MOCVD growth of GaN films on Si-rich SiNx

MOCVD growth of GaN films on Si-rich SiNx nanoislands patterned sapphire Zhilai Fang*, Shuping Li, Junyong Kang Semiconductor Photonics Research Center and Department of Physics, Xiamen University, Xiamen, China 361005 ABSTRACT We intentionally patterned Si-rich SiNx nanoislands on sapphire substrates and found the SiNx significantly influenced the subsequent growth of GaN films. Distinct GaN islands of triangular base were formed caused by the enhanced diffusion and regrowth anisotropy during the annealing processes of GaN nucleation layers. Subsequent growth of GaN epilayers at high temperature with initial low V/III ratios on the nucleated triangular islands resulted in island coarsening and shape variations from triangular to hexagonal due to the dominating gas phase transport growth mechanism and limited diffusion length. Further growth with high V/III ratios eventually resulted in layer-growth with surface roughness of ~2.6 Å. Both AFM and XRD results showed a significant improvement of the crystalline qualities with estimated threading dislocation (TD) density of about 1×108 cm-2 when Si-rich SiNx nanoislands patterning was performed. Photoluminescence measurements showed that the yellow and blue emissions were substantially suppressed. Keywords: Si-rich SiNx treatment; GaN; MOCVD; Surface morphology; Photoluminescence

1. INTRODUCTION The III-nitrides heteroepitaxy usually induces high density of threading dislocations (TDs) typically on the order of 1091011 cm-2, which have deleterious effects on the optical properties, reliability, and lifetime of the fabricated optoelectronic devices. Several methods, e.g. low-temperature (LT) buffers [1,2], selective area growth (SAG) [3], epitaxial lateral overgrowth (ELO) [4], have been employed to reduce the dislocation density. While conventional LT buffer layers are relatively technically simple and have been widely and intensively studied, it is still difficult to grow GaN films of very low dislocation density (e.g. 107 cm-2). In comparison, traditional ELO can reduce dislocation density to mid 106 cm-2 but requiring complicated ex situ lithography process and thus is time-consuming and expensive. Recently, in-situ SiH4 treatment prior to the growth of LT GaN buffer has been used to improve the crystalline qualities of GaN films and has been proposed to be a promising in-situ nanoscale ELO method with advantages of no ex situ processing and fast coalescence [5,6]. In this study, we performed epitaxial growth of GaN films on Si-rich SiNx nanoislands patterned c-sapphire substrates and investigated the GaN epilayers at different growth stages.

2. EXPERIMENTAL The sapphire substrates were prepared by thermal cleaning at 1060 oC followed by nitridation at 545 oC in a MOCVD system. Before the GaN growth an incomplete SiNx layer was in situ predeposited.. On the SiNx coated sapphire a conventional LT GaN nucleation layer (NL) was grown at 535 oC and 500 Torr followed by a high temperature (HT) annealing process. The subsequent growth of HT GaN epilayers was carried out at 1035 oC and 100 Torr. To study the growth processes of GaN on the Si-rich SiNx patterned sapphire substrates, different growth stages of GaN films were prepared by varying the growth time and other growth conditions. The surface morphologies of SiNx layers, LT GaN NLs, and HT GaN films at different growth stages were investigated by atomic force microscope (AFM, PicoSPM) and scanning electron microscope (SEM, LEO1530) equipped with an energy dispersive X-ray spectrometer (EDX). The surface chemical compositions were analyzed by X-ray photoelectron spectroscopy (XPS, PHI Quantum2000) and also SEM-EDX. The crystal structure was characterized by X-ray diffraction (XRD, Bede QC200). The photoluminescence (PL) excited by a 325 nm He-Cd laser was measured at room temperature (RT) for GaN films prepared on sapphire substrates with and without SiNx treatment. *[email protected]; phone 86 592 2184220; fax 86 592 2184220 Sixth International Conference on Thin Film Physics and Applications, edited by Wenzhong Shen, Junhao Chu, Proceedings of SPIE Vol. 6984, 69842V, (2008) · 0277-786X/08/$18 · doi: 10.1117/12.792371

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3. RESULTS AND DISCUSSION The in situ patterned SiNx layers showed islandlike surface structure with average island size, height, and density of ~100 nm, ~2 nm, and ~1.6×109 cm-2, respectively [7]. XPS studies showed that the SiNx layers are highly Si-rich with a Si/N ratio of ~3.21. SEM-EDX measurements showed that the sapphire substrate surface was not fully covered by the SiNx nanoislands, which made it possible for the GaN epitaxy on the exposed sapphire surface as the dielectric SiNx masks would block the GaN growth. In table 1 below we list the typical growth stages we have investigated. The surface morphologies of GaN at different growth stages (from Stage “A” to “F”) are shown in Figures 1a, 1b, 2a, 2b, 2c, and 2d, respectively. As shown in Figure 1a, before annealing the surface of LT GaN NLs is smooth with a roughness of ~2 nm. Taking into account the deposition thickness of GaN NLs (~25 nm) and the SiNx layers (~2 nm), we conclude that before annealing the patterned Table 1. Description of different growth stages Growth Stages

A B C D E F

Description of the GaN Growth Stages

LT GaN NLs on Si-rich SiNx patterned sapphire without annealing LT GaN NLs on Si-rich SiNx patterned sapphire after annealing Subsequent HT growth with low V/III ratios for 100s on Stage “B” Subsequent HT growth with low V/III ratios for 400s on Stage “B” Further growth at HT with high V/III ratios for ~1000 s on Stage “D” Further growth at HT with high V/III ratios for ~4000 s on Stage “D”

(a)

1 µm

(b)

500 nm

Fig. 1. The AFM images of the surface morphologies of the GaN nucleation layers grown on Si-rich SiNx nanoislands patterned sapphire surface (a) before annealing and (b) after annealing at 1035oC.

SiNx islands were fully covered by the LT GaN NLs. After HT annealing of the LT GaN NLs (Stage “B”) isolated islands were formed which resulted in an increase of the surface roughness. The GaN decomposition was enhanced at HT and hence fractional bare sapphire surface was exposed. Without Si-rich SiNx pretreatment the annealed GaN NLs typically form nucleated islands of small size (~100 nm) and lack of distinct island shapes. Interestingly in our studies with the Si-rich SiNx pretreatment, the annealed LT GaN NLs recrystallized and formed distinct 3D pyramids of triangular base as shown in Figure 1b. The triangular pyramids are of large average lateral size of ~220 nm, height of ~25 nm, and density of ~1.4×109 cm-2. The GaN island density is close to that of the predeposited SiNx islands (~1.6×109 cm-2); from SEM-EDX measurements we found that Si was detected at the GaN island sites whereas not detectable at the other sites; therefore we infer that GaN preferentially nucleates on the bare sapphire surface surrounding the SiNx islands. The excess Si atoms are expected to out diffuse and incorporate into GaN at the subsurface Ga substitutional sites and thus more Ga-rich surface could be formed. As a result, the adatom diffusion barrier would be reduced [7-9]. Furthermore, Si-doping would play roles in the compressive strain relaxation for GaN on sapphire [10]. Consequently, the generation of misfit dislocations would be suppressed by the Si-doping induced strain relaxation and enhanced diffusion. According to the wurtzite GaN crystal structure, there exists different bonding geometry for the edge atoms of the hexagonal units and thus different growth and diffusion anisotropy. The diffusion and reincorporation anisotropy of adatoms (during the annealing process) is generally not very effective (due to the limited diffusion length) and thus generally indistinct island shapes were observed for the conventional annealed GaN NLs. In our growth studies the Sirich SiNx patterning induced Si-Ga exchange process (Si incorporation into GaN islands at the Ga substitutional site) would increase the surface Ga coverage and thus reduce the Ga adatom diffusion barrier. Furthermore, at HT the GaN decomposition and reincorporation of Ga adatoms would also be enhanced. As a result, the diffusion and regrowth

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anisotropy were enhanced and triangular islands of distinct faceting were formed. Therefore the kinetic mechanism for the triangular GaN island shaping can be attributed to enhanced diffusion and regrowth anisotropy [7,9,11].

(a)

1 µm

(b)

1µm

(d)

200 nm

(c)

20 µm

Fig. 2. The surface morphologies of the GaN layers (a) with subsequent HT growth for 100 s (Stage “C”), (b) for 400 s (Stage “D”) with initial low V/III ratios, (c) subsequently grown on Stage “D” at HT with high V/III ratios for ~1000 s (i.e. for Stage “E”), and (d) further growth on Stage “D” at HT with high V/III ratios for ~4000 s (i.e. for Stage “F”).

Subsequent HT GaN growth for 100 s (e.g. Stage “C”) with low V/III ratios resulted in the further increase of surface roughness due to the coarsening of the isolated islands (with increases of lateral size to 200 - 400 nm and decrease of island density from ~1.4×109 cm-2 to ~8×107 cm-2). The island shape changed from triangular to quasi-hexagonal and hexagonal as observed in Figure 2a. Generally the processes of island shaping and shape variations are governed by the competition between growth rate of island edge facets and diffusion rate along island peripheries. Under HT growth conditions the gas dissociation (and thus growth rate) is very fast. Furthermore, with increase of island size the adatom diffusion becomes less effective due to the limited diffusion length compared with the island side length. Hence the growth rate of island edges by gas phase transport becomes dominating in island shaping and shape variations [7,12]. According to the CCS planetary reactor used in our MOCVD system, in which during the HT growth there is the same mass transport rate for the six side facets of the hexagonal units, the triangular (0001) facets of the coarsening islands were gradually changed and eventually hexagonal (0001) facets emerged. In Figure 2b, i.e. the initial HT growth of GaN with low V/III ratios for 400 s for Stage “D”, we found further increase of island size (~1 µm) and height to several hundred nm and further reduction of island density to ~4×105 cm-2 due to the island coalescence and annihilation of some islands. More distinct hexagonal islands with sidewall facets were formed and some of the isolated islands coalesced to form bigger islands. In addition to the observation of quasi-triangular or quasi-hexagonal islands and hexagonal islands, we also observed the phenomenon of island coalescence, e.g. appearance of initially-coalesced islands with clear coalescence boundary, partly-coalesced islands, and fully-coalesced elongated hexagonal islands. Further growth at HT with high V/III ratios for ~1000 s resulted in a flat surface of very large “rounded islands” with smooth top facet and lateral size of ~10 µm (see Figure 2c for Stage “E”). The formation of round shape has further confirmed the isotropic growth rate of island edges at HT in a CCS reactor. This has become evident that the adatom diffusion became less effective at large length scale and the gas phase transport rates of island edges became dominating in island shaping and shape variations under the HT growth conditions. Further HT growth with high V/III ratios for ~4000 s (Stage “F”) resulted in a surface morphology of layer structures as shown in Figure 2d. The surface is smooth with atomic flatness (~2.6 Å). By counting the etched pits of the AFM images we can roughly estimated the dislocation density of less than 1×108 cm-2. We performed ω-scan of the [002] and [102] reflections for GaN films on Si-rich SiNx coated sapphire by use of the Bede QC200 double crystal high-resolution XRD. The typical FWHMs of the [002] and [102] reflections are 240 and 340 arcsec, respectively. Accordingly we estimate the dislocation density of GaN films (~2.5 µm) prepared with SiNx pretreatment to be about 1×108 cm-2. By


optimization of the growth conditions of the SiNx layers, LT GaN NLs, and HT GaN epilayers, further reduction of the TDs density could be achieved. The room-temperature photoluminescence of the GaN films has also been measured. Obviously, as shown in Figure 3, for GaN films without SiNx treatment (the line curve) the yellow-band (YB) and blue-band (BB) emissions are clearly visible. With Si-rich SiNx treatment, the YB and BB emission intensity are substantially decreased to be invisible. The photoluminescence results have further supported the improvement of crystalline qualities and optical properties by means of in-situ Si-rich SiNx patterning.

Intensity (arb. unit)

with Si-rich SiNx treatment without SiNx treatment

Blue Band Yellow Band

350

400

450

500

550

600

650

700

Wavelength (nm)

Fig. 3. Room-temperature photoluminescence of GaN films grown on sapphire with and without Si-rich SiNx patterning.

4. CONCLUSION The growth behaviour of GaN was significantly influenced by the emergence of the Si-rich SiNx nanoislands. During the annealing process the decomposed GaN from the GaN NLs preferentially nucleated on the exposed sapphire surface and surrounding the SiNx nanoislands. The adatom diffusion was enhanced due to the increase of the surface Ga coverage by Si-Ga exchange processes and the Ga-rich growth conditions. Further, Si-doping may partially relieve the compressive strain of GaN on sapphire. Both strain relaxation and enhanced diffusion modified the surface morphology and crystalline qualities of GaN films. The AFM and XRD results indicated a significant improvement of the crystalline qualities for GaN films on Si-rich SiNx patterned sapphire surface, i.e. the dislocation density reduced to ~1×108 cm-2, which is lower than that of GaN films prepared on LT GaN buffers without Si-rich SiNx treatment (typically higher than 109 cm-2). The photoluminescence measurements also showed an enhancement of the optical properties for GaN films on Si-rich SiNx patterned sapphire substrates.

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