pubs.acs.org/crystalPublished on Web 05/06/2010r 2010 American Chemical Society

DOI: 10.1021/cg901546t

2010, Vol. 102581–2584

Growth andCharacteristics of Zinc-Blende andWurtzite GaNJunctioned

Branch Nanostructures

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 chemicalvapor depositionmethods via the vapor-liquid-solid process for the first time.We observed catalystmovement and the regrowthof nanowire zinc-blende phases. This phenomenon was believed to occur in order to reduce the lattice mismatches between thewurtzite phaseGaNandAuplanes. The growth route could synthesize the c- and h-GaN junctioned nanostructures. The emissionvalues for the zinc-blende phase GaN nanosturtures in cathodoluminescence were shifted a few meV higher than the reportedvalues because the zinc-blende phase GaN epitaxially grew on wurtzite phase GaN without the residual strain.

1. Introduction

In recent years, III-nitride nanostructures, such as nanopar-ticles, nanowires, and nanotubes, have attracted extensiveattention because of their great potential for use in novelnanoelectronic devices due to their unique electronic andoptical properties. Gallium nitride (GaN) has a wide directbandgap of 3.4 eV at room temperature and is a promisingcandidate material for short wavelength optoelectronic devi-ces, such as light emitting diodes and laser diodes, as well ashigh power and high temperature operation devices, on acc-ount of its high melting temperature, high breakdown field,and high saturation drift velocity.1,2 In these regards, GaNnanostructures have receiveda considerable amountof interestbecause of the great potential for application in optoelectronicand electronic devices.3-6

GaNnanostructure researchwith variousmorphologies hasfocused on the thermodynamically stable hexagonal wurtzitephase (h-GaN). Meanwhile, the metastable zinc-blende cubicphase (c-GaN) has also been investigated, and compared toh-GaN, it has a highmobility, resulting from its lower phononscattering in a higher crystallographic symmetry, a high p-typeconductivity due to the easier doping process, and a highoptical gain from its quantum wells structures.7,8 However,the c-GaNnanostructures are difficult to obtain under thermo-dynamic equilibrium conditions. Only a few reports have syn-thesized c-GaN nanoparticles and nanotubes or the cubicphase embedded in h-GaN nanowires (NWs).9-12

In this work, branched GaN nanostructures with both thewurtzite and zinc-blende phases were observed for the firsttime. The c- and h-GaN junctioned nanostructures were syn-thesized using the thermal chemical vapor deposition (CVD)methods via the vapor-liquid-solid (VLS) process. Thegrowth routes were studied along with the structural andoptical properties. These structures could potentially be usedin the fabrication of nanoscale functional devices.

2. Experimental Section

The branched GaN nanostructures were synthesized using thethermal CVDmethod. The startingmaterials were amixture ofGaNpowder (99.999%, High Purity Chem.) and molten Ga (99.9999%,9Digit Co. Ltd.) at a weight ratio of 1:1. The c-Al2O3 substrate wascleaned through sonication in acetone, and then Au thin films with athickness of 1 nm were coated onto the substrate using the thermalevaporation system. These thin films were used as a catalyst forbranchedGaNnanostructure growth.High purityAr (99.999%) andNH3 (99.99%) were introduced into the reactor as carrier gas andreaction gas, respectively. The Au-coated Al2O3 substrates wereplaced on the mixture of GaN powder and molten Ga sources inan alumina boat, positioned in the center of the quartz tube. Thesynthesis experiments were performed at 950 �C. During the maingrowth of the branch GaN nanostructure, Ar gas and NH3 gas wereintroduced into the quartz tube at flow rates of 1000 and 20 sccm,respectively. The reaction chamber was maintained under a vacuumof 200 Torr. The main growth time was varied from 15 to 90 min.After the main growth, the samples were naturally cooled to roomtemperature. GaN branch nanostructures were sythesized by adjust-ing the growth time.

The shape andmorphology of the GaN nanostructures were obser-ved using field emission scanning electron microscopy (FESEM,JSM7500F). Transmission electron microscopy (HR-TEM) was car-ried out using JEM2100Fwith an accelerating voltage of 200 kV. Thecrystallinity and structure of the GaN nanostructures were investi-gated using synchrotron X-ray diffraction (XRD) measurementscarried out with a beamline 3C2 at the Pohang Light Source. Thesynchrotron X-ray was vertically focused using a mirror, and adouble bounce Si (111) monochromator was used to monochroma-tize the X-ray to a wavelength of 1.239 A. The spatial localizationof the emission was determined using the monochromatic cathodo-luminescence (CL) measurements. CL was performed through FES-EM using a GATAN MONO CL3þ system with an accelationvoltage of 10 kV.

3. Results and Discussion

The GaN nanostructures with both the wurtzite and zinc-blende phases were synthesized using the CVD methods.Figure 1 shows the FESEM images of the GaN nanostruc-tures as a function of growth time. Figure 1a-d shows the*Towhomcorrespondence should be addressed. E-mail: dhyoon@skku.edu.

2582 Crystal Growth & Design, Vol. 10, No. 6, 2010 Kang et al.

initial-, two middle-, and the end-stage images of the GaNnanostructures during their growth. In Figure 1a, a largenumber of randomly oriented GaN NWs with diametersranging from30 to70nmand lengths of 2μmwere synthesizedthrough the VLS routes. The morphology of the GaN NWswas straight with a triangular cross-section and a smoothsurface all over the c-Al2O3 substrates. The GaN NWs grewalong the [0110] direction and were enclosed by the (0001),(2112), and (2112) side planes.13 The random orientation oftheGaNNWs indicated that the growthoccurred in aGa-richenvironment. Previous research conducted by several groupsshowed that theGa/N reactant ratio in the vapor phase playeda role in determining themorphology and growth direction ofthe GaN nanostructures.14,15 Then, the GaN NWs started togrow laterally toward the enclosed side planes, and the Aucatalyst moved on the surface of the GaN nanorods (NRs)with diameters ranging from130 to 160 nm andwithout chan-ging the length. After the Au catalyst movement on the GaNNRs, the GaN NWs regrew through the VLS routes only onthe edge of the GaN NRs. The growth routes synthesized thebranched GaN nanostructures which consist of trunk GaNNRs and thin branchGaNNWs, as shown inFigure 1b and c.Finally, the branched GaN NWs started to grow laterallyagain (Figure 1d).

The detailed crystallography information of the branchedGaN nanostructure was observed using TEM and selectedarea electron diffraction (SAED) analysis. Figure 2a and cshows the branched GaN nanostructures with wurtzite andzinc-blende phases. Figure 2a and b shows the upper sideimage and the SAEDpattern of the individual branchedGaNstructures. A single crystalline Au catalyst was observed atthe endof c-GaNNWs, andperiodically,many stacking faultsand microtwins were present in the c-GaN NWs. The SAEDpattern taken along the [110] zone axis indicated that thegrowthwas along the [111] direction,whichwas perpendicularto the (111) plane of the c-GaN NWs. This pattern alsorevealed that the c-GaN NWs epitaxially grew along the Au(111) plane.Figure 2c shows thebottom imageof the individual

branchedGaNstructures.Heteroepitaxial growthof the cubicphase GaN NWs was observed on the hexagonal phase withstacking faults. The fast Fourier transform (FFT) patternobtained from the box area in Figure 2c shows that the c-GaNNWs grew epitaxially, as shown in Figure 2d. In addition tothe hexagonal diffraction spots, extra cubic diffraction spotswere also observed as indicated by the arrows. The 1/3 orderspots were interpreted as cubic (220) diffraction spots. Theh-GaN (0002) spot and the adjacent extra spots were close tothe c-GaN (111) plane. Therefore, the h-GaNNRs (0001) andthe c-GaN NWs (111) were oriented parallel to each other.Figure 2e shows the high resolution TEM images at theinterface of the Au catalyst and c-GaN NWs. The visiblelattice fringes were quite perfect, and the crystallographicorientation was clear and uniform. The interplanar spacingsof Au (111) and c-GaN (111) were 0.23 and 0.25 nm,respectively, which agreed well with the reported values forthe bulk crystals. The h-GaNNRs (0001), c-GaNNWs (111),andAu (111) epitaxially grew parallel to each other. Figure 2fshows the high resolution TEM images at the interface of the

Figure 1. FESEM images of the branch GaN nanostructures as afunction of growth time. (a) Initial stage: the GaNNWswere grownusing the VLS routes. (b and c)Middle stages: theGaNNWs startedto grow laterally toward the enclosed side planes, and the Aucatalyst moved on the surface of the GaN NRs. The GaN NWsregrew using the VLS routes. (d) End stage: the branch GaN NWsstarted to grow laterally again. The inset in b is a low magnificationTEM image of branched GaN nanostructures.

Figure 2. (a) Top of the HRTEM image for the branch GaNnanostructures. (b) SAED pattern of the c-GaN NWs. (c) Bottomof the HRTEM image for the branched GaN nanostructures.(d) FFT pattern of the branched GaN nanostructure. Enlargedimages of the interface (e) between the Au catalyst and c-GaN and(f) between c-GaN and h-GaN.

Article Crystal Growth & Design, Vol. 10, No. 6, 2010 2583

c-GaN NWs and the h-GaN NRs. A face centered cubicstacking sequence (ABCABC) was observed in the c-GaNNW region, and a hexagonal stacking sequence (ABAB) wasobserved in the h-GaN NRs region.

Figure 3 shows a schematic diagram of each route that wasused in the formation of the branched GaN nanostructurebased on the SEM and TEM results. In the early stage, theGaN NWs were synthesized through the VLS routes with atriangular cross-section. The GaN NWs began to grow later-ally 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 toreduce the lattice mismatches between the h-GaN (0110) andAu (111) planes and between the h-GaN (0001) and Aucatalyst (111) planes, which were 17.28 and 10.1%, respec-tively. The latticemismatch between theAu catalyst (111) andthe other {212} side planes was 58.88%. Therefore, the differ-ence of lattice mismatch difference was a key point for Aucatalyst movement, where the branchGaNNWs grew via theVLS route on the (0001) plane of the GaNNRs. After the Aucatalyst movement on theGaNNRs, the c-GaNNWs regrewalong the [111] direction via the VLS routes because the NWswith stacking defects or a twinning plane, which were high-energy sites, exposed more nucleation sites at the growthinterface of the catalyst and NWs than the NWs without def-ects, which played an important role during branched NWsgrowth.16,17 The difference in thermodynamic driving force isresponsible for the shape choice in the different growth stagessuch as nucleation, growth, and structural transition of nano-structures.18 The lattice mismatch difference between theh-GaN (0001) and Au catalyst (111) planes and between Au(111) and c-GaN (111) planes were 10.1 and 6.98%, respec-tively.The other routemayalso expect a direct consequence ofthe Ga-stabilized polar surface and formation of liquid phasemetal catalyst. If the Ga source incorporated into the Aucatalyst faster thanGa is consumedby the growingGaNNRs,the tip particle can become enriched with Ga and grow indiameter, which provides the opportunity for the Ga-richliquid to leave the tip particle, migrate along the surface ofthe growing GaN NRs, and ultimately nucleate branches.19

Finally, the branch c-GaNNWs started to grow laterally, andcubic phase embedded in h-GaN nanowires transitioned tothe hexagonal phase.

Figure 4a shows the synchrotron XRD pattern of thebranched GaN nanostructure on c-Al2O3. All of the diffrac-tion patterns were indexed as either the hexagonal structure(JCPDS card number 50-0792) or the cubic GaN (JCPDScard 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 bothh- and c-GaN phases, which was in good agreement with theTEM results. No other oxide phases were observed.

Figure 5a, b, and c showanFESEM image, the correspond-ing CL image, and the CL spectrum of the branched GaNnanostructures taken at 77 K, respectively. The CL spectrumclarified the difference between the optical properties of h- and

Figure 3. Schematic diagram of the growth routes for the branchGaN nanostructures. (a) Initial stage, (b and c) middle stages, and(d) end stage.

Figure 4. Synchrotron X-ray scattering measurement result for thebranch GaN nanostructures.

Figure 5. (a) FESEM and (b) corresponding CL images for thebranched GaN nanostructures. (c) The CL spectrum was taken at77 K.

2584 Crystal Growth & Design, Vol. 10, No. 6, 2010 Kang et al.

c-GaN. Five resolved peaks were observed, two for h-GaNand three for c-GaN. In the spectrum, the two h-GaN peakpositions corresponded to the emissions at 3.469 and3.398 eV,which were attributed to the neutral donor-bound exciton(DoX) emission and its first longitudinal-optical phononreplica (DoX-1LO) emission, respectively. Three peaks inthe spectrum were located at 3.278, 3.198, and 3.10 eVemissions. We assigned the 3.278 and 3.198 eV emissions asthe excitonic transition (DoX) and the donor-acceptor (DA)pair transition, respectively. The 3.10 eV emission was assig-ned to the longitudinal-optical phonon replica of theDApairtransition (DAP-1LO) because the energy spacing betweenthe emission and the DA pair transition amounts to about90 meV.20 The excitonic emission had the strongest intensityin the spectrum. The CL peaks of c-GaN were much broaderthan the corresponding h-GaN peaks because of the presenceof defects originating from stacking faults and microtwins.However, the peakpositions for c-GaNfitwellwithapreviouswork regarding the bulk c-GaN emission properties.21

4. Conclusions

The c- and h-GaN junctioned branch nanostructures weregrown via the VLS routes using the thermal CVD method.The catalyst movement and the regrowth of NWs on the sitesof the c-GaN phase were observed. The growth route synthe-sized the c- and h-GaN junctioned nanostructures. Theh-GaN NRs (0001) and c-GaN NWs (111) epitaxially grewparallel to each other. The emission values for the c-GaN inCL were shifted a few meV higher than the reported valuesbecause the c-GaN crystal epitaxially grew on h-GaNwithoutthe residual strain. The branched GaN nanostructures mayhave significant fundamental and technological implicationsfor the fabrication of nanoscale functional devices.

Acknowledgment. WethankChul-Ho Jung for performingthe synchrotron XRD studies.

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