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aPhys. Status Solidi C 8, No. 7–8, 2123–2126 (2011) / DOI 10.1002/pssc.201001026

Free-standing GaN substrates fabricated by a combination of substrate fracturing and chemical lift-off Hiroki Goto, Hyun-Jae Lee*, Seogwoo Lee, Meoung-Whan Cho, and Takafumi Yao*

Center for Interdisciplinary Research, Tohoku University, Aramaki-Aoba, Sendai 980-8578, Japan

Received 17 September 2010, revised 18 October 2010, accepted 21 October 2010 Published online 17 May 2011

Keywords free-standing GaN, chemical lift-off, HVPE, GaN substrate * Corresponding authors: e-mail kalos@cir.tohoku.ac.jp, Phone: +81-90-795-4404, Fax: +81-90-795-7810; * e-mail tyao@cir.tohoku.ac.jp, Phone: +81-90-795-4400, Fax: +81-90-795-7810

Free-standing GaN thick films are successfully fabricated by employing intentional mechanical fracturing of Al2O3 substrates during HVPE growth followed by chemical lift-off of Al2O3 substrates after growth. To do so, ther-mal stress induced in the Al2O3 substrates and overgrown GaN layers are controlled by adjusting the thickness of

the Al2O3 substrates. The grown layers however consist of multi domains, which deteriorate crystal quality. Do-mains coalesce to form larger domains as the growth pro-ceed, which results in an improvement in crystal quality. Bending of free-standing GaN thick layers is of concave shape and its typical bending radius is 2.5 m.

© 2011 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

1 Introduction So far, commercially available GaN substrates have been grown solely by hydride vapor phase epitaxy (HVPE). The first step toward the fabrication of GaN substrates is the preparation of free-standing GaN thick layers. There are several methods to prepare free-standing GaN thick layers by HVPE. These include laser lift-off technique [1], in which UV laser is irradiated to the interface between GaN and sapphire substrate to detach the sapphire substrate. It is also possible to remove sapphire substrate by polishing after the growth of GaN thick layer [2]. GaN thick layers can be chemically lift-off by dissol-ving GaAs substrates, on which GaN is grown [3]. Very recently, thermal stress induced by the difference of ther-mal expansion coefficients of GaN and sapphire substrate is utilized to detach sapphire substrate [4], where GaN thick layers are grown on TiN deposited GaN/sapphire templates and GaN thick layers are detached at the GaN/TiN interface due to thermal stress, thus enabling in-situ lift-off process of GaN thick layers. However, those processes contain expensive additional processes such as preparation of GaN/sapphire templates and Ti deposition, which might have hampered lowering the price of GaN substrates.

In this study, we propose a new method to detach sapphire substrate by combining substrate fracturing and

chemical lift-off processes, which we believe to be a more cost-friendly method. If we use thin sapphire substrate (say 150 μm thick) in stead of conventional 450 μm-thick sapphire substrate, the sapphire substrate will be fractured during cooling down after growth by tensile thermal stress. Regarding the buffer layer which can be chemically etched, we have already demonstrated that CrN can serve as a buf-fer layer for the growth of high-quality GaN layers [5] and that CrN can be etched [6]. Hence the CrN buffer has been applied to chemical-lift-off of thick GaN layers [6] and to the fabrication of vertical LEDs [7].

The purpose of this article is to fabricate free-standing GaN thick layers through the following procedures: (1) GaN growth on CrN buffer predeposited on thin sapphire substrates; (2) intentional fracturing of sapphire substrates by thermal stress during cooling down after GaN growth, (3) which is followed by chemical etching to detach sapphire substrates.

2 Experimental First we need to determine the thickness of sapphire

substrate to induce intentional fracturing. Figure 1 shows the calculated in-plane tensile stress induced in the sapphi-re substrate at the GaN/sapphire interface during cooling

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down after the growth of 500 μm thick GaN layers. Here the in-plane stress is calculated based on a bimetallic mo-del [8]. The parameters used in the calculation are: Growth temperature was 1040 oC; thermal expansion coefficients of GaN and sapphire are 5.45x10-6 /K and 7.50x10-6 /K, respectively; Young modulus of GaN and sapphire are 1.96x1011 Pa and 1.85x1011 Pa, respectively; Poisson ratios of both GaN and sapphire are 0.3. Sapphire substrate will be fractured, if the in-plane tensile stress exceeds the yield strength of sapphire (355 MPa). The critical thickness for fracturing of sapphire substrate is estimated to be 160 μm. Hence, we used 150 μm-thick sapphire substrates in this study. It should be mentioned that sapphire substrates thin-ner than 100 μm cannot be recommended because of hand-ling problems.

Figure 1 Calculated in-plane thermal stress in sapphire substrate based on a bimetallic model [8] as a function of sapphire thickness. The thickness of GaN is 500 μm and the growth tem-perature is 1050 oC.

Details of the growth procedures of GaN on

CrN/sapphire are described elsewhere [6]. Briefly, a metal-lic Cr layer pre-deposited on sapphire substrates was nitri-ded under NH3 ambient in our HVPE system to form a CrN layer. Deposition of low-temperature buffer layer of GaN was followed by high-temperature growth of GaN thick layer. After cooling down, fractures of thin sapphire substrates were removed by chemical lift-off using Ce(NH4)2(NO3)6 as an etchant to obtain thick GaN layers.

In order to evaluate crystallinity of free-standing GaN thick layers, X-ray diffraction (XRD), cathodolum-inescence (CL), and low-temperature photoluminescence (PL) measurements were performed. Scanning electron microscopy (SEM), optical microscopy, and atomic force microscopy (AFM) were used to study the surface and growth morphology of the crystal.

3 Results and discussion Figure 2 shows (a) as-grown GaN/sapphire sample before detached, where the thickness of the GaN layer is 500 μm, (b) the back surface of an as-grown GaN/sapphire sample, and (c) front view of a self-standing GaN layer. Sapphire

substrates were already cracked and parts of the sapphire substrates were already detached. Some of those cracks ex-tended into the back surface of the GaN. Such cracking of GaN layers presumably occurred not only during cooking down, where the sapphire substrate was fractures, but du-ring growth in particular through coalescence of domains, where tensile stress exerted on the GaN layer [9]. Such cracking could lead to the domain structures as observed by X-ray diffraction and cathodoluminessence (CL) mea-surements. It is likely that such cracking of GaN layers could be minimized by optimizing growth procedures, such as adjusting the onset of lateral growth, but obviously need more investigation.

Figure 2 (a) As-grown GaN/sapphire sample before being de-tached, where the thickness of the GaN layer is 500 μm. (b) Back surface of as-grown GaN/sapphire sample. (c) Front view of a self-standing GaN layer.

Figure 3 shows the (a) (0002) and (b) (10-10) X-ray ω-

rocking curves of the free-standing 500 μm-thick GaN lay-er at the front surface. The (0002) X-ray ω-rocking curve, which reflects the density of screw dislocations, shows a multi-peak behavior with the apparent full width at half maximum (FWHM) value of 1520 arcsec presumably due to the existence of misoriented domains. However, the diffraction peak consists of very sharp peaks with typical FWHM value of several tens of arcsec, which indicates that the constituent domains have high crystal quality. The (10-10) X-ray ω-rocking curve, which reflects the density of edge dislocations, shows a two sharp-peak behavior with a FWHM value of 202.7 arcsec, indicating again the existence of domains with two major in-plane rotation angles.

The (0002) X-ray ω-rocking curve at the back surface (Fig. 3(c)) shows an asymmetric line shape with an appar-ent FWHM value of 992 arcsec. This FWHM value is even narrower than that of the front surface, indicating that the distribution of tilted domains is even narrower at the growth seed, where the low-temperature buffer layer was deposited as a seed layer, rather than at the growth front, where high-quality layer is grown at high temperature. Such strange difference in FWHM value suggests that the misoriented domains may have developed as the growth proceeded and can be diminished by optimizing growth conditions. The Since it was very difficult to detect the

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(10-10) diffraction at the back surface, we show the (30-32) X-ray ω-rocking curve in Fig. 3(d) to get insight into in-plane rotation of domains at the growth seed. The diffraction curve shows again asymmetric line shape with a FWHM value of 586 arcsec, indicating the existence of rotational domains as well.

Figure 3 (a) (0002) and (b) (10-10) omega rocking curves of the front surface of a free-standing GaN thick layer. (c) (0002) and (d) (30-32) omega rocking curves of the back surface of GaN thick layer.

Bending of as-grown free-standing 500 μm-thick GaN

layer was of concave shape and its typical bending radius measured 2.5 m with stylus measurement system. We note that this value is even larger than reported values of as-grown free-standing GaN layer obtained by laser lift-off (1.21 m) [10] and by self-separation (0.43 m) [10].

Figure 4 shows room-temperature CL spectra of free-standing GaN thick layers with various thicknesses. The CL spectra taken at an acceleration voltage of 25 kV with a sample current of 70 μA show dominant near band edge emission peak at around 3.5 eV, which can be attributed to band-edge emission. The next strong emission is observed at around 2.3 eV, which is associated with Ga vacancy. We note that the emission intensity increases with GaN layer thickness. Such dependence can be tentatively explained in terms reduction of non-radiative process as the growth proceeds. If the radiative life time is denoted as τr, and the non-radiative lifetime is deonoted as τnr, then the probabil-ity of radiative recombination is given by η = τr

-1 / { τr-1 +

τnr-1} [11], which indicates that as the nonradiative recom-

bination rate τnr-1 is decreased, the radiative recombination

probability is enhanced. In a simple one-trap model, τnr is inversely proportional to the density of trap. Therefore, the radiative recombination probability is enhanced as the trap density is decreased, even if τr is unchanged. It is most li-kely than the radiative recombination rate is enhanced, while the nonradiative recomnination rate is decreased, as the crystal quality is improved. Thus, the observed increase

in CL intensity with the layer thickness suggests an impro-vement in crystal quality as the growth proceeds.

Figure 4 Room-temperature CL spectra of GaN thick layers with various thicknesses: 500 μm, 200 μm, and 40 μm from the high-est intensity curve to the lowest intensity one.

Figure 5 shows SEM images of the front surface of GaN layers with a thickness of (a) 40 μm, (b) 200 μm, and (c) 500 μm and mapping images of the room temperature CL emission at 364 nm (Band-edge emission) from the GaN layers with a thickness of (d) 40 μm, (e) 200 μm, and (f) 500 μm. When the GaN thickness is thin (40 μm), sur-face defects (dark spots) with a typical radius of 1μm are observed on the surface (Fig. 5(a)). As the thickness in-creases the number of those surface defects is reduced, and the 500 μm-thick GaN layer shows only few surface de-fects (Fig. 5(c)). Since the emission intensity from the 40 μm-thick sample is very weak, we could not obtain clear mapping image for this sample (Fig. 5(d)). As the thick-ness increases, the CL intensity increases and the CL map-ping show domain structures as shown in Figs. 5(e) and (f), where the darks lines are considered to indicate domain boundaries, since domain boundaries are expected to act as the source of non-radiative centers [11]. It is likely that as the growth proceeds, domains coalesce to form larger do-mains thus reducing domain boundaries as observed in the change in CL mapping from Fig. (e) to Fig. (f). It should be mentioned that the observed behavior in CL mapping is consistent with the discussion to explain the change of CL intensity with the GaN layer thickness.

GaN surfaces were etched by a solution of H2SO4 +H3PO4 at 200 oC to estimate dislocation density. Etch pits were predominantly observed at around the domain boun-daries. The etch-pit density was measured to be 5.5 x 108/cm2. Although this value can be compared with the dark spot density (DSD) in CL images, it was difficult to measure DSD due to insufficient resolution.

Hall measurements were performed to evaluate the carrier concentration and carrier mobility. The 500 μm-thick GaN was n-type with an electron concentration of 1.0 x 1017/cm3 and an electron mobility of 50 cm2/V sec. The mobility value is considerably smaller compared to that of

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conventional free-standing GaN (200-300 cm2/V sec), pre-sumably due to the presence of domains in the GaN layers.

Figure 5 SEM images of the front surface of GaN layers with a thickness of (a) 40 μm, (b) 200 μm, and (c) 500 μm. Mapping images of the room temperature CL emission at 364 nm (Band-edge emission) from the GaN layers with a thickness of (e) 40 μm, (f) 200 μm, and (g) 500 μm.

4 Summary

Free-standing GaN layers with a bending radius of 2.5 m are successfully obtained by a combination of thin sapphire substrates and chemical lift-off. XRD studies show the presence of multiple domains which caused mul-ti-peak diffraction behaviors. CL study indicates that the size of domains increases due to coalescence of domains through which crystal quality is improved as the growth proceeds. It is suggested that more optimization of growth procedure is needed to obtain single domain structure.

Acknowledgements The authors are grateful to Furukawa Electromechanical Co. Ltd. for CL measurements.

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