Growth of High Crystalline Quality HVPE-GaN Crystals with Controlled

Article pubs.acs.org/crystal

Growth of High Crystalline Quality HVPE-GaN Crystals with Controlled Electrical Properties J. A. Freitas, Jr.,*,† J. C. Culbertson,† N. A. Mahadik,† T. Sochacki,‡ M. Bockowski,‡ and M. Iwinska‡ †

Naval Research Laboratory, Washington, D.C. 20375, United States Institute of High Pressure Physics, Polish Academy of Science, 01-142 Warsaw, Poland

Downloaded by CENTRAL MICHIGAN UNIV on September 15, 2015 | http://pubs.acs.org Publication Date (Web): September 14, 2015 | doi: 10.1021/acs.cgd.5b00617

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ABSTRACT: Thick free-standing hydride vapor phase epitaxy (HVPE) GaN wafers were grown with close to perfect crystalline quality and with more than two orders of magnitude reduced free carrier concentration relative to ammonothermal GaN without resorting to deep level compensation. These new homoepitaxially grown crystals exceed the quality of any commercially available GaN substrates. High crystalline quality is confirmed by XRD rocking curve measurements having full width at half-maximum of only 16 arc s. Detailed Raman scattering spectroscopy and imaging show that the wafers are biaxially stress-free and have a uniform and low background doping. The stress-free, flat nature of these substrates is essential for high-yield device fabrication. High-resolution photoluminescence studies in combination with SIMS analysis identify Si and O as residual shallow donors with concentration levels orders of magnitude lower than those present in ammonothermal GaN substrates.

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INTRODUCTION Three decades of GaN semiconductor research have resulted in the realization of a number of optoelectronic and electronic devices that have impacted device technology and affected everyday life. Despite the fast development of GaN-based devices, mostly resulting from the unique combination of intrinsic GaN properties, growth of real bulk III−V nitrides has progressed at a much slower pace. About 5 years ago an ability to grow high crystalline quality bulk GaN crystals using the ammonothermal method was reported and substrates produced by this method became readily available.1 Ammonothermal GaN (Am-GaN) wafers commonly have a very high concentration of free electrons (typically >5 × 1018/cm3), are very flat, and have a very low dislocation density (typically ≤5 × 104/cm2).2 Am-GaN typically has a high concentration of gallium vacancies.3 To realize high yields of high-performance optoelectronic and electronic devices, full control of the substrate’s electrical properties is also required. To obtain semi-insulating ammonothermal wafers, it is necessary to reduce the free carrier concentration caused by the dominant shallow impurity (oxygen donors) by adding deep acceptor impurities. Thus, the number of foreign atoms in the wafer increases, changing the lattice constants and the thermal conductivity of the material. It has been reported that hydride vapor phase epitaxy (HVPE), a fast quasi-bulk growth technique, can be used to deposit thick films with relatively low free carrier background concentrations.4 Therefore, it seems that a combination of the HVPE technique (higher growth rate and higher purity) and the ammonothermal method (higher structural quality) could introduce a new approach to grow low defect density GaN with controllable © XXXX American Chemical Society

electronic properties. Sochacki et al. recently demonstrated that crack-free, high crystalline quality GaN with reduced free carrier concentrations can be deposited homoepitaxially on Am-GaN using the HVPE technique.5 However, it has not been demonstrated that free-standing wafers obtained from such crystals can be used as seeds for crystal growth or as substrates for epitaxial growth. The present work addresses these important issues, showing that thick HVPE-GaN previously deposited on an Am-GaN substrate can be processed like bulk crystals and used to grow high crystalline quality wafers with improved electrical transport properties. GaN substrates having such properties are required for the realization of a number of high-performance devices (e.g., high-voltage fast switches).6,7

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EXPERIMENTAL DETAILS

A home-built horizontal quartz HVPE reactor, described in detail elsewhere,8 was used as a growth apparatus. GaCl was supplied vertically over the surface of the seed with a showerhead type quartz nozzle, while NH3 was supplied by a second quartz nozzle located a few millimeters away from the susceptor and on the same level as the seed’s surface. The temperature of GaN crystallization was about 1045 °C, while the temperature of GaCl synthesis was about 870 °C. The growth was carried out with a HCl flow of 24 mL/min diluted in H2 and a NH3 flow of 480 mL/min diluted in H2 with H2 as the carrier gas. High-resolution X-ray diffraction (HR-XRD) rocking curves were measured using a Rigaku Smartlab diffractometer, equipped with a 9 kW Cu rotating anode, a collimating mirror, and a four bounce Ge (220) channel cut monochromator, which provides a low divergence Received: May 5, 2015 Revised: August 13, 2015

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DOI: 10.1021/acs.cgd.5b00617 Cryst. Growth Des. XXXX, XXX, XXX−XXX


Crystal Growth & Design

Article

pure Cu Kα1 incident beam. The sample was precisely aligned using a computer-controlled, Z and χ, θ stage. Symmetric and asymmetric reflections were measured to precisely obtain the a and c lattice parameters. Values of full width at half-maximum (FWHM) of the rocking curves were used to determine the crystalline quality of the wafers. The slit widths were 1 mm on both the incident and analyzing sides in order not to artificially reduce the intrinsic sample FWHM. Surface topography was measured using a Bruker-Nano Dimension FastScan AFM instrument with a closed-loop xyz scanning head. Confocal micro-Raman measurements were carried out using a single-mode 488 nm laser, a half-meter Acton spectrometer with an 1800 grooves/mm holographic grating, and a Princeton Instruments back-thinned, deep depleted, nitrogen cooled CCD (1340 × 400 pixel array). The spectral resolution of this configuration of the system is 2.5 cm−1, and the repeatability with which a line position can be determined is within 0.1 cm−1. The laser power was attenuated using neutral density filters. Sample excitation with the 488 nm laser was made confocal with the detection axis using a SEMRock dichroic beam splitter. The laser was focused and the Raman collected using a 50× Mitutoyo microscope objective having a numerical aperture of 0.65. Two SEMRock long-pass filters were placed in front of the spectrometer to filter out laser light. Samples were mounted on a precision, computer-controlled Aerotech XYZ translator having a bidirectional position accuracy of better than 0.1 μm. Raman spectra and the incident and reflected laser power were measured and stored for each position in each spatial map. Low-temperature (∼1.5−6 K) photoluminescence measurements were carried out with the samples placed in a continuous liquid He flow cryostat with a temperature variation capability between 1.5 and 300 K. Luminescence was excited with the 325 nm line of a He−Cd laser with a typical incident power between 0.5 and 2.0 mW. Neutral density filters were used to maintain the incident power within desired limits to avoid heating the samples and to allow the observation of recombination processes with different recombination times. The emitted light was dispersed by an 1800 grooves/mm 0.85 m doublegrating spectrometer and detected by a UV-sensitive GaAs photomultiplier coupled to a computer-controlled photon counter system.

the second. This difference was due to small changes in the geometry of the reactor and the distance between GaCl nozzles and the susceptor with the reagents’ flows remaining constant. The 2300 μm thick boule was sliced into three pieces: one sample having 130 μm of HVPE-GaN on a 600 μm thick AmGaN substrate (S) and two samples of ∼600 μm thick freestanding HVPE-GaN (W2 and W1, top and bottom, respectively), schematically represented in Figure 1a. Both

Figure 1. (a) Scheme of slicing of the homoepitaxial HVPE- GaN crystal into three parts: S (730 μm thick seed: 600 μm of Am-GaN and 130 μm of HVPE-GaN) and W1 and W2 (free-standing 600 μm thick HVPE-GaN). (b and c) Free-standing 340 μm thick HVPE-GaN crystals “A” and “B”, which were diced from the free-standing W2 wafer: (b) sample “A” with two overgrown pinholes visible. Grid line, 1 mm.

faces of the W1 and W2 wafers were mechanically polished. The W2 wafer was diced into small pieces. Ga-face and N-face of two selected crystals (samples “A” and “B”, depicted in Figure 1b) were prepared by mechanochemical polishing to an epi-ready state. The thicknesses of A and B samples were 340 μm. The surface morphologies of the Ga-polar and N-polar faces of samples A and B were examined by tapping mode atomic force microscopy (AFM) after their surfaces were mechanically lapped, mechanochemically polished, and cleaned. Figure 2a shows the topography of the Ga-polar face of sample A. The surface roughness (RMS) of the Ga-polar face is only 0.141 nm and clearly shows atomic steps, consistent with a very high quality surface preparation. The N-polar face of this sample, depicted in Figure 2b, has a much higher roughness (RMS = 1.01 nm), clearly indicating that an improved chemical etching procedure must be developed for preparation of the (0001̅) surface to the epi-ready state. Due to the strong chemical activity of the N-polar GaN surface it is much more difficult to prepare it to the epi-ready state than the Ga-polar face. The Gaface is much more inert and can be polished very uniformly on the entire polished surface. Figure 3 shows XRD rocking curves of the (004) reflection obtained from the Ga-polar faces of samples A and B, using a wide slit size of 0.5 mm. The FWHMs for both samples were only 16 arc s. These values are similar to the HR-XRD rocking curves previously reported for high crystalline quality Am-GaN.2 Such a small value of FWHM can be observed only on high crystalline quality materials characterized by very low dislocation density.2 Selective etching experiments carried out on similar HVPE epitaxial layers deposited on Am-GaN substrates show that the dislocation density is typically lower than 5 × 104/cm2.9 These HR-XRD results clearly indicate that the high crystalline quality of the Am-GaN substrates is preserved for the HVPE-GaN layers after multiple growths. All eight of the zone-center optical modes of the hexagonal phase of GaN (wurtzite structure belonging to the space group

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RESULTS AND DISCUSSION Two crystallization runs were carried out to verify that the high crystalline quality of ammonothermal substrates could be reproduced using the fast chemical vapor deposition method, with a simultaneous improvement of the material’s optical and electronic properties. GaN was deposited using the HVPE method on a 600 μm thick 1 in. Am-GaN substrate (seed) having a misorientation of 0.98 ± 0.2° toward the m-direction. The Am-GaN had a free carrier concentration ∼1 × 1018 cm−3. Before the growth, the (0001) surface of the seed was prepared by mechanochemical polishing and subsequently cleaned. The high structural quality of the Am-GaN seed was confirmed by a 50 arc s (instrument limited) X-ray rocking curve FWHM for the (002) reflection. The 50 arc s FWHM measurement was performed on a triple axis system having a higher divergence, as compared to the Rigaku Smartlab XRD system described previously that was used to evaluate the free-standing HVPE samples. The a and c lattice constants were 3.1892 ± 0.0002 and 5.1851 ± 0.0001 Å, respectively. A 1400 μm thick homoepitaxial HVPE-GaN layer with a few hillocks on the (0001) surface was obtained after a 9 h long growth run. The (0001) surface was then mechanically and mechanochemically polished and cleaned. This sample, used as a substrate for the second run, comprised of 900 μm of HVPEGaN on 600 μm of Am-GaN, was placed in the HVPE reactor to grow an additional 800 μm thick HVPE-GaN layer in 8 h. Thus, two homoepitaxial HVPE-GaN layers were grown with two different rates: 160 μm/h for the first and 100 μm/h for B




Article

Figure 2. AFM topography of the free-standing homoepitaxial HVPE-GaN, sample A, after removing it from the HVPE-GaN/Am-GaN template subsequent CMP and cleaning; (a and b) micrographs of the Ga-polar and N-polar faces, respectively. Scan area, 500 nm.

Figure 4. First order Raman spectra of the Ga-polar and N-polar faces of sample B. Note the shift of the A1(LO) phonon to higher frequency. Spectral details are discussed in the text.

Figure 3. Rocking curves of the (004) reflection acquired at Ga-polar faces of both free-standing HVPE homoepitaxial GaN samples, A and B.

concentration on the Ga-polar than that on the N-polar faces of the samples. To characterize the areal and volumetric uniformity of the samples, spatial maps of Raman spectra for the Ga- and N-polar faces of samples A and B were measured. Spatial maps of the Raman shift (k) of the E22 phonon across a 2 mm × 2 mm area (scan steps of 100 μm) of the sample B are represented in Figure 5a,b (Ga-polar and N-polar, respectively). It should be pointed out that the E22 phonon frequency is essentially the same across the whole area on both sides of the samples. This indicates that the sample is free of biaxial stress. Spatial maps of the Raman shift of the A1(LO) phonon are depicted in Figure 5c,d (Ga-polar and N-polar, respectively). It should be mentioned that there is a considerable difference in the A1(LO) average phonon frequencies measured on the Ga-polar and the N-polar faces. This difference is consistent with a higher concentration of free electrons on the Ga-polar face compared to the N-polar face. Parts c and d of Figure 5 were acquired at the same position on opposite faces of the same sample and show the same lateral gradient in free carrier concentration. Parts c and d of Figure 5 also show the existence

C46v) predicted by group theory have been observed by Raman scattering (RS) spectroscopy.10 It has been demonstrated that free electrons in GaN strongly couple with the A1(LO) phonon, affecting both its phonon frequency and FWHM.11 The phonon frequency of the E2 2 phonon has been demonstrated to be sensitive to strain12 and insensitive to changes in free carrier concentration.13 We used room temperature RS to obtain insight into residual stress, the freeelectron concentration, and their areal and volumetric variations across our samples. Figure 4 shows Raman spectra corresponding to the Z(Y,Y)-Z backscattering geometry, measured along the directions (0001)-GaN (front side or Gapolar face) and (0001̅)-GaN (back side or N-polar face) of sample A. It should be noted that there is no change in the E22 phonon frequency, line shape, and line width. However, from the N-face to the Ga-face significant increases in the A1(LO) phonon frequency and FWHM were observed. These observations, also verified in sample B, are consistent with a biaxial stress-free sample and a detectible higher free carrier C




Article

Figure 5. (a and b) Maps of the Raman shift of the stress-sensitive E22 phonon: (a) Ga-polar face; (b) N-polar face. (c and d) Maps of the Raman shift of the free carrier sensitive A1(LO) phonon: (c) Ga-polar face; (d) N-polar face. A detailed analysis of these 200 μm × 200 μm spatial scans is presented in the text.

with annihilation of excitons bound to neutral donors (DoX), leaving the donors in the ground and excited states, and their phonon replicas (nLO−DoX), near 3.47 eV. Also commonly observed in the spectra are weak zero-phonon lines from shallow-donor/shallow-acceptor pair (DAP) recombination at 3.25 eV and their phonon replicas (nLO−DAP) at lower energies, and the pervasive yellow band around 2.25 eV.10,20 It has been demonstrated that detailed high-resolution, lowtemperature photoluminescence (LTPL) studies of high-quality FS heteroepitaxial HVPE-GaN can be conveniently used to identify the chemical nature of the shallow donors.20,21 With that in mind, high-resolution LTPL at the Ga-polar and N-polar faces of the examined homoepitaxial HVPE samples were carried out. Figure 6 shows the PL spectra of both polar faces of sample A. Two well resolved lines were observed, SioXA1 and OoXA1, respectively associated with the neutral-donor-bound excitonic recombination process that leaves Si and O shallow impurities in the ground state, after the exciton annihilation.22 In addition, it can be observed that the emission line associated with the Si donor, the shallower impurity, is more intense than that associated with the O donor, at the lower energy, indicating that the concentration of the neutral Si impurity is larger than that of the neutral O impurity. Considering that shallow impurities are preferentially compensated for by deep levels, one should expect a larger concentration of shallower Si

of a vertical (and larger) gradient in free carrier concentration (consistent with SIMS depth profiling discussed later). Bulk GaN crystals grown by the ammonothermal method are typically n-type and degenerate, with temperature-independent carrier concentration and mobility typically between 1018 and 1019 cm−3 and between 130 and 60 cm2/(V s), respectively.14 This result is consistent with the strong LO-phonon/ conduction electron−plasmon coupling mode observed in the Raman spectra.2,11 Oxygen is assumed to be the shallow donor responsible for the observed high concentration of free electrons.15 The low-temperature PL spectra of such samples typically show an intense broad band with peak around 2.25 eV and a weak and relatively broad near band edge (NBE) emission near 3.5 eV.16 The large concentration of free carriers causes the NBE emission shift to higher energy (the so-called Burstein−Moss effect).17,18 Unintentionally doped (UID) thick free-standing (FS) heteroepitaxial HVPE-GaN films (grown on sapphire) are also typically n-type, but with much lower background free-electron concentration (≤3 × 1017/cm3) than ammono bulk GaN. Heteroepitaxial HVPE-GaN samples are characterized by dislocation densities of ≤1 × 107/cm2, room temperature carrier mobility of ∼1300 cm2/(V s), and free carrier concentrations ∼ 7 × 1015/cm3.19 Low-temperature PL spectra of such UID FS heteroepitaxial HVPE-GaN films typically show intense recombination emission lines associated D




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Article

SUMMARY We have shown that a 2.3 mm thick, crack-free homoepitaxial HVPE-GaN can be grown on ammonothermal single-crystal GaN substrates. It was also demonstrated that a homoepitaxial HVPE-GaN crystal, previously grown on an Am-GaN singlecrystal substrate, can be used as a seed for subsequent crystal growth. These second-grown homoepitaxial HVPE-GaN wafers reproduce the high crystalline quality and flatness of the original Am-GaN substrate, maintaining the very small values of FWHM of the XRD rocking curve, which implies a very low dislocation density. In addition, a multiple order of magnitude reduction of the concentration of the pervasive Si and O donors, in comparison with that of the Am-GaN, was verified. These new homoepitaxial HVPE substrates exceed the quality of any commercially available GaN substrates. Additional work is in progress to reduce concentrations of these impurities to a lower level which would allow obtaining substrates suitable for the fabrication of devices such as fast high-voltage switches.

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Figure 6. High-resolution, low-temperature PL spectra of the Ga-polar and N-polar faces of sample “A”. The broadening of the SioXA1 emission line observed in the spectrum of the Ga-polar face is associated with the increasing of the Si concentration.

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

donors relative to the deeper O donors than one could estimate by the relative intensity ratio of the related lines. The line width of the SioXA1 was much larger on the Ga-polar face than on the N-polar one. This indicates that the Ga-polar face has a larger concentration of the neutral Si impurity than the N-polar face, which is consistent with the present RS results.23 Despite the high sensitivity of LTPL, this technique is not adequate for impurity concentration measurements due to the difficulties in accounting for all radiative and nonradiative recombination paths and the partial compensation of the defects involved in the recombination processes. To address this uncertainty, both polar faces of sample A were probed by SIMS (secondary ion mass spectroscopy) depth analysis. It was verified that the Gapolar face has [Si] = 2 × 1017/cm−3 and [O] = 1 × 1016/cm−3, while the N-polar face has [Si] = 2 × 1016/cm−3 and [O] = 1 × 1016/cm−3. Carbon, a compensating acceptor impurity in GaN, had a concentration of [C] = 1 × 1016/cm−3, the SIMS detection limit. Considering that the concentration of O is the same in both polar faces, it can be stated that Si is the impurity responsible for the higher concentration of free carriers observed in the Ga-polar face of the samples. It is assumed that the presence of Si in the HVPE system comes from quartz elements (nozzles, susceptor, and tubes) of the reactor. They are etched during growth by the reactive gases at a relatively high temperature. Thus, in time, the concentration of Si close to the growth zone increases. This explains the gradient of the Si concentration in the samples. The N-face of sample A was grown 3.5 h earlier than the Gaface. Herein, it should be remarked that crystallization with a higher growth rate yields HVPE-GaN crystals of a higher purity.8 Growing GaN at the rate of 240 μm/h results in the incorporation of less Si (3 × 1016 cm−3 at the Ga-face). Thus, it seems that growing GaN with a very high growth rate yields high-purity material with donor impurity concentration lower than the SIMS detection limit. The Raman scattering measurements indicated that these samples have a relatively low and uniform free carrier concentration. Electrical properties of similar crystals have free carrier concentrations from 4 × 1016 to 8 × 1016 electrons/cm3 and mobilities of the order of 900− 1000 V s/cm2.24

The authors declare no competing financial interest.

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REFERENCES

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Growth of High Crystalline Quality HVPE-GaN Crystals with Controlled

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