures.
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This article is divided into the following sections: in Section 2 experimental details about the sample growth are given. The optical and structural studies are reported in Section 3, while the Ohmic contact formation and its application to T-Hall are described and discussed in Section 4. Discussion of the results and comparison with the literature is given in Section 5, and finally conclusions are given in Section 6. 2. Experimental details GaN layers were deposited on (0001) oriented (Al2O3) sapphire substrates using low pressure MOCVD horizontal quartz reactor [32]. The heater around the quartz reactor was composed of seven infrared lamps; the electrical power per lamp was 1600 W. The Al2O3 substrates were cleaned ultrasonically in organic solvents, chemically etched in a hot H3PO4:H2SO4 (1:3) solution for 10 min, and rinsed in de-ionized water before loading them into the reactor. First, the sapphire substrate temperature (TS) was increased to 900 °C in hydrogen flow. Afterwards, the hydrogen flow was changed for ammonia. These first steps were done for 15 min. A low temperature GaN buffer layer (520 °C) was grown for 5 min with a thickness of 20 nm using trimethylgallium (TMGa) and ammonia (NH3) as Ga and N precursors, respectively. The carrier gas was H2. The carrier gas was purified using a Pd-diffusion cell. After depositing the GaN buffer layer, TS was raised to 900 °C or to 950 °C, to grow GaN epitaxial layer for 90 min. The TMGa flow rate was kept constant, with 5.39 μmol/min for the first and second experiments (samples S900-A and S950-A in Table 1), while for the third experiment the TMGa flow rate was 8.39 μmol/min (sample S900-B). The hydrogen and ammonia flow rates were also kept constant at 3 slm and 0.5 slm, respectively. The growth pressure was kept at 10,132 Pa. Table 1 summarizes the experimental growth conditions. 3. Optical and structural analysis Photoluminescence (PL) measurements were obtained using a 1404 double grating spectrometer at RT. The light source is a He–Cd laser with a wavelength of 325 nm. Fig. 1 shows the PL spectra of GaN with different deposition temperatures for the epilayers measured at RT. The dominant emission peak at 366 nm corresponds to the band edge emission. A small emission can be seen at 425 nm (blue emission), which is characteristic of MOCVD grown material [15]. The yellow band emission was observed at 550 nm. It has been suggested that the yellow band luminescence could be associated with the dislocations or perhaps native point defects that nucleate at the dislocation sites [33]. PL spectra of samples grown at TS = 900 °C exhibit a broad yellow emission typical of poor-quality GaN layers. However, the weakness of the yellow emission for GaN films grown at TS = 950 °C is a sign attribute for improved quality of GaN grains and their clustering to islands with micrometer size. Variable Angle Spectroscopic Ellipsometry (Gaertner Scientific & Co.) was used to determine the refractive index and the thickness of GaN in an incidence angle range from 45° to 80°. The light source was He–Ne laser and two laser diodes with wavelengths of 632.8 nm and 824.3 nm, respectively. The refractive index was determined as 2.3, consistent with another report [34]. The thickness of the GaN layer was determined to be approximately 1 μm. Table 1 Growth parameters of the samples. Sample name
TMGa (μmol/min)
NH3 (slm)
H2 (slm)
TS (°C)
P (Pa)
Growth time (min)
Thickness (nm)
S900-A S950-A S900-B
5.39 5.39 8.39
0.5 0.5 0.5
3 3 3
900 950 900
10,132 10,132 10,132
90 90 90
1165 1330 1192
Fig. 1. Room-temperature photoluminescence spectra for samples S900-A, S950-A and S900-B.
X-ray diffraction (XRD) measurements were performed using an X-Ray diffractometer Panalytical X-Pert Pro MRD with Cu target operated at 40 kV and 20 mA. Fig. 2 shows the XRD pattern representative of all samples. The peak at 34.54° corresponds to GaN with wurtzite structure. The shoulder at smaller angles is due to dislocations. Atomic force microscopy (AFM) measurements were done using a Quesant 250 AFM, in contact mode. The studied area was 10 μm × 10 μm to observe the surface morphology of GaN layers grown at 900 °C and 950 °C, respectively. AFM images shown in Fig. 3 illustrate the effect of the growth temperature on the surface morphology of the films. The surface of the sample S900-A (grown at 900 °C) was composed of the planar regions and the three
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4. Electrical analysis: Ohmic contact formation and its application to T-Hall measurements
Fig. 2. X-ray diffraction pattern representative of all the GaN samples. The peak at 34.54° corresponds to (0001) plane of GaN with wurtzite structure.
dimensional (3D) islands, while the surface of the sample S950-A (grown at 950 °C) was mostly occupied by the 3D islands. The root mean square roughness for sample S900-A and sample S950-A were 95 nm and 190 nm, respectively. Scanning electron microscopy (SEM) was done using a Tescan Vega TS-5136SB SEM equipment. The applied voltage was 10 kV. The SEM analysis (Fig. 4) shows that the layer grown at 900 °C (S900-A) exhibits good morphology with small islands on the surface. As TS is increased to 950 °C (S950-A), the island size increases as well, forming a rough surface with holes at the surface, confirming the analysis carried out by AFM. Sample S900-A has a smaller grain size and higher island density. Obviously, it takes less time for the island to coalesce at lower substrate temperature and the surface will never reach a high degree of roughness, because the decreased mobility of surface species promotes a uniform distribution of islands that can effectively cover the substrate. Sample S950-A has larger grain sizes and lower density (Fig. 3). Here individual islands are developed before they start to coalesce (which imposes an increase of surface roughness). These studies suggest a dislocation density around 3 × 1012 cm− 2 and 1012 cm− 2 for S900-A and S950-A, respectively. Larger grain sizes and rougher nucleation morphology lead to an increased volume of defect-free columnar domain, a lower density of dislocations associated with the domain boundaries, and improved structural and electrical properties [35].
In order to control the size and position of the Ohmic contacts on the samples, the films were covered with 200 nm SiO2 by plasma enhanced chemical vapor deposition, and using standard photolithographic techniques (as those reported in Ref. [36]), 50 nm Ti and 100 nm Au were deposited in a van der Pauw geometry. No thermal treatment was done. The contacts were Ohmic at all temperatures. In this study, the surface is passivated by the SiO2; hence, the conduction happens through the bulk material. Au wires were soldered to these Ohmic contacts. The samples were glued to the sample holder of the Hall cryostat with Fixogum©, to ensure good thermal conductivity [37]. Afterwards, the chamber was pumped to 10− 6 mbar, and measurements were done in darkness between 35 K and 373 K. The magnetic field was set to 600 mT. In order to avoid any sample heating, the dissipated power due to the applied voltages and currents was controlled to be below 0.1 mW. For convenience, we express the relation of the Hall coefficient R and the resistivity ρ with the charge carrier density n and the conductivity σ: n=
1 eR
ð1Þ
1 : ρ
ð2Þ
and σ=
Finally, the Hall mobility μ is obtained by μ=
R : ρ
ð3Þ
These expressions are valid when charge transport is due to a single type of charge. This is the case in this study, as it will be seen. Our T-Hall setup yields ρ, μ and n. The ρ is calculated using the van der Pauw method, which consists of calculating the sheet resistance RS and multiplying it by the thickness of the layer [38,39] to obtain ρ. Results are given in Fig. 5. We briefly comment on the interpretation of the measured ρ. As was indicated in Section 2, our samples are grown in three distinct steps: the hydrogen and ammonia flow, the 20 nm buffer GaN layer deposition, and finally the deposition of the ~1 μm GaN bulk layer. Thus, a proper model to interpret the Hall results should include all
Fig. 3. Comparison of AFM images of GaN surface for samples prepared using 5.39 μmol/min and substrate temperatures of 900 °C (S900-A) and 950 °C (S950-A), in a 10 μm × 10 μm area.
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Fig. 4. Comparison of SEM images of GaN surface for samples prepared using 5.39 μmol/min and substrate temperatures of 900 °C (S900-A) and 950 °C (S950-A).
three layers [40]. The first step might form some AlN layer between the substrate and the GaN buffer layer. Even if it were present, for AlN layers ρ values of ~ 1013 Ω × cm or larger have been reported [41]. Regarding the buffer layer, in the optical and structural analysis section it was commented that it contains 3 × 1012 cm− 2 dislocations. In their paper, Weimann et al. analyze the effect of dislocations in transport measurements [16]. In their Fig. 6, the transverse mobility decreases with increasing dislocation density. The highest dislocation density they report is 3 × 1011 cm− 2. Thus, their analysis for their highest dislocation density yields an upper limit to the transverse mobility of our samples. In their Fig. 6, the highest mobility for a dislocation density of 3 × 1011 cm− 2 is 50 cm2 V− 1 s− 1. According to their Fig. 5, for this mobility and this dislocation density, the free carrier concentration is 1017 cm− 3. The resistivity in this case would be 1.25 Ω × cm. Dividing this value by the thickness of our buffer layer (20 nm) yields an upper limit of 6.25 × 106 Ω for the sheet resistance of our buffer layer, at all temperatures. On the other hand, the smallest resistivity we measure in our samples is around 4.63 × 10− 3 Ω × cm (sample S900-A in Fig. 5 at 373 K). Dividing this value by the sample thickness (1.1 μm), yields a sheet resistance for sample S900-A of 42.09 Ω at 373 K. Clearly, the measured sheet resistance of all the samples is six orders of magnitude smaller than the contribution of the buffer layer and at least twelve orders of magnitude smaller than the contribution of any possible AlN layer. The contribution of each sheet resistance to the total sheet resistance is described by a harmonic sum: the inverse of the total sheet resistance is the sum of the inverse of each sheet resistance [40]. Thus, the contributions of the
buffer layer and any possible AlN layer to the Hall measurements are negligible. Arrhenius plot of the conductivity for samples S900-A, S950-A and S900-B are shown in Fig. 6. Assuming two thermally activated processes, we obtain the following fits: S900−A
σ Fit
meV −1 −0:151meV −1 −68:58 kT kT = 149:23ðΩ×cmÞ + 19:53ðΩ×c mÞ ×e ×e
ð4Þ meV S950−A −1 −0:765kTmeV −1 −21:52 kT σ Fit = 38:405ðΩ×cmÞ + 40:417ðΩ×cmÞ ×e ×e
ð5Þ S900−B
σ Fit
meV −1 −0:245 −1 −21:2 meV kT = 70:3ðΩ×c mÞ + 30:35ðΩ×cmÞ ×e × e kT
ð6Þ where k is the Boltzmann constant. All samples are n-type nature at all temperatures. The mobilities for samples S900-A and S900-B did not show a clear trend with temperature: they fluctuate around 20 and 10 cm2 V− 1 s− 1, respectively, at all temperatures. Their charge carrier densities oscillate around 5.7× 1019 and 4.5× 1019 cm− 3, respectively. These values are half of those reported by Ma et al. [30]. This was not the case for sample S950-A, as can be seen in Fig. 7. Our sample S950-A shows a similar trend as the one reported by Ma et al. [30]. A semilog plot of the mobility vs inverse of the temperature shows a linear behavior, as can be seen in Fig. 8. The linear fit yields a thermally activated process of the form 2 −1 −1 −2:15 meV 46:15 · cm V s × e kT
ð7Þ
A semilog plot of the charge carrier density vs inverse of the temperature shows also a linear behavior, as can be seen in Fig. 9. This shows a thermally activated process of the form 18 −3 −5:41 meV 9:72 × 10 cm × e kT ð8Þ 5. Discussion and comparison with literature
Fig. 5. Resistivity vs temperature for samples S900-A, S950-A and S900-B.
In his work, Weimann et al. showed how the decreasing mobility as a function of decreasing free carrier concentration in GaN, can be explained by the scattering of electrons by threading dislocations [16]. They found that there is one dislocation per hexagonal island. Our SEM analysis showed that for samples S900-A and S900-B, the number of islands (and thus of dislocations) is in the order 3 × 1012 cm− 2.
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Fig. 7. Mobility and charge carrier density vs temperature for sample S950-A.
et al. between 400 and 700 °C [28] and Molnar et al. should have used a TS of 600 °C, according to their references [22]). In their paper, they obtained mobilities 50% larger than ours for similar charge carrier densities. This could be due to the better crystal quality obtained by MBE than by MOCVD growth techniques. Our measured conductivities show two thermally activated processes (Eqs. 4–6). The first term has a very small activation energy (b1 meV). This was observed also by Molnar et al. [22], and also in the charge carrier concentration by Ng et al. [20]. As Ng et al. pointed out [20], possible explanations are impurity band transport or a non-negligible contribution of the buffer layer to the transport at low temperatures [20]. This second argument is discarded in our case, due to the high resistivity of the GaN buffer layer as discussed. Interestingly, the mobility decreases below 150 K. The charge carrier density is constant between 150 K and 75 K, and increases below 75 K. Ilegems et al. [6] reported that the mobility of their samples decreased below 150 K, while the charge carrier density increases. They conclude that the conduction behavior of their samples should be explained by impurity band transport below 150 K, and by donors over 150 K. Something similar might be happening with our sample S950-A: impurity band transport contributes to the transport between 150 K and 75 K, and it becomes dominant below 75 K. Kusakabe et al. [27] and Ma et al. [30] have found a linear dependency between the mobility and temperature below 200 K. Nakamura et al. [21] reported seven points between 100 K and 40 K, while Molnar et al. showed fourteen points below 100 K down to 10 K [22]. Ng et al. reported only six values for temperatures below 100 K [20]. Unfortunately, in Fig. 6. Fitting of the conductivity vs temperature for samples S900-A, S950-A and S900-B.
Comparing with Fig. 6 of Ref. [16], this explains why for our samples S900-A and S900-B the mobility is constant and between 10 and 20 cm2 V− 1 s− 1: the number of dislocations is so large, that it dominates the scattering mechanism at all temperatures. This is similar to the RT results reported by Sasaki et al. [15] and Li et al. [29] for their MOCVD grown material. Both papers report that they first treated the substrates with H2 at high temperatures (Sasaki et al. at 1200 °C [15] and Li et al. at 1150 °C [29]). Li et al. grew a 25 nm GaN buffer layer, in contrast to Sasaki et al. [15,29]. Sasaki et al. used similar growth temperatures as we did, but larger ranges of TMGa flux (4 to 30 μmol min) and larger NH3 gas flows than ours (1 to 2.5 slm) [15]. Li et al. used a TS of 1200 °C but no further details are found about the growth [29]. Our result suggests that a lower TS of 900 °C for the H2 flow is enough to obtain the same effects. Our results are also comparable with the MBE results reported by Yoshida et al. [18], Eddy et al. [28] and Molnar et al. [22]. Interestingly, they used smaller TS than ours (Yoshida et al. used 700 °C [18], Eddy
Fig. 8. Semilog plot of the mobility vs the inverse of the temperature for sample S950-A and its linear fitting.
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Fig. 9. Semilog plot of the charge carrier density vs the inverse of the temperature for sample S950-A and its linear fitting.
none of these studies are there further details about how the Hall measurements were performed at these low temperatures: neither on whether the sample was cooled down using a gas, or by placing it on a cold finger in vacuum. For the latter case, one of us has recently reported that the thermal glue used to place the sample on a sample holder can have a deep impact in the T-Hall measurements, especially below 100 K [37]. This could be a source of error in the measurements described and could explain the unusual linear dependence between the mobility and the temperature reported by Kusakabe et al. [27] and Ma et al. [30]. Anyhow, 1=2 as Ng et al. [20] have pointed out, the mobility has a n dependency, and thus, a thermally activated process of the form ΔE μ ≈ exp − 2kT
ð9Þ
where ΔE is the thermally activated energy for the charge carrier density. This is what we observed with Eqs. (7) and (8): the activation energy for the mobility is approximately half of the value of the activation energy for the charge carrier density. This suggests that in our sample S950-A, scattering is controlled by charge trapped in the dislocations, for temperatures higher than 150 K. Also, our sample S950-A showed a smaller density of islands (and thus, of dislocations). SEM studies suggest they are in the order of 1012 cm− 2. Our T-Hall results fit well to Fig. 6 of Ref. [16]. According to Fig. 3 of Ref. [16], this means that around 90% of the traps due to the dislocations are filled. Our study thus suggests that, for our undoped MOCVD GaN samples, the scattering model proposed by Weimann et al. explains well the temperature dependence of the mobility between 150 K and 373 K [16]. We comment on the similarity of our results with some of the MBE grown samples reported by Yoshida et al. [18] and Molnar et al. [22]. Yoshida et al. grew an AlN layer before the deposition of their GaN material [18]. This was not done by Molnar et al. [22]. However, the results of Molnar et al. for their samples 114 and 115 are similar to our samples S900-B and S950-A, respectively [22]. The same happens with Yoshida et al. sample grown with an AlN buffer layer and our sample S950-A [18]. Also, Yoshida et al. observed a decrease of the blue and yellow luminescence when using an AlN layer before the deposition of GaN [18]. This is similar to what we observed between our samples S900-A and S950-A. In other words, a similar effect to the one obtained by depositing an AlN layer prior to the deposition of undoped GaN by MBE growth can be obtained by MOCVD growth, doing the hydrogen and ammonia preliminary flow treatments, and increasing the TS from 900 °C to 950 °C. Finally, we briefly mention why our sample deposited with a TS of 950 °C (S950-A) is expected to have different properties than those
samples deposited with a TS of 900 °C (S900-A and S900-B). In their work, Khan et al. grew all their samples with a TS of 900 °C [42], and in particular, their sample 2 showed similar resistivity, charge carrier density and mobility to our samples S900-A and S900-B. On the other hand, Sasaki et al. reported samples grown at TS values between 950 °C and 1100 °C and 10,132 Pa [43]. They reported TMGa, NH3 and H2 flows with values between 20 and 40 μmol/min, 2.5 slm and 4.4 slm, respectively, which are larger than our growth conditions [43]. Another important difference is that they deposited first an AlN buffer layer. Nevertheless, they mentioned that their growth conditions formed truncated-island-like hillock. This is similar to what we observed in our sample grown with a TS of 950 °C (Figs. 3 and 4). In other words, with our research, we have shown that using smaller TMGa, NH3 and H2 flows, and no AlN buffer layer, we can obtain the same results as reported by Sasaki et al. [43]. It can be expected that the use of larger TS values like 1000 °C or 1050 °C, would yield larger mobilities and smaller charge carrier densities. This is what is usually observed as the growth temperature is increased [43]. TS values larger than 1100 °C should be avoided, as they cause thermal pitting on the GaN surface [43]. 6. Conclusions In this paper the electrical, optical and structural properties of undoped GaN films grown on sapphire substrates by LP-MOCVD, under different growth conditions, were reported. PL spectra of GaN films measured at RT showed that there is strong emission at 366 nm, which is related to the GaN band gap. In addition, weak blue emission, characteristic of MOCVD grown material, and emission in yellow range (550 nm) are seen. The surface morphology and the optical luminescence of GaN layers depend sensitively on the growth temperature. The layer grown with a TS of 900 °C shows a broad yellow luminescence originated from the defect levels, while at a larger growth temperature (with a TS of 950 °C) the quality of epitaxial GaN improves and the yellow emission is reduced. A TS of 950 °C reduces the yellow luminescence and increases the roughness of the surface, by the coalescence of islands. This supports the hypothesis that the yellow luminescence is related to dislocations or native point defects as these are reduced with the formation of islands. Also, it is only for the samples grown at 950 °C that some trend can be seen for the mobility and the charge carrier density. The temperature dependence of the mobility and charge carrier densities are well explained by the scattering mechanism due to charge trapped at threading dislocations. For material grown at 900 °C, the mobilities do not depend on temperature and their values are around 20 and 10 cm2 V− 1 s− 1, decreasing with increasing TMGa flow. The charge carrier densities are also temperature independent, and decrease from 5.7 × 1019 to 4.5 × 1019 cm− 3 as the TMGa flow increases. This is explained by a high density of dislocations (3 × 1012 cm− 2). For material grown at 950 °C, the islands coalesce and the dislocations are reduced. Still the dominant scattering mechanism is scattering by charge trapped at the threading dislocations, but with a thermally activated dependence for the mobility and charge carrier density. The activation energies are 2.15 and 5.41 meV, respectively. Our results confirm the scattering model by Weimann et al. [16]. Acknowledgements Victor-Tapio Rangel-Kuoppa gratefully acknowledges the Fonds zur Förderung der Wissenschaftlichen Forschung, Vienna, Austria and the National Council for Science and Technology (CONACyT) of Mexico, postdoctoral fellowship 78965. Victor Sánchez Reséndiz expresses his gratitude to the National Council for Science and Technology (CONACyT) from México, grant 56486. The authors would like to thank Elisabeth Pachinger, Alma Halilovic, Ursula Kainz, Daniel Ramirez, Miguel Avendaño, Rogelio
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Fragoso, Eckehard Nusko, Otmar Fuchs, and Stephan Bräuer for technical assistance. The manuscript revision by Prof. Alberta Bonanni, Prof. Máximo López-López, Prof. Jonathan Farley and Dr. Andrea Navarro-Quezada is gratefully acknowledged. Special gratitude is expressed to Prof. Wolfgang Jantsch for his scientific support and advice. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13]
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Thin Solid Films j o u r n a l h o m e p a g e : w w w. e l s ev i e r. c o m / l o c a t e / t s f
Structural, optical and electrical study of undoped GaN layers obtained by metalorganic chemical vapor deposition on sapphire substrates Victor-Tapio Rangel-Kuoppa a,⁎, Cesia Guarneros Aguilar b, Victor Sánchez-Reséndiz b,⁎ a
Institute of Semiconductor and Solid State Physics, Johannes Kepler Universität, A-4040 Linz, Austria Sección de Electrónica del Estado Sólido, Departamento de Ingeniería Eléctrica, Centro de Investigación y de Estudios Avanzados del Instituto Politécnico Nacional, A.P. 14740, C.P. 07360, México, Distrito Federal, Mexico b
a r t i c l e
i n f o
Article history: Received 24 March 2010 Received in revised form 15 October 2010 Accepted 28 October 2010 Available online 3 November 2010 Keywords: Gallium nitride Metal-organic chemical vapor deposition Scanning electron microscopy X-ray diffraction Photoluminescence Hall effect measurements Atomic force microscopy
a b s t r a c t We investigate optical, structural and electrical properties of undoped GaN grown on sapphire. The layers were prepared in a horizontal reactor by low pressure metal organic chemical vapor deposition at temperatures of 900 °C and 950 °C on a low temperature grown (520 °C) GaN buffer layer on (0001) sapphire substrate. The growth pressure was kept at 10,132 Pa. The photoluminescence study of such layers revealed a band-to-band emission around 366 nm and a yellow band around 550 nm. The yellow band intensity decreases with increasing deposition temperature. X-ray diffraction, atomic force microscopy and scanning electron microscopy studies show the formation of hexagonal GaN layers with a thickness of around 1 μm. The electrical study was performed using temperature dependent Hall measurements between 35 and 373 K. Two activation energies are obtained from the temperature dependent conductivity, one smaller than 1 meV and the other one around 20 meV. For the samples grown at 900 °C the mobilities are constant around 10 and 20 cm2 V−1 s− 1, while for the sample grown at 950 °C the mobility shows a thermally activated behavior with an activation energy of 2.15 meV. © 2010 Elsevier B.V. All rights reserved.
1. Introduction The III–V–N material system has entered the focus of research due to their interesting properties, which makes it suitable for light emitting devices in the UV and visible range [1], metal–semiconductor field effect transistors [2], sun blind solar detectors [3], and more. In particular, GaN has a large direct band gap and high saturation velocity, which makes it a candidate for high temperature and high power electron devices, such as the high electron mobility transistors [4]. Despite these promising applications, very few studies have been done on its electrical transport properties, and in particular, only some articles can be found on temperature dependent Hall (T–Hall) measurements on undoped and n-type material. Among the first electrical studies of undoped material, Maruska et al. [5] reported room temperature (RT) Hall measurements, while the Hall results of Ilegems et al. [6] were performed between 10 K and 1125 K. In both cases the material was grown by vapor deposition on (0001) sapphire. Later on, Hall studies were done mainly at RT [7–17]. Other researchers extended their experiments also to 77 K [18,19]. Some research was performed on the temperature dependent Hall (T-Hall) measurements on material grown on the (0001) plane (the c plane). They were Si-doped grown by either molecular beam epitaxy (MBE) ⁎ Corresponding authors. E-mail addresses: [email protected] (V.-T. Rangel-Kuoppa), [email protected] (V. Sánchez-Reséndiz). 0040-6090/$ – see front matter © 2010 Elsevier B.V. All rights reserved. doi:10.1016/j.tsf.2010.10.053
or metal organic chemical vapor deposition (MOCVD) [20,21] or undoped or highly resistive samples grown by MBE [22,23]. Also, some results on Hydride Vapor Phase Epitaxy and on Low Pressure Chemical Vapor Deposition grown material can be found for the temperature range of 13 K to 600 K [24–26]. T-Hall results were obtained also on GaN grown on the (1120) plane (the r plane). These samples were Si-doped by MOCVD [27] and investigated between 70 K and 400 K, or undoped by MBE [28] from 10 K to 300 K. Recently, electrical measurements have been done on c plane GaN material grown on r plane, by MOCVD. In one case, the material was undoped and only the RT Hall results were given [29], while on the other case the material was Si-doped and the Hall measurements were done between 77 K and 300 K [30]. Also, some theoretical studies have appeared: Weimann et al. [16] calculated the Hall coefficient for 300 K and Shrestha et al. [31] published theoretical results for the temperature dependent Hall effect. The intention of this article is to complement the temperature dependent electrical transport study of undoped GaN grown by MOCVD on the c plane of sapphire substrates, under different trimethylgallium (TMGa) flow rates and growth temperatures. T-Hall measurements were done between 35 K and 373 K. We also report the structural and optical properties of our GaN samples. An important difference between our study and the one done by Ilegems et al. [6] is that they use the vapor phase reaction between GaCl and NH3 and He as the carrier gas. Their growth temperature is around 1050 °C. In our study we use TMGa, NH3, H2 and lower temperatures.
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This article is divided into the following sections: in Section 2 experimental details about the sample growth are given. The optical and structural studies are reported in Section 3, while the Ohmic contact formation and its application to T-Hall are described and discussed in Section 4. Discussion of the results and comparison with the literature is given in Section 5, and finally conclusions are given in Section 6. 2. Experimental details GaN layers were deposited on (0001) oriented (Al2O3) sapphire substrates using low pressure MOCVD horizontal quartz reactor [32]. The heater around the quartz reactor was composed of seven infrared lamps; the electrical power per lamp was 1600 W. The Al2O3 substrates were cleaned ultrasonically in organic solvents, chemically etched in a hot H3PO4:H2SO4 (1:3) solution for 10 min, and rinsed in de-ionized water before loading them into the reactor. First, the sapphire substrate temperature (TS) was increased to 900 °C in hydrogen flow. Afterwards, the hydrogen flow was changed for ammonia. These first steps were done for 15 min. A low temperature GaN buffer layer (520 °C) was grown for 5 min with a thickness of 20 nm using trimethylgallium (TMGa) and ammonia (NH3) as Ga and N precursors, respectively. The carrier gas was H2. The carrier gas was purified using a Pd-diffusion cell. After depositing the GaN buffer layer, TS was raised to 900 °C or to 950 °C, to grow GaN epitaxial layer for 90 min. The TMGa flow rate was kept constant, with 5.39 μmol/min for the first and second experiments (samples S900-A and S950-A in Table 1), while for the third experiment the TMGa flow rate was 8.39 μmol/min (sample S900-B). The hydrogen and ammonia flow rates were also kept constant at 3 slm and 0.5 slm, respectively. The growth pressure was kept at 10,132 Pa. Table 1 summarizes the experimental growth conditions. 3. Optical and structural analysis Photoluminescence (PL) measurements were obtained using a 1404 double grating spectrometer at RT. The light source is a He–Cd laser with a wavelength of 325 nm. Fig. 1 shows the PL spectra of GaN with different deposition temperatures for the epilayers measured at RT. The dominant emission peak at 366 nm corresponds to the band edge emission. A small emission can be seen at 425 nm (blue emission), which is characteristic of MOCVD grown material [15]. The yellow band emission was observed at 550 nm. It has been suggested that the yellow band luminescence could be associated with the dislocations or perhaps native point defects that nucleate at the dislocation sites [33]. PL spectra of samples grown at TS = 900 °C exhibit a broad yellow emission typical of poor-quality GaN layers. However, the weakness of the yellow emission for GaN films grown at TS = 950 °C is a sign attribute for improved quality of GaN grains and their clustering to islands with micrometer size. Variable Angle Spectroscopic Ellipsometry (Gaertner Scientific & Co.) was used to determine the refractive index and the thickness of GaN in an incidence angle range from 45° to 80°. The light source was He–Ne laser and two laser diodes with wavelengths of 632.8 nm and 824.3 nm, respectively. The refractive index was determined as 2.3, consistent with another report [34]. The thickness of the GaN layer was determined to be approximately 1 μm. Table 1 Growth parameters of the samples. Sample name
TMGa (μmol/min)
NH3 (slm)
H2 (slm)
TS (°C)
P (Pa)
Growth time (min)
Thickness (nm)
S900-A S950-A S900-B
5.39 5.39 8.39
0.5 0.5 0.5
3 3 3
900 950 900
10,132 10,132 10,132
90 90 90
1165 1330 1192
Fig. 1. Room-temperature photoluminescence spectra for samples S900-A, S950-A and S900-B.
X-ray diffraction (XRD) measurements were performed using an X-Ray diffractometer Panalytical X-Pert Pro MRD with Cu target operated at 40 kV and 20 mA. Fig. 2 shows the XRD pattern representative of all samples. The peak at 34.54° corresponds to GaN with wurtzite structure. The shoulder at smaller angles is due to dislocations. Atomic force microscopy (AFM) measurements were done using a Quesant 250 AFM, in contact mode. The studied area was 10 μm × 10 μm to observe the surface morphology of GaN layers grown at 900 °C and 950 °C, respectively. AFM images shown in Fig. 3 illustrate the effect of the growth temperature on the surface morphology of the films. The surface of the sample S900-A (grown at 900 °C) was composed of the planar regions and the three
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4. Electrical analysis: Ohmic contact formation and its application to T-Hall measurements
Fig. 2. X-ray diffraction pattern representative of all the GaN samples. The peak at 34.54° corresponds to (0001) plane of GaN with wurtzite structure.
dimensional (3D) islands, while the surface of the sample S950-A (grown at 950 °C) was mostly occupied by the 3D islands. The root mean square roughness for sample S900-A and sample S950-A were 95 nm and 190 nm, respectively. Scanning electron microscopy (SEM) was done using a Tescan Vega TS-5136SB SEM equipment. The applied voltage was 10 kV. The SEM analysis (Fig. 4) shows that the layer grown at 900 °C (S900-A) exhibits good morphology with small islands on the surface. As TS is increased to 950 °C (S950-A), the island size increases as well, forming a rough surface with holes at the surface, confirming the analysis carried out by AFM. Sample S900-A has a smaller grain size and higher island density. Obviously, it takes less time for the island to coalesce at lower substrate temperature and the surface will never reach a high degree of roughness, because the decreased mobility of surface species promotes a uniform distribution of islands that can effectively cover the substrate. Sample S950-A has larger grain sizes and lower density (Fig. 3). Here individual islands are developed before they start to coalesce (which imposes an increase of surface roughness). These studies suggest a dislocation density around 3 × 1012 cm− 2 and 1012 cm− 2 for S900-A and S950-A, respectively. Larger grain sizes and rougher nucleation morphology lead to an increased volume of defect-free columnar domain, a lower density of dislocations associated with the domain boundaries, and improved structural and electrical properties [35].
In order to control the size and position of the Ohmic contacts on the samples, the films were covered with 200 nm SiO2 by plasma enhanced chemical vapor deposition, and using standard photolithographic techniques (as those reported in Ref. [36]), 50 nm Ti and 100 nm Au were deposited in a van der Pauw geometry. No thermal treatment was done. The contacts were Ohmic at all temperatures. In this study, the surface is passivated by the SiO2; hence, the conduction happens through the bulk material. Au wires were soldered to these Ohmic contacts. The samples were glued to the sample holder of the Hall cryostat with Fixogum©, to ensure good thermal conductivity [37]. Afterwards, the chamber was pumped to 10− 6 mbar, and measurements were done in darkness between 35 K and 373 K. The magnetic field was set to 600 mT. In order to avoid any sample heating, the dissipated power due to the applied voltages and currents was controlled to be below 0.1 mW. For convenience, we express the relation of the Hall coefficient R and the resistivity ρ with the charge carrier density n and the conductivity σ: n=
1 eR
ð1Þ
1 : ρ
ð2Þ
and σ=
Finally, the Hall mobility μ is obtained by μ=
R : ρ
ð3Þ
These expressions are valid when charge transport is due to a single type of charge. This is the case in this study, as it will be seen. Our T-Hall setup yields ρ, μ and n. The ρ is calculated using the van der Pauw method, which consists of calculating the sheet resistance RS and multiplying it by the thickness of the layer [38,39] to obtain ρ. Results are given in Fig. 5. We briefly comment on the interpretation of the measured ρ. As was indicated in Section 2, our samples are grown in three distinct steps: the hydrogen and ammonia flow, the 20 nm buffer GaN layer deposition, and finally the deposition of the ~1 μm GaN bulk layer. Thus, a proper model to interpret the Hall results should include all
Fig. 3. Comparison of AFM images of GaN surface for samples prepared using 5.39 μmol/min and substrate temperatures of 900 °C (S900-A) and 950 °C (S950-A), in a 10 μm × 10 μm area.
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Fig. 4. Comparison of SEM images of GaN surface for samples prepared using 5.39 μmol/min and substrate temperatures of 900 °C (S900-A) and 950 °C (S950-A).
three layers [40]. The first step might form some AlN layer between the substrate and the GaN buffer layer. Even if it were present, for AlN layers ρ values of ~ 1013 Ω × cm or larger have been reported [41]. Regarding the buffer layer, in the optical and structural analysis section it was commented that it contains 3 × 1012 cm− 2 dislocations. In their paper, Weimann et al. analyze the effect of dislocations in transport measurements [16]. In their Fig. 6, the transverse mobility decreases with increasing dislocation density. The highest dislocation density they report is 3 × 1011 cm− 2. Thus, their analysis for their highest dislocation density yields an upper limit to the transverse mobility of our samples. In their Fig. 6, the highest mobility for a dislocation density of 3 × 1011 cm− 2 is 50 cm2 V− 1 s− 1. According to their Fig. 5, for this mobility and this dislocation density, the free carrier concentration is 1017 cm− 3. The resistivity in this case would be 1.25 Ω × cm. Dividing this value by the thickness of our buffer layer (20 nm) yields an upper limit of 6.25 × 106 Ω for the sheet resistance of our buffer layer, at all temperatures. On the other hand, the smallest resistivity we measure in our samples is around 4.63 × 10− 3 Ω × cm (sample S900-A in Fig. 5 at 373 K). Dividing this value by the sample thickness (1.1 μm), yields a sheet resistance for sample S900-A of 42.09 Ω at 373 K. Clearly, the measured sheet resistance of all the samples is six orders of magnitude smaller than the contribution of the buffer layer and at least twelve orders of magnitude smaller than the contribution of any possible AlN layer. The contribution of each sheet resistance to the total sheet resistance is described by a harmonic sum: the inverse of the total sheet resistance is the sum of the inverse of each sheet resistance [40]. Thus, the contributions of the
buffer layer and any possible AlN layer to the Hall measurements are negligible. Arrhenius plot of the conductivity for samples S900-A, S950-A and S900-B are shown in Fig. 6. Assuming two thermally activated processes, we obtain the following fits: S900−A
σ Fit
meV −1 −0:151meV −1 −68:58 kT kT = 149:23ðΩ×cmÞ + 19:53ðΩ×c mÞ ×e ×e
ð4Þ meV S950−A −1 −0:765kTmeV −1 −21:52 kT σ Fit = 38:405ðΩ×cmÞ + 40:417ðΩ×cmÞ ×e ×e
ð5Þ S900−B
σ Fit
meV −1 −0:245 −1 −21:2 meV kT = 70:3ðΩ×c mÞ + 30:35ðΩ×cmÞ ×e × e kT
ð6Þ where k is the Boltzmann constant. All samples are n-type nature at all temperatures. The mobilities for samples S900-A and S900-B did not show a clear trend with temperature: they fluctuate around 20 and 10 cm2 V− 1 s− 1, respectively, at all temperatures. Their charge carrier densities oscillate around 5.7× 1019 and 4.5× 1019 cm− 3, respectively. These values are half of those reported by Ma et al. [30]. This was not the case for sample S950-A, as can be seen in Fig. 7. Our sample S950-A shows a similar trend as the one reported by Ma et al. [30]. A semilog plot of the mobility vs inverse of the temperature shows a linear behavior, as can be seen in Fig. 8. The linear fit yields a thermally activated process of the form 2 −1 −1 −2:15 meV 46:15 · cm V s × e kT
ð7Þ
A semilog plot of the charge carrier density vs inverse of the temperature shows also a linear behavior, as can be seen in Fig. 9. This shows a thermally activated process of the form 18 −3 −5:41 meV 9:72 × 10 cm × e kT ð8Þ 5. Discussion and comparison with literature
Fig. 5. Resistivity vs temperature for samples S900-A, S950-A and S900-B.
In his work, Weimann et al. showed how the decreasing mobility as a function of decreasing free carrier concentration in GaN, can be explained by the scattering of electrons by threading dislocations [16]. They found that there is one dislocation per hexagonal island. Our SEM analysis showed that for samples S900-A and S900-B, the number of islands (and thus of dislocations) is in the order 3 × 1012 cm− 2.
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Fig. 7. Mobility and charge carrier density vs temperature for sample S950-A.
et al. between 400 and 700 °C [28] and Molnar et al. should have used a TS of 600 °C, according to their references [22]). In their paper, they obtained mobilities 50% larger than ours for similar charge carrier densities. This could be due to the better crystal quality obtained by MBE than by MOCVD growth techniques. Our measured conductivities show two thermally activated processes (Eqs. 4–6). The first term has a very small activation energy (b1 meV). This was observed also by Molnar et al. [22], and also in the charge carrier concentration by Ng et al. [20]. As Ng et al. pointed out [20], possible explanations are impurity band transport or a non-negligible contribution of the buffer layer to the transport at low temperatures [20]. This second argument is discarded in our case, due to the high resistivity of the GaN buffer layer as discussed. Interestingly, the mobility decreases below 150 K. The charge carrier density is constant between 150 K and 75 K, and increases below 75 K. Ilegems et al. [6] reported that the mobility of their samples decreased below 150 K, while the charge carrier density increases. They conclude that the conduction behavior of their samples should be explained by impurity band transport below 150 K, and by donors over 150 K. Something similar might be happening with our sample S950-A: impurity band transport contributes to the transport between 150 K and 75 K, and it becomes dominant below 75 K. Kusakabe et al. [27] and Ma et al. [30] have found a linear dependency between the mobility and temperature below 200 K. Nakamura et al. [21] reported seven points between 100 K and 40 K, while Molnar et al. showed fourteen points below 100 K down to 10 K [22]. Ng et al. reported only six values for temperatures below 100 K [20]. Unfortunately, in Fig. 6. Fitting of the conductivity vs temperature for samples S900-A, S950-A and S900-B.
Comparing with Fig. 6 of Ref. [16], this explains why for our samples S900-A and S900-B the mobility is constant and between 10 and 20 cm2 V− 1 s− 1: the number of dislocations is so large, that it dominates the scattering mechanism at all temperatures. This is similar to the RT results reported by Sasaki et al. [15] and Li et al. [29] for their MOCVD grown material. Both papers report that they first treated the substrates with H2 at high temperatures (Sasaki et al. at 1200 °C [15] and Li et al. at 1150 °C [29]). Li et al. grew a 25 nm GaN buffer layer, in contrast to Sasaki et al. [15,29]. Sasaki et al. used similar growth temperatures as we did, but larger ranges of TMGa flux (4 to 30 μmol min) and larger NH3 gas flows than ours (1 to 2.5 slm) [15]. Li et al. used a TS of 1200 °C but no further details are found about the growth [29]. Our result suggests that a lower TS of 900 °C for the H2 flow is enough to obtain the same effects. Our results are also comparable with the MBE results reported by Yoshida et al. [18], Eddy et al. [28] and Molnar et al. [22]. Interestingly, they used smaller TS than ours (Yoshida et al. used 700 °C [18], Eddy
Fig. 8. Semilog plot of the mobility vs the inverse of the temperature for sample S950-A and its linear fitting.
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Fig. 9. Semilog plot of the charge carrier density vs the inverse of the temperature for sample S950-A and its linear fitting.
none of these studies are there further details about how the Hall measurements were performed at these low temperatures: neither on whether the sample was cooled down using a gas, or by placing it on a cold finger in vacuum. For the latter case, one of us has recently reported that the thermal glue used to place the sample on a sample holder can have a deep impact in the T-Hall measurements, especially below 100 K [37]. This could be a source of error in the measurements described and could explain the unusual linear dependence between the mobility and the temperature reported by Kusakabe et al. [27] and Ma et al. [30]. Anyhow, 1=2 as Ng et al. [20] have pointed out, the mobility has a n dependency, and thus, a thermally activated process of the form ΔE μ ≈ exp − 2kT
ð9Þ
where ΔE is the thermally activated energy for the charge carrier density. This is what we observed with Eqs. (7) and (8): the activation energy for the mobility is approximately half of the value of the activation energy for the charge carrier density. This suggests that in our sample S950-A, scattering is controlled by charge trapped in the dislocations, for temperatures higher than 150 K. Also, our sample S950-A showed a smaller density of islands (and thus, of dislocations). SEM studies suggest they are in the order of 1012 cm− 2. Our T-Hall results fit well to Fig. 6 of Ref. [16]. According to Fig. 3 of Ref. [16], this means that around 90% of the traps due to the dislocations are filled. Our study thus suggests that, for our undoped MOCVD GaN samples, the scattering model proposed by Weimann et al. explains well the temperature dependence of the mobility between 150 K and 373 K [16]. We comment on the similarity of our results with some of the MBE grown samples reported by Yoshida et al. [18] and Molnar et al. [22]. Yoshida et al. grew an AlN layer before the deposition of their GaN material [18]. This was not done by Molnar et al. [22]. However, the results of Molnar et al. for their samples 114 and 115 are similar to our samples S900-B and S950-A, respectively [22]. The same happens with Yoshida et al. sample grown with an AlN buffer layer and our sample S950-A [18]. Also, Yoshida et al. observed a decrease of the blue and yellow luminescence when using an AlN layer before the deposition of GaN [18]. This is similar to what we observed between our samples S900-A and S950-A. In other words, a similar effect to the one obtained by depositing an AlN layer prior to the deposition of undoped GaN by MBE growth can be obtained by MOCVD growth, doing the hydrogen and ammonia preliminary flow treatments, and increasing the TS from 900 °C to 950 °C. Finally, we briefly mention why our sample deposited with a TS of 950 °C (S950-A) is expected to have different properties than those
samples deposited with a TS of 900 °C (S900-A and S900-B). In their work, Khan et al. grew all their samples with a TS of 900 °C [42], and in particular, their sample 2 showed similar resistivity, charge carrier density and mobility to our samples S900-A and S900-B. On the other hand, Sasaki et al. reported samples grown at TS values between 950 °C and 1100 °C and 10,132 Pa [43]. They reported TMGa, NH3 and H2 flows with values between 20 and 40 μmol/min, 2.5 slm and 4.4 slm, respectively, which are larger than our growth conditions [43]. Another important difference is that they deposited first an AlN buffer layer. Nevertheless, they mentioned that their growth conditions formed truncated-island-like hillock. This is similar to what we observed in our sample grown with a TS of 950 °C (Figs. 3 and 4). In other words, with our research, we have shown that using smaller TMGa, NH3 and H2 flows, and no AlN buffer layer, we can obtain the same results as reported by Sasaki et al. [43]. It can be expected that the use of larger TS values like 1000 °C or 1050 °C, would yield larger mobilities and smaller charge carrier densities. This is what is usually observed as the growth temperature is increased [43]. TS values larger than 1100 °C should be avoided, as they cause thermal pitting on the GaN surface [43]. 6. Conclusions In this paper the electrical, optical and structural properties of undoped GaN films grown on sapphire substrates by LP-MOCVD, under different growth conditions, were reported. PL spectra of GaN films measured at RT showed that there is strong emission at 366 nm, which is related to the GaN band gap. In addition, weak blue emission, characteristic of MOCVD grown material, and emission in yellow range (550 nm) are seen. The surface morphology and the optical luminescence of GaN layers depend sensitively on the growth temperature. The layer grown with a TS of 900 °C shows a broad yellow luminescence originated from the defect levels, while at a larger growth temperature (with a TS of 950 °C) the quality of epitaxial GaN improves and the yellow emission is reduced. A TS of 950 °C reduces the yellow luminescence and increases the roughness of the surface, by the coalescence of islands. This supports the hypothesis that the yellow luminescence is related to dislocations or native point defects as these are reduced with the formation of islands. Also, it is only for the samples grown at 950 °C that some trend can be seen for the mobility and the charge carrier density. The temperature dependence of the mobility and charge carrier densities are well explained by the scattering mechanism due to charge trapped at threading dislocations. For material grown at 900 °C, the mobilities do not depend on temperature and their values are around 20 and 10 cm2 V− 1 s− 1, decreasing with increasing TMGa flow. The charge carrier densities are also temperature independent, and decrease from 5.7 × 1019 to 4.5 × 1019 cm− 3 as the TMGa flow increases. This is explained by a high density of dislocations (3 × 1012 cm− 2). For material grown at 950 °C, the islands coalesce and the dislocations are reduced. Still the dominant scattering mechanism is scattering by charge trapped at the threading dislocations, but with a thermally activated dependence for the mobility and charge carrier density. The activation energies are 2.15 and 5.41 meV, respectively. Our results confirm the scattering model by Weimann et al. [16]. Acknowledgements Victor-Tapio Rangel-Kuoppa gratefully acknowledges the Fonds zur Förderung der Wissenschaftlichen Forschung, Vienna, Austria and the National Council for Science and Technology (CONACyT) of Mexico, postdoctoral fellowship 78965. Victor Sánchez Reséndiz expresses his gratitude to the National Council for Science and Technology (CONACyT) from México, grant 56486. The authors would like to thank Elisabeth Pachinger, Alma Halilovic, Ursula Kainz, Daniel Ramirez, Miguel Avendaño, Rogelio
V.-T. Rangel-Kuoppa et al. / Thin Solid Films 519 (2011) 2255–2261
Fragoso, Eckehard Nusko, Otmar Fuchs, and Stephan Bräuer for technical assistance. The manuscript revision by Prof. Alberta Bonanni, Prof. Máximo López-López, Prof. Jonathan Farley and Dr. Andrea Navarro-Quezada is gratefully acknowledged. Special gratitude is expressed to Prof. Wolfgang Jantsch for his scientific support and advice. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13]
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