DOI: 10.1021/cg901546t
Growth and Characteristics of Zinc-Blende and Wurtzite GaN Junctioned Branch Nanostructures
2010, Vol. 10 2581–2584
Sammook Kang,† Bong Kyun Kang,† Sang-Woo Kim,†,‡,§ and Dae Ho Yoon*,†,§ †
School of Advanced Materials Science and Engineering, Sungkyunkwan University, 300 Cheoncheon-dong, Jangan-gu, Suwon, Gyeonggi-do 440-746, Republic of Korea, Center for Human Interface Nanotechnology (HINT), Sungkyunkwan University, 300 Cheoncheon-dong, Jangan-gu, Suwon, Gyeonggi-do 440-746, Republic of Korea, and § SKKU Advanced Institute of Nanotechnology (SAINT), Sungkyunkwan University, 300 Cheoncheon-dong, Jangan-gu, Suwon, Gyeonggi-do 440-746, Republic of Korea ‡
Received December 9, 2009; Revised Manuscript Received March 23, 2010
ABSTRACT: Both wurtzite and zinc-blende phase junctioned GaN nanostructures were synthesized using thermal chemical vapor deposition methods via the vapor-liquid-solid process for the first time. We observed catalyst movement and the regrowth of nanowire zinc-blende phases. This phenomenon was believed to occur in order to reduce the lattice mismatches between the wurtzite phase GaN and Au planes. The growth route could synthesize the c- and h-GaN junctioned nanostructures. The emission values for the zinc-blende phase GaN nanosturtures in cathodoluminescence were shifted a few meV higher than the reported values because the zinc-blende phase GaN epitaxially grew on wurtzite phase GaN without the residual strain.
1. Introduction
2. Experimental Section
In recent years, III-nitride nanostructures, such as nanoparticles, nanowires, and nanotubes, have attracted extensive attention because of their great potential for use in novel nanoelectronic devices due to their unique electronic and optical properties. Gallium nitride (GaN) has a wide direct bandgap of 3.4 eV at room temperature and is a promising candidate material for short wavelength optoelectronic devices, such as light emitting diodes and laser diodes, as well as high power and high temperature operation devices, on account of its high melting temperature, high breakdown field, and high saturation drift velocity.1,2 In these regards, GaN nanostructures have received a considerable amount of interest because of the great potential for application in optoelectronic and electronic devices.3-6 GaN nanostructure research with various morphologies has focused on the thermodynamically stable hexagonal wurtzite phase (h-GaN). Meanwhile, the metastable zinc-blende cubic phase (c-GaN) has also been investigated, and compared to h-GaN, it has a high mobility, resulting from its lower phonon scattering in a higher crystallographic symmetry, a high p-type conductivity due to the easier doping process, and a high optical gain from its quantum wells structures.7,8 However, the c-GaN nanostructures are difficult to obtain under thermodynamic equilibrium conditions. Only a few reports have synthesized c-GaN nanoparticles and nanotubes or the cubic phase embedded in h-GaN nanowires (NWs).9-12 In this work, branched GaN nanostructures with both the wurtzite and zinc-blende phases were observed for the first time. The c- and h-GaN junctioned nanostructures were synthesized using the thermal chemical vapor deposition (CVD) methods via the vapor-liquid-solid (VLS) process. The growth routes were studied along with the structural and optical properties. These structures could potentially be used in the fabrication of nanoscale functional devices.
The branched GaN nanostructures were synthesized using the thermal CVD method. The starting materials were a mixture of GaN powder (99.999%, High Purity Chem.) and molten Ga (99.9999%, 9Digit Co. Ltd.) at a weight ratio of 1:1. The c-Al2O3 substrate was cleaned through sonication in acetone, and then Au thin films with a thickness of 1 nm were coated onto the substrate using the thermal evaporation system. These thin films were used as a catalyst for branched GaN nanostructure growth. High purity Ar (99.999%) and NH3 (99.99%) were introduced into the reactor as carrier gas and reaction gas, respectively. The Au-coated Al2O3 substrates were placed on the mixture of GaN powder and molten Ga sources in an alumina boat, positioned in the center of the quartz tube. The synthesis experiments were performed at 950 °C. During the main growth of the branch GaN nanostructure, Ar gas and NH3 gas were introduced into the quartz tube at flow rates of 1000 and 20 sccm, respectively. The reaction chamber was maintained under a vacuum of 200 Torr. The main growth time was varied from 15 to 90 min. After the main growth, the samples were naturally cooled to room temperature. GaN branch nanostructures were sythesized by adjusting the growth time. The shape and morphology of the GaN nanostructures were observed using field emission scanning electron microscopy (FESEM, JSM7500F). Transmission electron microscopy (HR-TEM) was carried out using JEM2100F with an accelerating voltage of 200 kV. The crystallinity and structure of the GaN nanostructures were investigated using synchrotron X-ray diffraction (XRD) measurements carried out with a beamline 3C2 at the Pohang Light Source. The synchrotron X-ray was vertically focused using a mirror, and a double bounce Si (111) monochromator was used to monochromatize the X-ray to a wavelength of 1.239 A˚. The spatial localization of the emission was determined using the monochromatic cathodoluminescence (CL) measurements. CL was performed through FESEM using a GATAN MONO CL3þ system with an accelation voltage of 10 kV.
*To whom correspondence should be addressed. E-mail:
[email protected]. r 2010 American Chemical Society
3. Results and Discussion The GaN nanostructures with both the wurtzite and zincblende phases were synthesized using the CVD methods. Figure 1 shows the FESEM images of the GaN nanostructures as a function of growth time. Figure 1a-d shows the Published on Web 05/06/2010
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Figure 1. FESEM images of the branch GaN nanostructures as a function of growth time. (a) Initial stage: the GaN NWs were grown using the VLS routes. (b and c) Middle stages: the GaN NWs started to grow laterally toward the enclosed side planes, and the Au catalyst moved on the surface of the GaN NRs. The GaN NWs regrew using the VLS routes. (d) End stage: the branch GaN NWs started to grow laterally again. The inset in b is a low magnification TEM image of branched GaN nanostructures.
initial-, two middle-, and the end-stage images of the GaN nanostructures during their growth. In Figure 1a, a large number of randomly oriented GaN NWs with diameters ranging from 30 to 70 nm and lengths of 2 μm were synthesized through the VLS routes. The morphology of the GaN NWs was straight with a triangular cross-section and a smooth surface all over the c-Al2O3 substrates. The GaN NWs grew along the [0110] direction and were enclosed by the (0001), (2112), and (2112) side planes.13 The random orientation of the GaN NWs indicated that the growth occurred in a Ga-rich environment. Previous research conducted by several groups showed that the Ga/N reactant ratio in the vapor phase played a role in determining the morphology and growth direction of the GaN nanostructures.14,15 Then, the GaN NWs started to grow laterally toward the enclosed side planes, and the Au catalyst moved on the surface of the GaN nanorods (NRs) with diameters ranging from 130 to 160 nm and without changing the length. After the Au catalyst movement on the GaN NRs, the GaN NWs regrew through the VLS routes only on the edge of the GaN NRs. The growth routes synthesized the branched GaN nanostructures which consist of trunk GaN NRs and thin branch GaN NWs, as shown in Figure 1b and c. Finally, the branched GaN NWs started to grow laterally again (Figure 1d). The detailed crystallography information of the branched GaN nanostructure was observed using TEM and selected area electron diffraction (SAED) analysis. Figure 2a and c shows the branched GaN nanostructures with wurtzite and zinc-blende phases. Figure 2a and b shows the upper side image and the SAED pattern of the individual branched GaN structures. A single crystalline Au catalyst was observed at the end of c-GaN NWs, and periodically, many stacking faults and microtwins were present in the c-GaN NWs. The SAED pattern taken along the [110] zone axis indicated that the growth was along the [111] direction, which was perpendicular to the (111) plane of the c-GaN NWs. This pattern also revealed that the c-GaN NWs epitaxially grew along the Au (111) plane. Figure 2c shows the bottom image of the individual
Figure 2. (a) Top of the HRTEM image for the branch GaN nanostructures. (b) SAED pattern of the c-GaN NWs. (c) Bottom of the HRTEM image for the branched GaN nanostructures. (d) FFT pattern of the branched GaN nanostructure. Enlarged images of the interface (e) between the Au catalyst and c-GaN and (f) between c-GaN and h-GaN.
branched GaN structures. Heteroepitaxial growth of the cubic phase GaN NWs was observed on the hexagonal phase with stacking faults. The fast Fourier transform (FFT) pattern obtained from the box area in Figure 2c shows that the c-GaN NWs grew epitaxially, as shown in Figure 2d. In addition to the hexagonal diffraction spots, extra cubic diffraction spots were also observed as indicated by the arrows. The 1/3 order spots were interpreted as cubic (220) diffraction spots. The h-GaN (0002) spot and the adjacent extra spots were close to the c-GaN (111) plane. Therefore, the h-GaN NRs (0001) and the c-GaN NWs (111) were oriented parallel to each other. Figure 2e shows the high resolution TEM images at the interface of the Au catalyst and c-GaN NWs. The visible lattice fringes were quite perfect, and the crystallographic orientation was clear and uniform. The interplanar spacings of Au (111) and c-GaN (111) were 0.23 and 0.25 nm, respectively, which agreed well with the reported values for the bulk crystals. The h-GaN NRs (0001), c-GaN NWs (111), and Au (111) epitaxially grew parallel to each other. Figure 2f shows the high resolution TEM images at the interface of the
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Figure 3. Schematic diagram of the growth routes for the branch GaN nanostructures. (a) Initial stage, (b and c) middle stages, and (d) end stage.
c-GaN NWs and the h-GaN NRs. A face centered cubic stacking sequence (ABCABC) was observed in the c-GaN NW region, and a hexagonal stacking sequence (ABAB) was observed in the h-GaN NRs region. Figure 3 shows a schematic diagram of each route that was used in the formation of the branched GaN nanostructure based on the SEM and TEM results. In the early stage, the GaN NWs were synthesized through the VLS routes with a triangular cross-section. The GaN NWs began to grow laterally toward the enclosed side planes without vertical growth, and the Au catalyst moved on the surface of the GaN NRs. Almost all of the catalyst moved only along the h-GaN (0001) plane. This phenomenon was believed to occur in order to reduce the lattice mismatches between the h-GaN (0110) and Au (111) planes and between the h-GaN (0001) and Au catalyst (111) planes, which were 17.28 and 10.1%, respectively. The lattice mismatch between the Au catalyst (111) and the other {212} side planes was 58.88%. Therefore, the difference of lattice mismatch difference was a key point for Au catalyst movement, where the branch GaN NWs grew via the VLS route on the (0001) plane of the GaN NRs. After the Au catalyst movement on the GaN NRs, the c-GaN NWs regrew along the [111] direction via the VLS routes because the NWs with stacking defects or a twinning plane, which were highenergy sites, exposed more nucleation sites at the growth interface of the catalyst and NWs than the NWs without defects, which played an important role during branched NWs growth.16,17 The difference in thermodynamic driving force is responsible for the shape choice in the different growth stages such as nucleation, growth, and structural transition of nanostructures.18 The lattice mismatch difference between the h-GaN (0001) and Au catalyst (111) planes and between Au (111) and c-GaN (111) planes were 10.1 and 6.98%, respectively. The other route may also expect a direct consequence of the Ga-stabilized polar surface and formation of liquid phase metal catalyst. If the Ga source incorporated into the Au catalyst faster than Ga is consumed by the growing GaN NRs, the tip particle can become enriched with Ga and grow in diameter, which provides the opportunity for the Ga-rich liquid to leave the tip particle, migrate along the surface of the growing GaN NRs, and ultimately nucleate branches.19 Finally, the branch c-GaN NWs started to grow laterally, and cubic phase embedded in h-GaN nanowires transitioned to the hexagonal phase.
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Figure 4. Synchrotron X-ray scattering measurement result for the branch GaN nanostructures.
Figure 5. (a) FESEM and (b) corresponding CL images for the branched GaN nanostructures. (c) The CL spectrum was taken at 77 K.
Figure 4a shows the synchrotron XRD pattern of the branched GaN nanostructure on c-Al2O3. All of the diffraction patterns were indexed as either the hexagonal structure (JCPDS card number 50-0792) or the cubic GaN (JCPDS card number 80-0012). The diffraction peak for the c-GaN (111) plane was observed, and other peaks such as the (100), (002), and (101) planes in the pattern corresponded to h-GaN. The XRD pattern clearly confirmed the presence of both h- and c-GaN phases, which was in good agreement with the TEM results. No other oxide phases were observed. Figure 5a, b, and c show an FESEM image, the corresponding CL image, and the CL spectrum of the branched GaN nanostructures taken at 77 K, respectively. The CL spectrum clarified the difference between the optical properties of h- and
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c-GaN. Five resolved peaks were observed, two for h-GaN and three for c-GaN. In the spectrum, the two h-GaN peak positions corresponded to the emissions at 3.469 and 3.398 eV, which were attributed to the neutral donor-bound exciton (DoX) emission and its first longitudinal-optical phonon replica (DoX-1LO) emission, respectively. Three peaks in the spectrum were located at 3.278, 3.198, and 3.10 eV emissions. We assigned the 3.278 and 3.198 eV emissions as the excitonic transition (DoX) and the donor-acceptor (DA) pair transition, respectively. The 3.10 eV emission was assigned to the longitudinal-optical phonon replica of the DA pair transition (DAP-1LO) because the energy spacing between the emission and the DA pair transition amounts to about 90 meV.20 The excitonic emission had the strongest intensity in the spectrum. The CL peaks of c-GaN were much broader than the corresponding h-GaN peaks because of the presence of defects originating from stacking faults and microtwins. However, the peak positions for c-GaN fit well with a previous work regarding the bulk c-GaN emission properties.21 4. Conclusions The c- and h-GaN junctioned branch nanostructures were grown via the VLS routes using the thermal CVD method. The catalyst movement and the regrowth of NWs on the sites of the c-GaN phase were observed. The growth route synthesized the c- and h-GaN junctioned nanostructures. The h-GaN NRs (0001) and c-GaN NWs (111) epitaxially grew parallel to each other. The emission values for the c-GaN in CL were shifted a few meV higher than the reported values because the c-GaN crystal epitaxially grew on h-GaN without the residual strain. The branched GaN nanostructures may have significant fundamental and technological implications for the fabrication of nanoscale functional devices. Acknowledgment. We thank Chul-Ho Jung for performing the synchrotron XRD studies.
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