Electrical, optical, and structural properties of GaN films prepared by hydride vapor phase epitaxy

Journal of Alloys and Compounds 617 (2014) 200–206

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Electrical, optical, and structural properties of GaN films prepared by hydride vapor phase epitaxy A.Y. Polyakov a,b, N.B. Smirnov a,c, E.B. Yakimov d, A.S. Usikov e,f, H. Helava e, K.D. Shcherbachev b, A.V. Govorkov c, Yu N. Makarov e, In-Hwan Lee b,⇑ a

National University of Science and Technology MISiS, Moscow, Russia School of Advanced Materials Engineering and Research Center of Advanced Materials Development, Chonbuk National University, Jeonju, Republic of Korea Institute of Rare Metals, Moscow, Russia d Institute of Microelectronics Technology and High Purity Materials Russian Academy of Science, Chernogolovka, Russia e Nitride Crystals, Inc., Deer Park, USA f Saint-Petersburg National Research University of Information Technologies, Mechanics and Optics, Saint Petersburg, Russia b c

a r t i c l e

i n f o

Article history: Received 1 May 2014 Received in revised form 26 July 2014 Accepted 30 July 2014 Available online 7 August 2014 Keywords: Electrical compensation Residual donors Deep traps Hydride vapor phase epitaxy GaN

a b s t r a c t Two sets of undoped GaN films with the thickness of 10–20 lm were prepared by hydride vapor phase epitaxy (HVPE) and characterized by capacitance–voltage (C–V) profiling, microcathodoluminescence (MCL) spectra measurements, MCL imaging, electron beam induced current (EBIC) imaging, EBIC dependence on accelerating voltage, deep levels transient spectroscopy, high resolution X-ray diffraction measurements. The difference in growth conditions was mainly related to the lower (850 °C, group 1) or higher (950 °C, group 2) growth temperature. Both groups of samples showed similar crystalline quality with the dislocation density close to 108 cm 2, but very different electrical and optical properties. In group 1 samples the residual donors concentration was 1017 cm 3 or higher, the MCL spectra were dominated by the band-edge luminescence, and the diffusion length of charge carriers was close to 0.1 lm. Group 2 samples had a 2–4.5 lm thick highly resistive layer on top, for which MCL spectra were determined by green, yellow and red defect bands, and the diffusion length was 1.5 times higher than in group 1. We also present brief results of growth at the ‘‘standard’’ HVPE growth temperature of 1050 °C that show the presence of a minimum in the net donor concentration and deep traps density as a function of the growth temperature. Possible reasons for the observed results are discussed in terms of the electrical compensation of residual donors by deep traps. Ó 2014 Elsevier B.V. All rights reserved.

1. Introduction Hydride vapor phase epitaxy (HVPE) is the well-established growth technique for fabrication of high crystalline quality thick GaN films and crystals of polar and nonpolar orientations (see e.g. Refs. [1–7]). Currently it is the main method of preparation of native GaN substrates for epitaxial growth of GaN-based light emitting diode (LED) and injection laser diode (LD) structures [1–7]. The research in this direction has been concentrated on developing new ways of effective separation of GaN HVPE films from sapphire, preventing cracking and bowing of the free-standing wafers, and improving their crystalline quality. In recent years, serious progress has been achieved on this front owing to the creative use of epitaxial lateral overgrowth technique, the ⇑ Corresponding author. E-mail address: [email protected] (I.-H. Lee). http://dx.doi.org/10.1016/j.jallcom.2014.07.208 0925-8388/Ó 2014 Elsevier B.V. All rights reserved.

development of pit-producing buffer layers, and laser treatment of sapphire substrates (see e.g. Refs. [8–10]). As a result fabrication of good quality large diameter (up to 4 in.) GaN substrates became a reality. HVPE technique shows promise not only for preparation of substrates, but also it has been demonstrated to be useful for growth of actual device structures, such as LEDs, field effect transistors, high-voltage rectifiers and radiation detectors [11–13]. In such structures, much thinner HVPE-grown layers are employed and, for these applications, a good control of residual donors concentration and of deep traps spectra of the films is critical. However, little is as yet known about possible ways of obtaining this control. As a rule, in undoped HVPE GaN films, the concentration of uncompensated residual donors is around 1017 cm 3 or even higher, although several groups have demonstrated the possibility to reduce this level by about an order of magnitude or even lower [13–18].

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Undoped GaN films in this work were grown by HVPE on basal plane sapphire substrates. The growth process was performed in argon ambient with metallic Ga as a source material, and hydrogen chloride and ammonia as active gases. The growth was performed at temperatures of 850–950 °C that were somewhat lower than in a typical HVPE process of 1020–1040 °C [1–10]. The growth rate was 0.2–1 lm/min and the total thickness of the layers was 10–20 lm. The growth was initiated at low growth rate and then continued with increased growth rate. The two groups of samples studied in this paper mainly differed by the growth temperature. Group 1 samples were prepared at growth temperature of 850 °C, and group 2 samples were grown at 950 °C. In addition, we briefly present the results of characterization of the unintentionally doped n-GaN films grown at our standard HVPE growth temperature of 1050 °C (group 3). In all cases the sequence of growth rate changes and the actual values of the growth rate at each stage were kept the same, so that the growth temperature was the main variable. The shallow donors distribution and the deep traps spectra in the grown films were characterized by capacitance–voltage (C–V) profiling (HP4192A C–V meter for the 10–10 MHz frequency range), by deep levels transient spectroscopy with electrical (DLTS) and optical (ODLTS) injection (see e.g. Ref. [22]). For these measurements, a completely automated setup based on HP4280A C–V/C–t meter with external pulse generator HP8112A and the Oxford Instruments gas-flow cryostat was employed; at each point the entire capacitance relaxation curve was stored and the transients could be processed with any chosen set of time windows providing all the necessary information from one temperature scan. The grown films were also characterized by microcathodoluminescence (MCL) spectra measurements. The distribution of deep centers across the wafers and along the growth direction was determined by plan-view and cross-sectional monochromatic MCL imaging, using both band-edge luminescence and the luminescence in defect bands. These MCL measurements were carried out at room temperature in a JSM 6490 scanning electron microscope (Jeol) equipped with the MonoCL3 system (Gatan). The Hamamatsu photomultiplier was used as a detector. The spectral resolution depended on the signal intensity but in all measurements it was not worse than a few nm. In most cases the beam current of 0.1–1 nA was used for such experiments. The overall assessment of the deep defects presence and type was also done by electron beam induced current (EBIC) and by low temperature C–V measurements performed in the dark and after illumination. These low temperature C–V experiments were done to determine the concentration of deep electron traps with a barrier for capture of electrons and to estimate the density of deep hole traps located below the Fermi level [19–21]. Schottky diodes for these experiments and for C–V/DLTS studies were prepared by vacuum evaporation of Ni through a shadow mask (diode diameter typically 0.8 mm). Steady-state and pulsed illumination in C–V measurements and ODLTS was provided by high-power light emitting diodes (LEDs) with the peak wavelength ranging from 660 nm to 365 nm. Experimental setups were described previously [19–21]. EBIC experiments were used to determine the diffusion length Ld of non-equilibrium charge carriers. These experiments were performed using the JSM 840A scanning electron microscope, at zero bias applied to the sample. Two different methods were used. In the first, the dependence of collection efficiency of EBIC current on accelerating voltage of the probing electron beam was studied and the experimental dependence was fitted with the results of theoretical modeling using the Ld value as a fitting parameter. The method is described in some detail in Refs. [23,24], in which approximations and numerical values of the material parameters used are also discussed. The second method is based on linear scanning of the probing beam across the surface away from the edge of the Schottky diode and determining the diffusion length from the dependence of the EBIC signal on the distance to the diode edge (see e.g. Ref. [25]). Structural perfection of the films was studied by triple-axis high-resolution X-ray diffraction (HRXRD), by scanning electron microscope (SEM) imaging in the secondary electron (SE), monochromatic MCL, and EBIC imaging modes. HRXRD measurements were performed using D8 Discover (Bruker-AXS, Germany) tripleaxis diffractometer with Cu Ka radiation. Rocking curves for the symmetric GaN (0 0 0 4) reflection and asymmetric GaN (10–14) reflection were measured with widely open slits before the detector. For the symmetric (0 0 0 4) reflection, rocking curves were in addition measured with a narrow slit of 0.075 mm before the detector while scanning the reciprocal space in the direction normal to the vector of diffraction (qx) and parallel to the diffraction vector (qz). The diffraction intensity maps around the GaN (0 0 0 4) and GaN (10–14) nods in reciprocal space were also measured.

3.1. C–V profiling, DLTS, MCL spectra, and Ld measurements The electrical and optical properties of the samples of group 1 and group 2 were very different. Group 1 samples (growth temperature 850 °C) showed a high concentration of uncompensated shallow donors of 1017 cm 3 or higher in C–V profiles. For group 2 samples (growth temperature 950 °C), the top 2.5–4.5 lm of the films (the width different in different samples) were totally depleted at zero bias on Schottky diodes. Further down into the samples the C–V concentration was relatively constant with thickness. The concentration here varied from mid-1014 to 1016 cm 3 for various samples. Fig. 1 presents a typical concentration profile calculated for one of the group 2 samples. DLTS and ODLTS spectra of the samples of these two groups were also quite different. DLTS spectra are compared in Fig. 2. For one of the group 1 samples with the uncompensated shallow donors concentration of 1017 cm 3, the spectrum taken with reverse bias of 1 V and forward bias pulse of 0 V showed deep electron traps with activation energy 0.25 eV (ET1 traps, concentration 9.9  1013 cm 3), 0.35 eV (ET2 traps, 1.2  1014 cm 3), 0.6 eV (ET3 traps, 1.8  1014 cm 3) for our standard time windows settings. Fig. 2 shows the spectrum for time windows 100 ms/ 1000 ms. Also present there was a deeper trap with an activation energy of 1 eV (ET4, concentration of 1.9  1014 cm 3) that could be detected for longer time windows settings. Fig. 2 presents the spectrum taken with time windows 1000 ms/10,000 ms. The overall total deep electron traps concentration in this group 1 sample was 5.9  1014 cm 3. For one of the group 2 samples with the donor concentration 1016 cm 3 (Fig. 1) no signal from traps ET1 and ET2 was detected, traps ET3 and ET4 were still present (concentrations 6.4  1013 and 6.7  1013 cm 3, respectively) and an additional trap with the activation energy of 0.53 eV (ET5, 7.1  1013 cm 3) was observed (see Fig. 2) The aggregate concentration of the electron traps in the group 2 sample in Fig. 2 was about 3 times lower than for the group 1 sample (2  1014 cm 3 versus 5.9  1014 cm 3). At that, one has to bear in mind that (a) the traps concentrations in these two cases were sampled in the region 0.15 lm from the surface for the group 1 sample and 4.5 lm from the surface for the group 2 sample (generally the traps density decreases toward the surface, i.e. with increasing distance from the substrate) and (b) in other group 2 samples with lower donor concentration we could not detect any signal from electron traps, so the comparison in Fig. 2 represents the worst case scenario.

1.8x10

-3

2. Experimental

3. Results

Concentration (cm )

Although the concentration of deep electron traps in HVPE films can be made reasonably low, the density of deep hole traps is unacceptably high causing strong charge trapping in HVPE GaN-based devices (see e.g. Refs. [13,19–21]). In what follows we describe our work on understanding the ways to reduce the donor concentration and the deep traps density in undoped films of HVPE GaN.

1.6x10

1.4x10

1.2x10

1.0x10

16

16

16

16

16

4.75

4.76

4.77

4.78

4.79

4.80

Depth (µm) Fig. 1. Room temperature concentration profile of uncompensated donors for one of the group 2 samples.

A.Y. Polyakov et al. / Journal of Alloys and Compounds 617 (2014) 200–206

ODLTS spectra of the same group 1 sample as in Fig. 2 are presented in Fig. 3. The high-power LED with the peak wavelength of 365 nm was used, and the reverse bias was 0.5 V. With our usual time windows settings, hole traps with the activation energy of 0.6 eV and 0.9 eV could be observed. Fig. 3 (solid line) presents the spectrum taken with time windows of 100 ms/1000 ms. With longer time windows one more hole trap with the activation energy of 1.2 eV and hole capture cross section 10 12 cm2 could be detected (see the dashed line spectrum in Fig. 3 corresponding to time windows of 2500 ms/25,000 ms). The traps observed are quite similar to hole traps H3 (0.6 eV), H1 (0.9 eV), and H5 (1.1– 1.2 eV) reported previously for thick free-standing HVPE GaN crystals [13,21]. The H1 traps were attributed to gallium vacancy complexes with Si donors, and the H5 traps were ascribed to gallium vacancy complexes with oxygen donors (see Ref. [26]). For the group 2 sample from Fig. 2 the ODLTS spectrum showed the presence of hole traps similar to the H3, H1, and H5 traps, although for H5 traps the activation energy was slightly lower (1.1 eV versus 1.2 eV) and the hole capture cross section slightly smaller (10 13 cm2 versus 10 12 cm2) than for the group 1 sample. In addition, one more hole trap with the activation energy of 0.5 eV resembling the H2 hole traps of Ref. [13] was observed. As for comparing the concentrations of the hole traps in group 1 and group 2 samples it is not quite straightforward. For the group 1 sample, the situation is reasonably clear. If one assumes that the light pulse fully recharges the deep hole traps throughout the entire space charge region and takes the shallow donor concentration as 1017 cm 3 from C–V profiling one gets the concentrations of the three detected hole traps as 3  1014, 5  1014, and 5  1014 cm 3, respectively. The total concentration of deep hole traps lying below the Fermi level in the space charge region can also be determined from the persistent decrease of the voltage intercept of the low temperature (85 K in our case) 1/C2 versus voltage plots after illumination. This decrease is caused by the hole capture by the deep hole traps during illumination. The non-equilibrium increase in the built-in charge cannot be removed at low temperatures other than by passing the forward current that restores the initial voltage intercept (the thermal emission of holes from the deep traps is too slow at low temperature). The overall concentration of deep hole traps can be calculated from the voltage intercept decrease as described in Refs. [19–21]. We do not have space to show the actual data, but the concentration of hole traps thus obtained is 1.5  1015 cm 3, in good agreement with the sum of hole traps concentrations in ODLTS measurements.

0.25

DLTS signal (Arb. units)

ET5 0.20

ET4

ET3

0.15

Group 2 0.10 0.05

ET4

ET3 ET2

ET1

Group 1 0.00

100

150

200

250

300

350

400

Temperature (K) Fig. 2. DLTS spectra measured for one of the group 1 samples with reverse bias 1 V, forward bias pulse of 0 V and time windows 100 ms/1000 ms (solid line, red) and 1500 ms/15,000 ms (dashed line, blue); also shown the DLTS spectra measured for one of the group 2 samples at 1 V, with the forward bias pulsed to 0 V, with the time windows 100 ms/1000 ms (solid line, blue) and 1500 ms/15,000 ms (dashed line, blue). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

0

ODLTS signal (Arb. units)

202

group 1(x3)

-1

H3

-2

H1 H1

H5

-3 -4

H3 H2

-5

H5

H1 -6 -7

50

100

150

200

250

300

350

400

Temperature (K) Fig. 3. ODLTS spectra measured for one of the group 1 samples at 0.5 V, 365-nm wavelength LED excitation and time windows 100 ms/1000 ms (solid line, red) and 2500 ms/25,000 ms (dashed line, red); also shown is the ODLTS spectrum measured for one of the group 2 samples at 0.5 V, with time windows 500 ms/5000 ms, with 365-nm wavelength LED excitation (solid line, blue). (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

For the group 2 sample the situation is more complicated because the deep hole traps are recharged in the top highly compensated layer where we do not know the concentration of residual uncompensated donors. But from Fig. 3 it is obvious that the hole traps density in this sample is considerably lower than in the group 1 sample, even if one takes for the shallow donor concentration the concentration from the C–V profile in Fig. 1, which clearly leads to overestimating the hole traps density. Electrical measurements on the sample grown at our standard temperature of 1050 °C (group 3) gave the net shallow donors concentration from C–V profiling as 9.7  1016 cm 3. The dominant electron traps were the ET3 (0.6 eV) traps with the concentration of 4.2  1014 cm 3 and the ET4 (1 eV) traps with concentration of 1.4  1015 cm 3. The hole traps spectra were dominated by the H1 (0.9 eV) traps with the concentration of 9.6  1015 cm 3 and H5 (1.1 eV) traps with the concentration of 1.9  1016 cm 3. Thus, both the net donor concentration and the density of electron and hole traps showed a clear minimum for the growth temperature of 950 °C. Room temperature MCL spectra taken with the probing beam accelerating voltage 10 kV and probing beam current of 10 9 A are compared for the representative group 1 and group 2 samples in Fig. 4 (these are the same samples as the ones in Figs. 1–3). It can be seen that, for group 1 samples, the spectra were dominated by the band-edge line at 3.35 eV, with minor contributions from the defect bands peaked at 2.5 eV (green band, marked as GB in the figure), and a broad peak with the high energy and low energy shoulders that can be deconvoluted into the band peaked at 2.25 eV (yellow band, marked as YB in the figure) and the band peaked at 1.8–1.9 eV (red band, marked as RB). For the group 2 sample, the MCL spectrum was dominated by the defect green, yellow, and red bands, with the band-edge line intensity quite low (Fig. 4). At that, the integrated MCL intensity from all peaks was not significantly different for the group 1 and group 2 samples suggesting that one deals here with re-pumping of the MCL intensity from the band-edge line into the defect bands lines rather than with the suppression of the band-edge luminescence due to the increased impact of non-radiative recombination. This re-pumping comes most likely from the much lower density of shallow donors in the group 2 sample. MCL spectra for the sample grown at 1050 °C were very similar to the spectra obtained for group 1 samples. Comparing the MCL images of the cleaved group 1 and group 2 samples helps to better understand the above results. For the

A.Y. Polyakov et al. / Journal of Alloys and Compounds 617 (2014) 200–206

group 1 sample, the image taken for the band-edge luminescence band (365 nm) is shown in Fig. 5(a). The sapphire substrate is at the bottom, the bright line near the substrate relates to the ‘‘buffer’’ layer, approximately 1 lm-thick, grown at the start of the process with a low growth rate of 0.2 lm/min. This layer has a high density of residual donors and is responsible for the lowering of the apparent electron mobility in Hall measurements in HVPE GaN, which is well documented (see e.g. Ref. [15]). The increased bandedge MCL intensity is a natural consequence of the higher residual donors density in this ‘‘buffer’’ layer. This region is followed by the region, approximately 3 lm-thick, with a high density of dark defects that are usually attributed to dislocations. Above that layer is the uniformly bright region interspersed by dark vertical lines caused by threading dislocations. MCL imaging of the cleaved sample using the defect bands lines resulted in a dark more or less featureless picture due to the low intensity of the defect bands. The actual image is not presented here to save space, but can be found in the Supplementary material as Fig. S1(b) alongside the image taken in the secondary electrons observation mode, Fig. S1(a). For the group 2 sample, the images taken using the band-edge 365 nm line and the major defect band 510 nm line are presented in Fig. 5(b) and (c), respectively. The band-edge image is similar to the image of the group 1 sample in Fig. 5, showing the bright ‘‘buffer’’ layer (approximately 1 lm-thick), the about 3 lm-thick highdislocation density layer near the substrate, and the regular film with dark, mainly vertical, dislocation lines above. However, the important difference here is that the top approximately 4 lm-thick part of the film shows a much lower band-edge line intensity due to the lower concentration of residual donors. When mapped using the dominant defect band line at 510 nm this region looks brighter than the rest of the film (Fig. 5(c); Fig. S2(a) in the Supplementary material shows the cross section image of sample 2 taken in the secondary electrons observation mode alongside the image taken using the 510 nm MCL band; mind that the intensity of the defect band is in this case also relatively high in the near-substrate region with high dislocation density (the substrate is at the bottom): the contrast compared to sample 1 imaged in the same regime is obviously magnified owing to the low residual donor density). The thickness of the compensated layer in the MCL images in Fig. 5(b) and (c) is reasonably close to the width of the fully depleted at 0 V heavily compensated region obtained from C–V profiling (see Fig. 1), although they do not have to be exactly the same for obvious reasons. Measurements of the diffusion length from the EBIC signal dependence on the accelerating voltage of the probing electron 4

Bandedge MCL intensity (Arb. units)

GB YB

3

group 1 2

RB

group 2

1

0 1.5

2.0

2.5

3.0

3.5

203

beam gave the values of Ld equal to 0.1 lm for the group 1 sample and 0.15 lm for the group 2 sample (the material parameters used in fitting were taken from Ref. [24]; we do not show the actual data to save space). Measurements of diffusion length by the alternative method, i.e. by doing a linear scan of the EBIC signal while moving to or from the edge of the Schottky diode (see e.g. Ref. [25]) gave the similar relation between the diffusion lengths of the group 1 and group 2 samples, but an order of magnitude higher actual Ld values (1 lm for the group 1 sample, 1.5 lm for the group 2 sample). Such a strong discrepancy could be owing to (a) the secondary excitation caused by the luminescence induced by the probing beam and (b) the band bending at the surface caused by the Fermi level pinning. For case (a) one expects the apparent ‘‘diffusion length’’ to increase with increasing the probing beam accelerating voltage, both because of the effective increase of the footprint of the beam excitation volume projected to the surface and the increase of the secondary MCL intensity. This was not happening and the actual effective Ld values were the same for the probing beam accelerating voltage of 5 kV and 35 kV. This favors the (b) case above. This case will be realized if the Fermi level is pinned at the surface even without the Schottky diode. The resulting band bending causes the spatial separation of electrons and holes produced by the probing beam even far from the Schottky diode edge and increases the effective Ld value. Similar phenomenon has been observed by us for undoped GaN films grown by epitaxial lateral overgrowth (ELOG) technique [23]. 3.2. Comparison of the structural perfection of the group 1 and group 2 samples The results of HRXRD measurements for the (0 0 0 4) and (10– 14) rocking curves are presented in Table 1. These results indicate similar overall crystalline perfection of the group 1 and group 2 films. The density of screw and edge dislocations estimated from the FWHM values for the (0 0 0 4) reflection and the (10–14) reflection (see e.g. Ref. [27]) were around 5  107 cm 2 and 5  108 cm 2, respectively. As seen from the band-edge MCL images of the cleaved samples in Fig. 5(a) and (b) the density of threading dislocations counted as the density of dark line defects varies with thickness. It is high near the substrate and is close to 108 cm 2 near the surface for both sample types. Plan-view EBIC images measured with different accelerating voltages also confirm the different density of dark spots usually associated with threading dislocations. Fig. 6(a) and (b) shows the EBIC images for the group 2 sample for the probing beam accelerating voltage of 10 kV and 35 kV, respectively. It can be clearly seen that the dislocation density increases from about 108 cm 2 for 10 kV (the probed depth of about 0.75 lm) to about 5  108 cm 2 for the accelerating voltage of 35 kV (the probed depth of about 6.5 lm, see e.g. Ref. [23]). Very similar results were obtained for the group 1 sample. In HRXRD, the effective dislocation density is averaged over the excitation volume, which perhaps explains the observed discrepancy in the absolute values of the dislocation densities estimated by the two techniques. For the sample grown at 1050 °C the FWHM for the (0 0 0 4) reflection was 0.075 degrees. For the asymmetric (10–14) reflection the FWHM was 0.082 degrees. The dislocation density in the surface region determined from MCL and EBIC imaging was close to 108 cm 2. Thus, the structural quality of the films grown in the studied growth temperature range did not vary significantly.

Photon energy (eV) Fig. 4. Room temperature MCL spectra for two typical group 1 (red line) and group 2 (blue line) samples; accelerating voltage of the probing beam 10 kV, beam current 1 nA. (For interpretation of the references to colour in this figure legend, the reader is referred to the web version of this article.)

4. Discussion We have shown that, under our growth conditions, HVPE samples grown at 950 °C show a much lower density of

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Fig. 5. (a) Cross section 365 nm MCL line image of one of the group 1 samples; (b) cross section 365 nm MCL line image of one of the group 2 samples; (c) cross section 510 nm MCL line image of one of the group 2 samples; the scale bar on all figures corresponds to 5 lm, the accelerating voltage is 10 kV.

uncompensated shallow donors, a lower concentration of deep electron and hole traps and a measurably higher diffusion length of non-equilibrium charge carriers than the samples grown at 850 °C (and 1050 °C). The residual donor background in HVPE is mainly determined by contamination with Si and O shallow donors (see e.g. [2,7]) and by the density of native defects producing shallow donors, primarily nitrogen vacancies (VN). The level of Si and O concentration is the function of the presence of these impurities in the gas mixture near the growing surface and of the incorporation efficiency from the gas phase which is generally governed by the availability of suitable vacant sites in the lattice and the site competition considerations (i.e. the O donor impurity in the gas phase competes for the place in the N sublattice with N atoms, while the Si donors compete with Ga atoms for incorporation in the Ga sublattice). The concentrations of impurities in the gas phase are not expected to vary strongly in the gas phase for the growth temperatures range studied in this paper. However, the effective V/III components ratio near the surface should increase with increasing the growth temperature owing to the more complete decomposition of ammonia near the surface. This should decrease the incorporation of oxygen donors (and shallow donors due to nitrogen vacancies), but can simultaneously promote the formation of Ga vacancies (VGa) and thus the incorporation of Si donors. The latter are responsible for the formation of deep acceptor centers H1 near Ev + 1 eV either by themselves or when complexed with shallow donor impurities. Complexes of VGa with oxygen are responsible for the H5 hole traps with the acceptor level near Ev + (1.1–1.2) eV. These gallium-vacancyrelated acceptors play a prominent role in the conductivity compensation in undoped GaN films (see e.g. a review in Ref. [28] and references therein). The overall density of these deep

acceptor traps is determined by both the densities of the shallow donors and the gallium vacancies. Clearly, these conflicting tendencies should result in the concentration of uncompensated donors showing a minimum as a function of growth temperature. In our growth conditions this minimum is situated close to 950 °C. For higher (1020–1050 °C) and lower (850 °C) temperatures the residual donor density increased to about 1017 cm 3 or higher due to more efficient Si incorporation at higher temperature or oxygen incorporation and enhanced VN concentration at lower temperature. The major deep acceptors detected in our ODLTS spectra, the H1 and H5 hole traps, have been attributed to complexes of VGa with, respectively, Si and O [13,21,26]. Hence, it is expected that their densities should also show a minimum with the growth temperature, in our case approximately coinciding with the minimum in the donor concentration dependence on temperature. These centers are known to give rise to the major defect bands at 2.5 eV, 2.3 eV, and 1.9 eV that dominate the defect luminescence spectra in our samples and similarly grown HVPE samples [13,21,26,29]. Some of the dominant electron traps in undoped GaN have also been associated with complexes involving shallow donors [26] and thus it is not surprising that we see the overall density of deep electron traps to be lower in the group 2 samples with the lowest donor density. The presence of heavily compensated layers in the top portion of the group 2 samples is most likely related to the lower donor background level that tends to decrease toward the surface becoming at some point comparable to the density of compensating deep acceptors. The depth corresponding to the cross-over point of switching from conducting to high-resistivity parts of the films is determined by the interplay between the density of shallow

Fig. 6. (a) Plan-view EBIC image of one of the group 2 samples taken with accelerating voltage of 10 kV; (b) plan-view EBIC image of one of the group 2 samples taken with accelerating voltage of 35 kV.

A.Y. Polyakov et al. / Journal of Alloys and Compounds 617 (2014) 200–206 Table 1 Full widths at half maximum (FWHM) for symmetric (0 0 0 4) and asymmetric (10–14) reflections obtained for group 1 and group 2 samples. Sample

FWHM (0 0 0 4) qx (degrees)

FWHM (0 0 0 4) qz (degrees)

FWHM (0 0 0 4) fully open detector window (degrees)

FWHM (10–14) fully open detector window (degrees)

Group 1 Group 2

0.065

0.020

0.067

0.073

0.067

0.024

0.072

0.079

donors and the density of compensating deep acceptors whose concentration in turn depends on donor concentration. Generally, we observed that, for similar total layer thickness in group 2 samples, the thickness of the top high-resistivity layer was lower when the residual donor concentration decreased (it decreased from 4.5 lm for the donor concentration in the conducting part close to 1016 cm 3 to 2.5 lm for the donor concentration of 1015 cm 3). For group 1 (850 °C) and group 3 (1050 °C) samples the formation of the heavily compensated layer was not observed because the residual donor density was much higher while the concentration of the compensating deep acceptors was limited by the availability of vacancies. Since the density of deep traps never exceeded 1016 cm 3 they were not contributing significantly to compensation in the films with the shallow donors density of 1017 cm 3 or higher. Previously reported results of the diffusion length dependence on dislocation density in HVPE GaN (see e.g. Ref. [30]) were explained by the dislocations being the dominant nonradiative recombination centers, so that the diffusion length value was determined by the distance between the dislocations or the average dimensions of the dislocation cells in the samples where dislocations formed a cellular structure. However, the dislocation densities in our group 1 and group 2 samples were quite similar, yet the diffusion length was measurably higher in group 2 films. This suggests a sizeable impact of recombination via deep traps associated with point defects. They are in our case most likely the ET3 and ET4 electron traps, and the H1 and H5 hole traps with high capture cross sections and relatively high concentration. The results of our structural studies indicate that the density of extended defects in HVPE is not strongly affected by changing the growth temperature from 850 °C to1050 °C, although, for lower growth temperatures, we observed serious deterioration of the structural quality. 5. Conclusions We have shown that, in HVPE growth, an optimum growth temperature exists that corresponds to a minimum in both the concentration of residual donors and the density of deep electron and hole traps. The existence of such a minimum in donor concentration is related to the site competition between the incorporation of the major contaminants, oxygen and silicon, and of constituent atoms, respectively, gallium and nitrogen. For the deep traps the minimum is due to the fact that the major deep centers in GaN often are complexes between the native defects and the shallow donors impurities. In our conditions the minimum in question corresponded to the growth temperature of about 950 °C. At this temperature the residual donor concentration in the portion of the films near the middle of the samples could be reduced to the level of mid-1014–1016 cm 3, while the top portions of the films were highly compensated because, in this region, the concentration of deep compensating acceptors became close to that of shallow donors. The thickness of the top compensated layer decreased with decreasing the residual donor concentration. At the same time, the results of structural characterization show that these HVPE films

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are not inferior in crystalline quality to the HVPE films grown at standard growth temperatures of 1020 °C–1050 °C. Therefore, with further decreasing the level of contamination and further optimization one can get the electrically uniform thick films with very low concentration of background donors and deep traps. Preliminary studies show that it is possible to grow this way HVPE films with the density of uncompensated donors of some 1014 cm 3 and the density of compensating deep hole traps in the mid1013 cm 3 range [31]. This is very promising for applications in high-power rectifiers and in GaN radiation detectors. Acknowledgments The work at NUST MISiS was supported in part by the Ministry of Education and Science of the Russian Federation in the framework of Increase Competitiveness Program of NUST «MISiS» (No R2-2014-055). The work at Chonbuk National University was supported by National Research Foundation of Korea (NRF) funded by Ministry of Science, ICT & Future Planning (2013R1A2A2A070 67688, 2010-0019626). Appendix A. Supplementary material Supplementary data associated with this article can be found, in the online version, at http://dx.doi.org/10.1016/j.jallcom.2014. 07.208. References [1] K. Fujito, S. Kubo, H. Nagaoka, T. Mochizuki, H. Namita, S. Nagao, Bulk GaN crystals grown by HVPE, J. Cryst. Growth 311 (2009) 3011–3014. [2] V. Dmitriev, A. Usikov, in: Z.C. Feng (Ed.), III-Nitride Semi-conductor Materials, Imperial College Press, London, 2006, pp. 1–40. [3] K. Fujito, S. Kubo, I. Fujimura, Development of bulk GaN crystals and nonpolar/ semipolar substrates by HVPE, MRS Bull. 34 (2009) 313. [4] R. Kroger, T. Paskova, B. Monemar, S. Frigge, D. Hommel, A. Rosenauer, Phys. Status Solidi c 4 (2007) 2564. [5] T. Yoshida, Y. Oshima, T. Eri, K. Watanabe, M. Shibata, T. Mishima, Phys. Status Solidi a 205 (2008) 1053. [6] M.L. Reed, E.D. Readinger, H. Shen, M. Wraback, A. Syrkin, A. Usikov, O. Kovalenkov, V.A. Dmitriev, Appl. Phys. Lett. 93 (2008) 133505. [7] A. Usikov, V. Soukhoveev, O. Kovalenkov, A. Syrkin, L. Shapovalov, A. Volkova, V. Ivantsov, Jpn. J. Appl. Phys. 52 (08J) (2013) B22. [8] A. Usui, Gallium nitride crystals grown by hydride vapor phase epitaxy with dislocation reduction mechanism, ECS, J. Solid State Sci. Technol. 2 (8) (2013) N3045. [9] K. Yamano, M. Ueno, H. Furuya, N. Okada, K. Tadamoto, Successful natural stress-induced separation of hydride vapor phase epitaxy-grown GaN layers on sapphire substrates, J. Cryst. Growth 358 (2012) 1. [10] T. Sato, S. Okano, T. Goto, T. Yao, R. Seto, A. Sato, H. Goto, Nearly 4-inchdiameter free-standing GaN wafer fabricated by hydride vapor phase epitaxy with pit-inducing buffer layer, Jpn. J. Appl. Phys. 52 (2013) 08JA08. [11] S. Kurin, A. Antipov, I. Barash, A. Roenkov, H. Helava, S. Tarasov, E. Menkovich, I. Lamkin, Yu. Makarov, CHVPE growth of AlGaN-based UV LEDs, Phys. Status Solidi c 10 (2013) 289–293. [12] D. Tsvetkov, Yu. Melnik, A. Davydov, A. Shapiro, O. Kovalenkov, J.B. Lam, J.J. Song, V. Dmitriev, Growth of submicron AlGaN/GaN/AlGaN heterostructures by hydride vapor phase epitaxy, Phys. Status Solidi a 188 (1) (2011) 429. [13] In-Hwan Lee, A.Y. Polyakov, N.B. Smirnov, A.V. Govorkov, E.A. Kozhukhova, V.M. Zaletin, I.M. Gazizov, N.G. Kolin, S.J. Pearton, Electrical Properties, Deep trap spectra and radiation detector performance of free-standing bulk n-GaN with a thick compensated surface region, J. Vac. Sci. Technol. B 30 (2) (2012) 021205. [14] J. Jasinski, W. Swider, Z. Litiental-Weber, P. Visconti, K.M. Jones, M.A. Reshchikov, F. Yun, H. Morkoç, S.S. Park, K.Y. Lee, Appl. Phys. Lett. 68 (2001) 2297. [15] Z.-Q. Fang, D.C. Look, J. Jasinski, M. Benamara, Z. Liliental-Weber, R.J. Molnar, Appl. Phys. Lett. 78 (2001) 332. [16] K. Fujito, Kiyomi, T. Mochizuki, H. Oota, H. Namita, S. Nagao, I. Fujimura, Phys. Status Solidi a 205 (2008) 1056. [17] N. Okada, M. Imura, T. Nagai, K. Balakrishnan, M. Iwaya, S. Kamiyama, H. Amano, I. Akasaki, H. Maruyama, T. Noro, T. Takago, A. Bandoh, Phys. Status Solidi c 4 (2007) 2528. [18] P.R. Hageman, V. Kirilyuk, W.H. Corbeek, J.L. Weyher, B. Lucznik, M. Bochkowski, S. Porowski, S. Muller, Thick GaN layers grown by hydride vapor phase epitaxy: hetero-versus homo-epitaxy, J. Cryst. Growth 255 (2003) 241.

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