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Growth, structural and optical properties of ternary InGaN nanorods prepared by selective-

area metalorganic chemical vapor deposition

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2014 Nanotechnology 25 225602

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Growth, structural and optical properties ofternary InGaN nanorods prepared byselective-area metalorganic chemical vapordeposition

Jie Song, Benjamin Leung, Yu Zhang and Jung Han

Department of Electrical Engineering, Yale University, New Haven, CT 06511, USA

E-mail: jung.han@yale.edu

Received 5 September 2013, revised 4 March 2014Accepted for publication 1 April 2014Published 8 May 2014

AbstractTernary InGaN nanorods were prepared on dielectric-masked nano-holes with selective areametalorganic chemical vapor deposition. To overcome the tendency for random nucleation ofGaN at low temperatures, a pulsed growth procedure was introduced to enhance the diffusionlength of Ga adatoms on SiO2, resulting in good selectivity at typical temperature ranges forInGaN. Photoluminescence from the InGaN nanorods can be tuned from near ultraviolet(400 nm) to blue-green (∼500 nm). Microstructural properties were characterized bytransmission electron microscopy; threading dislocations from the underlying GaN templatewere terminated at the nanorod/template interface, resulting in dislocation-free nanorods. Theheight of dislocation-free InGaN nanorods is about 150 nm, which is much larger than thecritical thickness for the onset of misfit dislocations in planar InGaN growth with typicalthickness of less than 10 nm for an indium composition between 10 and 20%. The compositionprofile of In along the growth direction was examined by energy dispersive x-ray spectroscopicmapping and line scan. Oscillations of In composition along the growth direction were observedand are likely due to the kinetic competition between In and Ga adatoms. These InGaN nanorodsare expected to be useful as templates for growing higher In composition nano-light-emittingdiodes.

Keywords: InGaN, nanorods, MOCVD, pulsed growth

(Some figures may appear in colour only in the online journal)

1. Introduction

The internal quantum efficiency of conventional GaN-basedlight-emitting diodes (LEDs) drops precipitously towardslonger wavelengths as the In composition increases, a phe-nomenon known as the ‘green gap’ [1]. Explanations invokemechanisms such as the immiscibility gap of InGaN [2], thedifficulty of incorporating In under compressive strain(compositional pulling) [3], and large internal polarizationfields [4]. All these mechanisms are related directly orindirectly to the difficulty of planar heteroepitaxy of InGaNon GaN under a high compressive strain, which leads to poormorphology and deteriorating microstructures [2, 5, 6].

However, there have been a few intriguing reports of InGaNnanostructures, in the form of InGaN nanowires [7], InGaNnanorings and nanoarrays [8], and InGaN/GaN core-shellstructure nanorods [9, 10], which suggest that high In-contentInGaN is not inherently defective if the issues of strain andlattice mismatch can be circumvented. Thus we propose togrow InGaN nanorods as nano-scale templates to support thesynthesis of higher In-composition InGaN multiple quantumwells (MQWs) and LEDs. Unlike kinetics-dominated mole-cular beam epitaxy (MBE) [11] or catalyst-mediated growth[12], metalorganic chemical vapor deposition (MOCVD) doesnot support bottom-up growth of one-dimensional nanos-tructures. To create InGaN nanorods, we adopted a strategy in

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nanoscale confined epitaxy [13, 14] where nanoscalecylindrical holes are created through dielectric masks overGaN layers. InGaN is then ‘funneled’ into the nanoscale holesthrough selective epitaxy. Anticipated benefits of usingInGaN nanorods include (1) the elimination of the pre-exist-ing dislocations in the underlying template from aspect-ratioengineering [15], and (2) the possibility of strain relaxationelastically based on geometric effect [16] to minimize thecompositional pulling effect for high In compositional InGaN.Growth of binary GaN nanorods has been demonstrated byseveral groups [13, 17–20]. Core-shell nano-LEDs [10, 21]and axial InGaN MQWs [22] grown on GaN nanorods havealso been reported.

In this paper we report the growth of InGaN nanorods bynanoscale selective-area MOCVD. To overcome the chal-lenge of much reduced surface diffusion for Ga adatomsduring InGaN growth, a pulsed growth technique is demon-strated. InGaN nanorods with In compositions up to 20% areprepared at a thickness much greater than the critical thick-ness in planar epitaxy. Threading dislocations (TDs) propa-gating from the underlying GaN templates are terminated atthe nanorod/underlayer interface, resulting in dislocation-freeInGaN nanorods. The height of dislocation-free InGaNnanorods is about 150 nm, which is much larger than thecritical thickness for the onset of misfit dislocations in planarInGaN where growth is typically less than 10 nm for anindium composition between 10 to 20%. These InGaNnanorods are expected to be useful as templates for growinghigher In compositional nano-LEDs.

2. Experiment details

Nano-patterned substrates were fabricated on MOCVD grownGaN/sapphire templates according to a procedure illustratedin figure 1. First, a SiO2 dielectric mask with a thickness ofabout 0.2 μm was deposited on a 2-μm-thick unintentionallydoped c-plane GaN/sapphire template by plasma-enhancedchemical vapor deposition, as shown in figure 1(a). We notethat the thickness of SiO2 determines the height of thenanorods. Interference lithography (IL) was used to createperiodic arrays of nanoscale holes. As shown in figures 1(b)and (c), a spinning layer of hexamethyldisilazane (HMDS)with a thickness of 180 nm was exposed to a He-Cd layerwith a photo flux of 45 μJ for 6 min [23]. Once the nano-patterns by IL were developed, a Ti/Ni (5 nm/40 nm) Nietching mask was deposited followed by lifting off in acetoneand the residual nano-patterned metal as a mask was left, asshown in figures 1(d) and (e). The nano-pattern was thentransferred into SiO2 by using fluorine-based reactive ionetching (RIE), and finally the metal mask was removed bychlorine-based dry etching, as shown in figure 1(d).Figure 2(a) shows the SEM image of the processed substratewith nano-holes with a diameter of 100 nm and pitch of300 nm. Before loading the sample into the MOCVD reactorchamber, the substrates were cleaned in piranha (H2SO4:H2O2 = 3:1) and 15% HCl solution. The substrates were thenrinsed in deionized water for 5 min.

The growth was carried out by MOCVD (Aixtron 200-4RF/S). Trimethylgallium (TMGa), trimethylindium (TMIn),and ammonia (NH3) were used as Ga, In, and N precursors,respectively. A thin layer of GaN as nano-buffer was firstgrown at 1030 °C in a mixture of N2 (3 slm) and NH3 (3 slm)to ensure that InGaN nanoepitaxy was carried out in a con-sistent way. The GaN nano-buffer showed a flat top surfacewith a thickness of around 0.1 μm as shown in figure 2(b).The temperature was then lowered down to 780–850 °C togrow InGaN. Different In compositions of InGaN nanorodswere obtained by varying the growth temperature and indiumcomposition in the gas phase. Microscopic morphology wasexamined by a Hitachi scanning electron microscope (SEM).Double crystal x-ray diffraction (XRD) was carried out usinga Bede D1 diffractometer. Optical properties were measuredby photoluminescence (PL) using a 325 nm He-Cd laser atroom temperature (RT). Microstructural properties werecharacterized by transmission electron microscopy (TEM)and energy dispersive x-ray spectroscopy (EDS) with a FEIF20 microscope operating at an acceleration voltage of200 kV. TEM specimens were prepared by mechanical pol-ishing first followed by Ar-ion milling.

3. Results and discussion

The possibility of selective growth of InGaN was investigatedrecently by Shioda et al [24]. The issues involved include (1)the competition between the kinetic incorporation of In(Ga)Nand the thermodynamic formation of In droplets, (2) thetendency for In(N) desorption at elevated temperatures, exa-cerbated by the presence of hydrogen (from H2 or NH3), (3)

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Figure 1. Schematic diagrams of the fabrication method for nano-patterned substrates: (a) deposit 229 nm SiO2 on GaN/sapphiretemplate; (b) spin 180 nm HMDS and bake for 90 s at 110 °C; (c)develop nano-pattern in HMDS interference lithography and thendeposit metal mask (5 nm Ti + 40 nm Ni); (d) lift off the metal maskon HMDS in acetone; (e) transfer the nano-pattern into SiO2 by RIE;(f) remove residual metal mask by dry etching.

the random nucleation of GaN crystals on mask due to a highsupersaturation in typical InGaN growth conditions, and (4) alarge disparity in surface diffusion between In and Ga ada-toms. The critical task in the selective growth of InGaN istherefore to enhance the surface diffusion of Ga adatoms onthe SiO2 mask surface. As surface diffusion is intimatelyrelated to the atomistic configuration of a growing surface, aplausible way to control the surface diffusion is by alternatingthe precursor supply intervals in pulsed mode to allow sub-stantial changes of surface stoichiometry and reconstructionsnot typically attainable in continuous growth. Such a pulsedgrowth procedure has been found to be effective in variousmaterial systems [25–28]. During InGaN growth, the galliumprecursor (TEGa) was introduced in a pulsed way while TMInand NH3 flows were kept constant. The interruption durationof TEGa was chosen to be between 3 and 5 s. We found that ifthe interruption duration of TEGa was too short, the selec-tivity couldn’t be improved. If the interruption duration ofTEGa was too long, the In in the InGaN nanorods wouldreevaporate and accumulate on the surface to form In dro-plets. Figure 3(a) shows the morphology of InGaN nanorodsgrown at 780 °C with a gas-phase ratio of 59% under theconventional continuous-growth method. A high density ofpolycrystalline GaN was observed on the SiO2 mask surface;the majority of InGaN growth takes place as randomnucleation on the mask rather than proceeding epitaxiallywithin the nanopatterned holes (marked by arrows as shownin figure 3(a)). We also noted that the PL intensity of thesample corresponding to figure 3(a) was very weak and therewas no detectable InGaN peak in XRD. Based on the dis-tribution of polycrystalline GaN islands on the SiO2 masksurface around the opening holes, the surface diffusion lengthof Ga adatoms on SiO2 is estimated to be less than 50 nm,which is comparable with the result reported by Lee et al [29].The selectivity for InGaN growth over SiO2 was noticeablyimproved with the pulse procedure. Figure 3(b) shows theSEM image of the InGaN nanorods grown with pulsedgrowth while keeping all other parameters the same as thesample of figure 3(a). The selectivity of InGaN over SiO2 is

achieved, implying that the surface diffusion length of Gaadatoms on SiO2 is greater than about 160 nm

× −⎡⎣ ⎤⎦( )2 300 100 2 nm at 780 °C. Due to the contribution

of enhanced surface diffusion of Ga adatoms on SiO2, we alsoobserved that a greater amount of InGaN nanorods grew inthe holes with pulsed growth than in those without pulsedgrowth, as shown in figure 3(b). Figure 3(c) is the top-viewSEM image of the as-grown sample by pulsed growth andInGaN nanorods exhibiting pyramidal top surface bounded bysix {10-11} facets. Figure 3(d) shows a 45°-tilted SEM imageof the InGaN nanorods after removing the SiO2 where theheight of each nanorod is about 250 nm. Longer InGaNnanorods can conceivably be prepared with a thicker SiO2

nano-patterned layer using the same procedure for selectivegrowth.

The as-grown samples were characterized by PL at RT.Three samples A, B and C were grown at temperatures of850, 810 and 780 °C, respectively, with a constant In/(In +Ga) gas phase ratio of 59%. Sample D was grown at 780 °Cwith an increased gas phase ratio of 87%. As shown infigure 4(a), the PL peaks of A–D shift from 406 to 492 nm.We also tried lowering growth temperature further to 750 °Cand obtained higher In compositional InGaN nanorods withgood selectivity. However, PL intensity degraded due topossibly a high concentration of point defects under non-optimal growth conditions. Below 750 °C we started toobserve a loss of selectivity. Based on the PL results, Incompositions in InGaN nanorods were calculated using the

formula = + − − −( ) ( )E x x bx x0.67 3.42 1 1g , where Eg

denotes the band gap energy in eV of InGaN at RT and=b 1.7 is the bowing parameter chosen from the result of

Moses et al [30]. The In compositions of samples A, B, C andD are 8.5%, 11.7%, 17.7% and 20.5%, respectively. Thesample C grown at 780 °C was also characterized by XRD 2θ/ω scan as shown in figure 4(b). A peak located at 34.57° isattributed to the GaN template and the InGaN peak is locatedat ∼34.1°. A theoretical work predicts that, given the thick-ness and diameter of the InGaN nanorods, 90% of the

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Figure 2. (a) SEM image of nano-patterned substrate. (b) 45°-tilted SEM image of GaN nanorod buffer.

mismatched strain should be elastically relaxed [16]. Byassuming that the InGaN nanorods are fully relaxed, the Incomposition in InGaN nanorods is calculated to be 12.3%.The discrepancy between In compositions calculated by XRDresults and PL results is likely due to the In localization effect.

Microstructural properties were studied by cross-sec-tional TEM carried out on sample C which was grown at780 °C. Figure 5(a) shows the low magnification TEM imagetaken with g=<0002> and zone axis of near [1–100] and noTDs were observed in this image. A TEM image taken with

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Figure 3. (a) SEM image of InGaN nanorods without pulsed growth showing many polycrystals nucleated on the SiO2 mask surface. (b)SEM image of InGaN nanorods with pulsed growth showing improved selectivity. (c) Top-view SEM image of as-grown InGaN nanorodswith pulsed growth. (d) 45°-tilted SEM image of InGaN nanorods after removing the SiO2 mask.

Figure 4. (a) PL spectrums of InGaN nanorods with different In compositions measured at RT. (b) XRD 2θ/ω scanning curve.

g=<112̄0> is shown in figure 5(b) with pure edge- andmixed-type TDs visible. As shown in figure 5(b), the pureedge- or mixed-type TDs in the underlying GaN template areterminated by the SiO2 mask (labeled by blue arrows) orterminated at the interface of the nanorod/GaN template(labeled by the red arrow). Similar dislocation reduction hasalso been reported by Colby et al [31]. We also observed afew voids at the interface between some nanorods and theGaN template. These voids were attributed to the imperfectnucleation of nanorods grown on the processed GaN templatebecause the GaN template was probably contaminated by themetal mask during nano-pattern processing.

Mapping of the spatial distribution of elements in theInGaN nanorods was performed by EDS, as shown in fig-ure 6. Figure 6(a) shows the cross-sectional high-angleannular-dark-field (HAADF) scanning TEM (STEM) mode.The interface between GaN buffer and InGaN nanorod isindicated by the white arrow in figure 6(a). EDS mapping wascarried out and the distributions of N, Ga and In were shownfigures 6(b) to (d), respectively. In figure 6(d), the InGaNnanorod shows a sharp interface with the GaN nano-buffer(delineated by the dashed line). The In signals on the GaNtemplate, GaN nano-buffer, and SiO2 are attributed to thespattering of InGaN dusts during ion milling in the TEMsample preparation. The height of dislocation-free InGaNnanorods is about 150 nm. This observation is noteworthygiven that the critical thickness for the onset of misfit dis-locations in planar InGaN growth is typically less than 10 nmfor an indium composition between 10 and 20% [32, 33].

EDS line scan was also carried out to exhibit the In dis-tribution, as shown by the red array in figure 6(a), and the EDScount of In atoms is shown in figure 7. We observe that theEDS intensity of In atoms fluctuates, indicating the non-uni-formity of In composition in InGaN nanorods. The In com-position in InGaN nanorods shows an oscillatory behavioralong the growth direction with an amplitude of about 2.6%. Itis likely that the large difference in bond strength [34, 35] andbond length (strain) between GaN and InN, as well as thedifference in surface kinetics including adsorption, diffusion,desorption, and incorporation, work together to create a

dynamic situation where the concentration of In (or Ga) ada-toms oscillates periodically. Similar spontaneous compositionalfluctuation along the longitudinal axis has been reported inAlGaN nanowires grown by MBE [36]. This phenomenonhappens when one of the metal constituents has a desorptivetendency during growth. The surface concentration of thismetal then undergoes a cyclic process of accumulating (notincorporating) and depleting (incorporating) that accounts forthe observed cyclic fluctuations in composition, as illustrated infigure 9 of reference [36]. It should be noted that the observedfluctuation of In composition is only observable over the lengthscale of 100 nm and has not been reported in InGaN quantumwell samples.

4. Conclusions

In conclusion, ternary InGaN nanorods were prepared ondielectric-masked nano-holes with selective area MOCVD.To overcome the tendency for random nucleation of GaN atlow temperatures, a pulsed growth procedure wasintroduced to enhance the diffusion length of Ga adatomson SiO2, resulting in good selectivity at typical temperatureranges for InGaN. PL from the InGaN nanorods was tunedfrom near ultraviolet (400 nm) to blue-green (∼500 nm).Microstructural properties were characterized by TEM;TDs from the underlying GaN templates were terminated atthe nanorod/template interface, resulting in dislocation-freenanorods. The nanorod geometry is useful in extending thecritical thickness of InGaN heteroepitaxy on GaN. Thecomposition profile of In along the growth directionwas examined by EDS x-ray spectroscopic mapping andline scan. Oscillations of In composition along the growthdirection were observed and are likely due to thekinetic competition of incorporation between In andGa adatoms. These InGaN nanorods are expected to beuseful as templates for growing higher In compositionnano-LEDs.

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Figure 5. Cross-sectional TEM image with (a): g =<0002> and (b): g=<11–20>.

Acknowledgments

The authors gratefully thank Elison Matioli at MassachusettsInstitute of Technology for patterning nano-substrates byinterference lithography. This research was supported by theUS Department of Energy, Office of Basic Energy Sciences,Division of Materials Sciences and Engineering under Award#DE-SC0001134. Facilities used were supported by the YaleInstitute for Nanoscience and Quantum Engineering and NSFMRSEC DMR 1119826.

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