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Growth, structural and optical properties of ternary InGaN nanorods prepared by selectivearea metalorganic chemical vapor deposition
This content has been downloaded from IOPscience. Please scroll down to see the full text. 2014 Nanotechnology 25 225602 (http://iopscience.iop.org/0957-4484/25/22/225602) View the table of contents for this issue, or go to the journal homepage for more
Download details: IP Address: 95.68.109.7 This content was downloaded on 12/05/2014 at 12:27
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Nanotechnology Nanotechnology 25 (2014) 225602 (7pp)
doi:10.1088/0957-4484/25/22/225602
Growth, structural and optical properties of ternary InGaN nanorods prepared by selective-area metalorganic chemical vapor deposition Jie Song, Benjamin Leung, Yu Zhang and Jung Han Department of Electrical Engineering, Yale University, New Haven, CT 06511, USA E-mail: [email protected] Received 5 September 2013, revised 4 March 2014 Accepted for publication 1 April 2014 Published 8 May 2014 Abstract
Ternary InGaN nanorods were prepared on dielectric-masked nano-holes with selective area metalorganic chemical vapor deposition. To overcome the tendency for random nucleation of GaN at low temperatures, a pulsed growth procedure was introduced to enhance the diffusion length of Ga adatoms on SiO2, resulting in good selectivity at typical temperature ranges for InGaN. Photoluminescence from the InGaN nanorods can be tuned from near ultraviolet (400 nm) to blue-green (∼500 nm). Microstructural properties were characterized by transmission electron microscopy; threading dislocations from the underlying GaN template were terminated at the nanorod/template interface, resulting in dislocation-free nanorods. The height of dislocation-free InGaN nanorods is about 150 nm, which is much larger than the critical thickness for the onset of misfit dislocations in planar InGaN growth with typical thickness of less than 10 nm for an indium composition between 10 and 20%. The composition profile of In along the growth direction was examined by energy dispersive x-ray spectroscopic mapping and line scan. Oscillations of In composition along the growth direction were observed and are likely due to the kinetic competition between In and Ga adatoms. These InGaN nanorods are expected to be useful as templates for growing higher In composition nano-light-emitting diodes. Keywords: InGaN, nanorods, MOCVD, pulsed growth (Some figures may appear in colour only in the online journal) 1. Introduction
However, there have been a few intriguing reports of InGaN nanostructures, in the form of InGaN nanowires [7], InGaN nanorings and nanoarrays [8], and InGaN/GaN core-shell structure nanorods [9, 10], which suggest that high In-content InGaN is not inherently defective if the issues of strain and lattice mismatch can be circumvented. Thus we propose to grow InGaN nanorods as nano-scale templates to support the synthesis of higher In-composition InGaN multiple quantum wells (MQWs) and LEDs. Unlike kinetics-dominated molecular beam epitaxy (MBE) [11] or catalyst-mediated growth [12], metalorganic chemical vapor deposition (MOCVD) does not support bottom-up growth of one-dimensional nanostructures. To create InGaN nanorods, we adopted a strategy in
The internal quantum efficiency of conventional GaN-based light-emitting diodes (LEDs) drops precipitously towards longer wavelengths as the In composition increases, a phenomenon known as the ‘green gap’ [1]. Explanations invoke mechanisms such as the immiscibility gap of InGaN [2], the difficulty of incorporating In under compressive strain (compositional pulling) [3], and large internal polarization fields [4]. All these mechanisms are related directly or indirectly to the difficulty of planar heteroepitaxy of InGaN on GaN under a high compressive strain, which leads to poor morphology and deteriorating microstructures [2, 5, 6]. 0957-4484/14/225602+07$33.00
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nanoscale confined epitaxy [13, 14] where nanoscale cylindrical holes are created through dielectric masks over GaN layers. InGaN is then ‘funneled’ into the nanoscale holes through selective epitaxy. Anticipated benefits of using InGaN nanorods include (1) the elimination of the pre-existing dislocations in the underlying template from aspect-ratio engineering [15], and (2) the possibility of strain relaxation elastically based on geometric effect [16] to minimize the compositional pulling effect for high In compositional InGaN. Growth of binary GaN nanorods has been demonstrated by several groups [13, 17–20]. Core-shell nano-LEDs [10, 21] and axial InGaN MQWs [22] grown on GaN nanorods have also been reported. In this paper we report the growth of InGaN nanorods by nanoscale selective-area MOCVD. To overcome the challenge of much reduced surface diffusion for Ga adatoms during InGaN growth, a pulsed growth technique is demonstrated. InGaN nanorods with In compositions up to 20% are prepared at a thickness much greater than the critical thickness in planar epitaxy. Threading dislocations (TDs) propagating from the underlying GaN templates are terminated at the nanorod/underlayer interface, resulting in dislocation-free InGaN nanorods. The height of dislocation-free InGaN nanorods is about 150 nm, which is much larger than the critical thickness for the onset of misfit dislocations in planar InGaN where growth is typically less than 10 nm for an indium composition between 10 to 20%. These InGaN nanorods are expected to be useful as templates for growing higher In compositional nano-LEDs.
Figure 1. Schematic diagrams of the fabrication method for nano-
patterned substrates: (a) deposit 229 nm SiO2 on GaN/sapphire template; (b) spin 180 nm HMDS and bake for 90 s at 110 °C; (c) develop nano-pattern in HMDS interference lithography and then deposit metal mask (5 nm Ti + 40 nm Ni); (d) lift off the metal mask on HMDS in acetone; (e) transfer the nano-pattern into SiO2 by RIE; (f) remove residual metal mask by dry etching.
The growth was carried out by MOCVD (Aixtron 200-4 RF/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 first grown at 1030 °C in a mixture of N2 (3 slm) and NH3 (3 slm) to ensure that InGaN nanoepitaxy was carried out in a consistent way. The GaN nano-buffer showed a flat top surface with a thickness of around 0.1 μm as shown in figure 2(b). The temperature was then lowered down to 780–850 °C to grow InGaN. Different In compositions of InGaN nanorods were obtained by varying the growth temperature and indium composition in the gas phase. Microscopic morphology was examined by a Hitachi scanning electron microscope (SEM). Double crystal x-ray diffraction (XRD) was carried out using a Bede D1 diffractometer. Optical properties were measured by photoluminescence (PL) using a 325 nm He-Cd laser at room temperature (RT). Microstructural properties were characterized by transmission electron microscopy (TEM) and energy dispersive x-ray spectroscopy (EDS) with a FEI F20 microscope operating at an acceleration voltage of 200 kV. TEM specimens were prepared by mechanical polishing first followed by Ar-ion milling.
2. Experiment details Nano-patterned substrates were fabricated on MOCVD grown GaN/sapphire templates according to a procedure illustrated in figure 1. First, a SiO2 dielectric mask with a thickness of about 0.2 μm was deposited on a 2-μm-thick unintentionally doped c-plane GaN/sapphire template by plasma-enhanced chemical vapor deposition, as shown in figure 1(a). We note that the thickness of SiO2 determines the height of the nanorods. Interference lithography (IL) was used to create periodic 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 layer with a photo flux of 45 μJ for 6 min [23]. Once the nanopatterns by IL were developed, a Ti/Ni (5 nm/40 nm) Ni etching mask was deposited followed by lifting off in acetone and the residual nano-patterned metal as a mask was left, as shown in figures 1(d) and (e). The nano-pattern was then transferred into SiO2 by using fluorine-based reactive ion etching (RIE), and finally the metal mask was removed by chlorine-based dry etching, as shown in figure 1(d). Figure 2(a) shows the SEM image of the processed substrate with nano-holes with a diameter of 100 nm and pitch of 300 nm. Before loading the sample into the MOCVD reactor chamber, the substrates were cleaned in piranha (H2SO4: H2O2 = 3:1) and 15% HCl solution. The substrates were then rinsed in deionized water for 5 min.
3. Results and discussion The possibility of selective growth of InGaN was investigated recently by Shioda et al [24]. The issues involved include (1) the competition between the kinetic incorporation of In(Ga)N and the thermodynamic formation of In droplets, (2) the tendency for In(N) desorption at elevated temperatures, exacerbated by the presence of hydrogen (from H2 or NH3), (3) 2
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Nanotechnology 25 (2014) 225602
Figure 2. (a) SEM image of nano-patterned substrate. (b) 45°-tilted SEM image of GaN nanorod buffer.
the random nucleation of GaN crystals on mask due to a high supersaturation in typical InGaN growth conditions, and (4) a large disparity in surface diffusion between In and Ga adatoms. The critical task in the selective growth of InGaN is therefore to enhance the surface diffusion of Ga adatoms on the SiO2 mask surface. As surface diffusion is intimately related to the atomistic configuration of a growing surface, a plausible way to control the surface diffusion is by alternating the precursor supply intervals in pulsed mode to allow substantial changes of surface stoichiometry and reconstructions not typically attainable in continuous growth. Such a pulsed growth procedure has been found to be effective in various material systems [25–28]. During InGaN growth, the gallium precursor (TEGa) was introduced in a pulsed way while TMIn and NH3 flows were kept constant. The interruption duration of TEGa was chosen to be between 3 and 5 s. We found that if the interruption duration of TEGa was too short, the selectivity couldn’t be improved. If the interruption duration of TEGa was too long, the In in the InGaN nanorods would reevaporate and accumulate on the surface to form In droplets. Figure 3(a) shows the morphology of InGaN nanorods grown at 780 °C with a gas-phase ratio of 59% under the conventional continuous-growth method. A high density of polycrystalline GaN was observed on the SiO2 mask surface; the majority of InGaN growth takes place as random nucleation on the mask rather than proceeding epitaxially within the nanopatterned holes (marked by arrows as shown in figure 3(a)). We also noted that the PL intensity of the sample corresponding to figure 3(a) was very weak and there was no detectable InGaN peak in XRD. Based on the distribution of polycrystalline GaN islands on the SiO2 mask surface around the opening holes, the surface diffusion length of 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 noticeably improved with the pulse procedure. Figure 3(b) shows the SEM image of the InGaN nanorods grown with pulsed growth while keeping all other parameters the same as the sample of figure 3(a). The selectivity of InGaN over SiO2 is
achieved, implying that the surface diffusion length of Ga adatoms 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 also observed that a greater amount of InGaN nanorods grew in the holes with pulsed growth than in those without pulsed growth, as shown in figure 3(b). Figure 3(c) is the top-view SEM image of the as-grown sample by pulsed growth and InGaN nanorods exhibiting pyramidal top surface bounded by six {10-11} facets. Figure 3(d) shows a 45°-tilted SEM image of the InGaN nanorods after removing the SiO2 where the height of each nanorod is about 250 nm. Longer InGaN nanorods can conceivably be prepared with a thicker SiO2 nano-patterned layer using the same procedure for selective growth. The as-grown samples were characterized by PL at RT. Three samples A, B and C were grown at temperatures of 850, 810 and 780 °C, respectively, with a constant In/(In + Ga) gas phase ratio of 59%. Sample D was grown at 780 °C with an increased gas phase ratio of 87%. As shown in figure 4(a), the PL peaks of A–D shift from 406 to 492 nm. We also tried lowering growth temperature further to 750 °C and obtained higher In compositional InGaN nanorods with good selectivity. However, PL intensity degraded due to possibly a high concentration of point defects under nonoptimal growth conditions. Below 750 °C we started to observe a loss of selectivity. Based on the PL results, In compositions in InGaN nanorods were calculated using the formula Eg = 0.67x + 3.42 ( 1 − x ) − bx ( 1 − x ), 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 and D are 8.5%, 11.7%, 17.7% and 20.5%, respectively. The sample 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° is attributed to the GaN template and the InGaN peak is located at ∼34.1°. A theoretical work predicts that, given the thickness and diameter of the InGaN nanorods, 90% of the 3
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Nanotechnology 25 (2014) 225602
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 nanorods with 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.
mismatched strain should be elastically relaxed [16]. By assuming that the InGaN nanorods are fully relaxed, the In composition in InGaN nanorods is calculated to be 12.3%. The discrepancy between In compositions calculated by XRD results and PL results is likely due to the In localization effect.
Microstructural properties were studied by cross-sectional TEM carried out on sample C which was grown at 780 °C. Figure 5(a) shows the low magnification TEM image taken with g = and zone axis of near [1–100] and no TDs were observed in this image. A TEM image taken with 4
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Nanotechnology 25 (2014) 225602
Figure 5. Cross-sectional TEM image with (a): g = and (b): g = .
g = is shown in figure 5(b) with pure edge- and mixed-type TDs visible. As shown in figure 5(b), the pure edge- or mixed-type TDs in the underlying GaN template are terminated by the SiO2 mask (labeled by blue arrows) or terminated at the interface of the nanorod/GaN template (labeled by the red arrow). Similar dislocation reduction has also been reported by Colby et al [31]. We also observed a few voids at the interface between some nanorods and the GaN template. These voids were attributed to the imperfect nucleation of nanorods grown on the processed GaN template because the GaN template was probably contaminated by the metal mask during nano-pattern processing. Mapping of the spatial distribution of elements in the InGaN nanorods was performed by EDS, as shown in figure 6. Figure 6(a) shows the cross-sectional high-angle annular-dark-field (HAADF) scanning TEM (STEM) mode. The interface between GaN buffer and InGaN nanorod is indicated by the white arrow in figure 6(a). EDS mapping was carried out and the distributions of N, Ga and In were shown figures 6(b) to (d), respectively. In figure 6(d), the InGaN nanorod shows a sharp interface with the GaN nano-buffer (delineated by the dashed line). The In signals on the GaN template, GaN nano-buffer, and SiO2 are attributed to the spattering of InGaN dusts during ion milling in the TEM sample preparation. The height of dislocation-free InGaN nanorods is about 150 nm. This observation is noteworthy given that the critical thickness for the onset of misfit dislocations in planar InGaN growth is typically less than 10 nm for an indium composition between 10 and 20% [32, 33]. EDS line scan was also carried out to exhibit the In distribution, as shown by the red array in figure 6(a), and the EDS count of In atoms is shown in figure 7. We observe that the EDS intensity of In atoms fluctuates, indicating the non-uniformity of In composition in InGaN nanorods. The In composition in InGaN nanorods shows an oscillatory behavior along the growth direction with an amplitude of about 2.6%. It is likely that the large difference in bond strength [34, 35] and bond length (strain) between GaN and InN, as well as the difference in surface kinetics including adsorption, diffusion, desorption, and incorporation, work together to create a
dynamic situation where the concentration of In (or Ga) adatoms oscillates periodically. Similar spontaneous compositional fluctuation along the longitudinal axis has been reported in AlGaN nanowires grown by MBE [36]. This phenomenon happens when one of the metal constituents has a desorptive tendency during growth. The surface concentration of this metal then undergoes a cyclic process of accumulating (not incorporating) and depleting (incorporating) that accounts for the observed cyclic fluctuations in composition, as illustrated in figure 9 of reference [36]. It should be noted that the observed fluctuation of In composition is only observable over the length scale of 100 nm and has not been reported in InGaN quantum well samples.
4. Conclusions In conclusion, ternary InGaN nanorods were prepared on dielectric-masked nano-holes with selective area MOCVD. To overcome the tendency for random nucleation of GaN at low temperatures, a pulsed growth procedure was introduced to enhance the diffusion length of Ga adatoms on SiO2, resulting in good selectivity at typical temperature ranges for InGaN. PL from the InGaN nanorods was tuned from near ultraviolet (400 nm) to blue-green (∼500 nm). Microstructural properties were characterized by TEM; TDs from the underlying GaN templates were terminated at the nanorod/template interface, resulting in dislocation-free nanorods. The nanorod geometry is useful in extending the critical thickness of InGaN heteroepitaxy on GaN. The composition profile of In along the growth direction was examined by EDS x-ray spectroscopic mapping and line scan. Oscillations of In composition along the growth direction were observed and are likely due to the kinetic competition of incorporation between In and Ga adatoms. These InGaN nanorods are expected to be useful as templates for growing higher In composition nano-LEDs. 5
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Nanotechnology 25 (2014) 225602
Figure 6. (a) STEM HAADF image of a InGaN nanorod. (b)–(d) The EDS mapping results of N, Ga and In, respectively.
Acknowledgments The authors gratefully thank Elison Matioli at Massachusetts Institute of Technology for patterning nano-substrates by interference lithography. This research was supported by the US Department of Energy, Office of Basic Energy Sciences, Division of Materials Sciences and Engineering under Award #DE-SC0001134. Facilities used were supported by the Yale Institute for Nanoscience and Quantum Engineering and NSF MRSEC DMR 1119826. References [1] Krames M R, Shchekin O B, Mueller-Mach R, Muller G O, Zhou L, Harber G and Craford M G 2007 J. Display Technol. 3 160
Figure 7. EDS count of In taken from the line scan as shown by the
red array in figure 6(a).
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[20] Lin Y T, Yeh T W and Dapkus P D 2012 Nanotechnology 23 465601 [21] Hong Y J, Lee C H, Yoon A, Kim M, Seong H K, Chung H J, Sone C, Park Y J and Yi G C 2011 Adv. Mater. 23 3284 [22] Zang K Y, Wang Y D and Chua S J 2009 Phys. Stat. Sol. (c) 6 S514 [23] Matioli E, Brinkley S, Kelchner K M, Hu Y L, Nakamura S, Denbaars S, Speck J and Weisbuch C 2012 Light: Science & Applications 1 e22 [24] Shioda T, Sugiyama M, Shimogaki Y and Nakano Y 2009 J Cryst. Growth 311 2809 [25] Briones F and Horikoshi Y 1990 J. Appl. Phys. 29 1014 [26] Chen Z, Qhalid Fareed R S, Gaevski M, Adivarahan V, Yang J W, Khan A, Mei J and Ponce F A 2006 Appl. Phys. Lett. 89 081905 [27] Liu C, Shields P A, Denchitcharoen S, Stepanov S, Gott A and Wang W N 2007 J. Crystl. Growth 300 104 [28] Jindal V, Tripathi N, Tungare M, Paschos O, Haldar P and Shahedipour-Sandvik F 2008 Phys. Stat. Sol. (c) 5 1709 [29] Lee S C and Brueck S R J 2009 Appl. Phys. Lett. 94 153110 [30] Moses P G and Van de Walle C G 2010 Appl. Phys. Lett. 96 021908 [31] Colby R, Liang Z, Wildeson I H, Ewoldt D A, Sands T D, Edwin García R and Stach E A 2010 Nano Lett 10 1568 [32] Holec D, Costa P M F J, Kappers M J and Humphreys C J 2007 J. Crystl. Growth 303 314 [33] Reed M J, El-Masry N A, Parker C A, Roberts J C and Bedair S M 2000 Appl. Phys. Lett. 77 4121 [34] Adelmann C, Langer R, Feuillet G and Daudin B 1999 Appl. Phys. Lett. 75 3518 [35] Bedair S M, McIntosh F G, Roberts J C, Piner E L, Boutros K S and EI-Marsy N A 1997 J. Crystal. Growth 178 32 [36] Pierret A, Bougerol C, Murcia-Mascaros S, Cros A, Renevier H, Gayral B and Daudin B 2013 Nanotechnology 24 115704
[2] Rao M, Kim D and Mahajan S 2004 Appl. Phys. Lett. 85 1961 [3] Pereira S, Correia M R, Pereira E, O’Donnell K P, Trager-Cowan C, Sweeney F and Alves E 2001 Phys. Rev. B 64 205311 [4] Fuhrmann D, Netzel C, Rossow U, Hangleiter A, Ade G and Hinze P 2006 Appl. Phys. Lett. 88 071105 [5] Kawaguchi Y et al 1998 J. Cryst. Growth 189/190 24 [6] Cho H K, Lee J Y, Kim K S and Yang G M 2000 J. Cryst. Growth 220 197 [7] Kuykendall T, Ulrich P, Aloni S and Yang P 2007 Nat. Mater. 6 951 [8] Wang Y D, Zang K Y, Chua S J, Sander M S, Tripathy S and Fonstad C G 2006 J. Phys. Chem. B 110 11081 [9] Li Q and Wang G T 2010 Appl. Phys. Lett. 97 181107 [10] Sekiguchi H, Kishino K and Kikuchi A 2010 Appl. Phys. Lett. 96 231104 [11] Guo W, Zhang M, Banerjee A and Bhattacharya P 2010 Nano Lett. 10 3355 [12] Wu Z H, Mei X Y, Kim D, Blumin M and Ruda H E 2002 Appl. Phys. Lett. 81 5177 [13] Wang X, Sun X Y, Fairchild M and Hersee S D 2006 Appl. Phys. Lett. 89 233115 [14] Wang Y D, Zang K Y and Chua S J 2006 J. Appl. Phys. 100 054306 [15] Langdo T A, Leitz C W, Currie M T, Fitzgerald E A, Lochtefeld A and Antoniadis D A 2000 Appl. Phys. Lett. 76 3700 [16] Alizadeh A, Sharma P, Ganti S, LeBoeuf S F and Tsakalakos L 2004 J. Appl. Phys. 95 8199 [17] Sekiguchi H, Kishino K and Kikuchi A 2008 Appl. Phys. Express 1 124002 [18] Bergbauer W et al 2010 Nanotechnology 21 305201 [19] Bertness K A, Sanders A W, Rourke D M, Harvey T E, Roshko A, Schlager J B and Sanford N A 2010 Adv. Funct. Mater. 20 2911
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Growth, structural and optical properties of ternary InGaN nanorods prepared by selectivearea metalorganic chemical vapor deposition
This content has been downloaded from IOPscience. Please scroll down to see the full text. 2014 Nanotechnology 25 225602 (http://iopscience.iop.org/0957-4484/25/22/225602) View the table of contents for this issue, or go to the journal homepage for more
Download details: IP Address: 95.68.109.7 This content was downloaded on 12/05/2014 at 12:27
Please note that terms and conditions apply.
Nanotechnology Nanotechnology 25 (2014) 225602 (7pp)
doi:10.1088/0957-4484/25/22/225602
Growth, structural and optical properties of ternary InGaN nanorods prepared by selective-area metalorganic chemical vapor deposition Jie Song, Benjamin Leung, Yu Zhang and Jung Han Department of Electrical Engineering, Yale University, New Haven, CT 06511, USA E-mail: [email protected] Received 5 September 2013, revised 4 March 2014 Accepted for publication 1 April 2014 Published 8 May 2014 Abstract
Ternary InGaN nanorods were prepared on dielectric-masked nano-holes with selective area metalorganic chemical vapor deposition. To overcome the tendency for random nucleation of GaN at low temperatures, a pulsed growth procedure was introduced to enhance the diffusion length of Ga adatoms on SiO2, resulting in good selectivity at typical temperature ranges for InGaN. Photoluminescence from the InGaN nanorods can be tuned from near ultraviolet (400 nm) to blue-green (∼500 nm). Microstructural properties were characterized by transmission electron microscopy; threading dislocations from the underlying GaN template were terminated at the nanorod/template interface, resulting in dislocation-free nanorods. The height of dislocation-free InGaN nanorods is about 150 nm, which is much larger than the critical thickness for the onset of misfit dislocations in planar InGaN growth with typical thickness of less than 10 nm for an indium composition between 10 and 20%. The composition profile of In along the growth direction was examined by energy dispersive x-ray spectroscopic mapping and line scan. Oscillations of In composition along the growth direction were observed and are likely due to the kinetic competition between In and Ga adatoms. These InGaN nanorods are expected to be useful as templates for growing higher In composition nano-light-emitting diodes. Keywords: InGaN, nanorods, MOCVD, pulsed growth (Some figures may appear in colour only in the online journal) 1. Introduction
However, there have been a few intriguing reports of InGaN nanostructures, in the form of InGaN nanowires [7], InGaN nanorings and nanoarrays [8], and InGaN/GaN core-shell structure nanorods [9, 10], which suggest that high In-content InGaN is not inherently defective if the issues of strain and lattice mismatch can be circumvented. Thus we propose to grow InGaN nanorods as nano-scale templates to support the synthesis of higher In-composition InGaN multiple quantum wells (MQWs) and LEDs. Unlike kinetics-dominated molecular beam epitaxy (MBE) [11] or catalyst-mediated growth [12], metalorganic chemical vapor deposition (MOCVD) does not support bottom-up growth of one-dimensional nanostructures. To create InGaN nanorods, we adopted a strategy in
The internal quantum efficiency of conventional GaN-based light-emitting diodes (LEDs) drops precipitously towards longer wavelengths as the In composition increases, a phenomenon known as the ‘green gap’ [1]. Explanations invoke mechanisms such as the immiscibility gap of InGaN [2], the difficulty of incorporating In under compressive strain (compositional pulling) [3], and large internal polarization fields [4]. All these mechanisms are related directly or indirectly to the difficulty of planar heteroepitaxy of InGaN on GaN under a high compressive strain, which leads to poor morphology and deteriorating microstructures [2, 5, 6]. 0957-4484/14/225602+07$33.00
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Nanotechnology 25 (2014) 225602
nanoscale confined epitaxy [13, 14] where nanoscale cylindrical holes are created through dielectric masks over GaN layers. InGaN is then ‘funneled’ into the nanoscale holes through selective epitaxy. Anticipated benefits of using InGaN nanorods include (1) the elimination of the pre-existing dislocations in the underlying template from aspect-ratio engineering [15], and (2) the possibility of strain relaxation elastically based on geometric effect [16] to minimize the compositional pulling effect for high In compositional InGaN. Growth of binary GaN nanorods has been demonstrated by several groups [13, 17–20]. Core-shell nano-LEDs [10, 21] and axial InGaN MQWs [22] grown on GaN nanorods have also been reported. In this paper we report the growth of InGaN nanorods by nanoscale selective-area MOCVD. To overcome the challenge of much reduced surface diffusion for Ga adatoms during InGaN growth, a pulsed growth technique is demonstrated. InGaN nanorods with In compositions up to 20% are prepared at a thickness much greater than the critical thickness in planar epitaxy. Threading dislocations (TDs) propagating from the underlying GaN templates are terminated at the nanorod/underlayer interface, resulting in dislocation-free InGaN nanorods. The height of dislocation-free InGaN nanorods is about 150 nm, which is much larger than the critical thickness for the onset of misfit dislocations in planar InGaN where growth is typically less than 10 nm for an indium composition between 10 to 20%. These InGaN nanorods are expected to be useful as templates for growing higher In compositional nano-LEDs.
Figure 1. Schematic diagrams of the fabrication method for nano-
patterned substrates: (a) deposit 229 nm SiO2 on GaN/sapphire template; (b) spin 180 nm HMDS and bake for 90 s at 110 °C; (c) develop nano-pattern in HMDS interference lithography and then deposit metal mask (5 nm Ti + 40 nm Ni); (d) lift off the metal mask on HMDS in acetone; (e) transfer the nano-pattern into SiO2 by RIE; (f) remove residual metal mask by dry etching.
The growth was carried out by MOCVD (Aixtron 200-4 RF/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 first grown at 1030 °C in a mixture of N2 (3 slm) and NH3 (3 slm) to ensure that InGaN nanoepitaxy was carried out in a consistent way. The GaN nano-buffer showed a flat top surface with a thickness of around 0.1 μm as shown in figure 2(b). The temperature was then lowered down to 780–850 °C to grow InGaN. Different In compositions of InGaN nanorods were obtained by varying the growth temperature and indium composition in the gas phase. Microscopic morphology was examined by a Hitachi scanning electron microscope (SEM). Double crystal x-ray diffraction (XRD) was carried out using a Bede D1 diffractometer. Optical properties were measured by photoluminescence (PL) using a 325 nm He-Cd laser at room temperature (RT). Microstructural properties were characterized by transmission electron microscopy (TEM) and energy dispersive x-ray spectroscopy (EDS) with a FEI F20 microscope operating at an acceleration voltage of 200 kV. TEM specimens were prepared by mechanical polishing first followed by Ar-ion milling.
2. Experiment details Nano-patterned substrates were fabricated on MOCVD grown GaN/sapphire templates according to a procedure illustrated in figure 1. First, a SiO2 dielectric mask with a thickness of about 0.2 μm was deposited on a 2-μm-thick unintentionally doped c-plane GaN/sapphire template by plasma-enhanced chemical vapor deposition, as shown in figure 1(a). We note that the thickness of SiO2 determines the height of the nanorods. Interference lithography (IL) was used to create periodic 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 layer with a photo flux of 45 μJ for 6 min [23]. Once the nanopatterns by IL were developed, a Ti/Ni (5 nm/40 nm) Ni etching mask was deposited followed by lifting off in acetone and the residual nano-patterned metal as a mask was left, as shown in figures 1(d) and (e). The nano-pattern was then transferred into SiO2 by using fluorine-based reactive ion etching (RIE), and finally the metal mask was removed by chlorine-based dry etching, as shown in figure 1(d). Figure 2(a) shows the SEM image of the processed substrate with nano-holes with a diameter of 100 nm and pitch of 300 nm. Before loading the sample into the MOCVD reactor chamber, the substrates were cleaned in piranha (H2SO4: H2O2 = 3:1) and 15% HCl solution. The substrates were then rinsed in deionized water for 5 min.
3. Results and discussion The possibility of selective growth of InGaN was investigated recently by Shioda et al [24]. The issues involved include (1) the competition between the kinetic incorporation of In(Ga)N and the thermodynamic formation of In droplets, (2) the tendency for In(N) desorption at elevated temperatures, exacerbated by the presence of hydrogen (from H2 or NH3), (3) 2
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Figure 2. (a) SEM image of nano-patterned substrate. (b) 45°-tilted SEM image of GaN nanorod buffer.
the random nucleation of GaN crystals on mask due to a high supersaturation in typical InGaN growth conditions, and (4) a large disparity in surface diffusion between In and Ga adatoms. The critical task in the selective growth of InGaN is therefore to enhance the surface diffusion of Ga adatoms on the SiO2 mask surface. As surface diffusion is intimately related to the atomistic configuration of a growing surface, a plausible way to control the surface diffusion is by alternating the precursor supply intervals in pulsed mode to allow substantial changes of surface stoichiometry and reconstructions not typically attainable in continuous growth. Such a pulsed growth procedure has been found to be effective in various material systems [25–28]. During InGaN growth, the gallium precursor (TEGa) was introduced in a pulsed way while TMIn and NH3 flows were kept constant. The interruption duration of TEGa was chosen to be between 3 and 5 s. We found that if the interruption duration of TEGa was too short, the selectivity couldn’t be improved. If the interruption duration of TEGa was too long, the In in the InGaN nanorods would reevaporate and accumulate on the surface to form In droplets. Figure 3(a) shows the morphology of InGaN nanorods grown at 780 °C with a gas-phase ratio of 59% under the conventional continuous-growth method. A high density of polycrystalline GaN was observed on the SiO2 mask surface; the majority of InGaN growth takes place as random nucleation on the mask rather than proceeding epitaxially within the nanopatterned holes (marked by arrows as shown in figure 3(a)). We also noted that the PL intensity of the sample corresponding to figure 3(a) was very weak and there was no detectable InGaN peak in XRD. Based on the distribution of polycrystalline GaN islands on the SiO2 mask surface around the opening holes, the surface diffusion length of 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 noticeably improved with the pulse procedure. Figure 3(b) shows the SEM image of the InGaN nanorods grown with pulsed growth while keeping all other parameters the same as the sample of figure 3(a). The selectivity of InGaN over SiO2 is
achieved, implying that the surface diffusion length of Ga adatoms 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 also observed that a greater amount of InGaN nanorods grew in the holes with pulsed growth than in those without pulsed growth, as shown in figure 3(b). Figure 3(c) is the top-view SEM image of the as-grown sample by pulsed growth and InGaN nanorods exhibiting pyramidal top surface bounded by six {10-11} facets. Figure 3(d) shows a 45°-tilted SEM image of the InGaN nanorods after removing the SiO2 where the height of each nanorod is about 250 nm. Longer InGaN nanorods can conceivably be prepared with a thicker SiO2 nano-patterned layer using the same procedure for selective growth. The as-grown samples were characterized by PL at RT. Three samples A, B and C were grown at temperatures of 850, 810 and 780 °C, respectively, with a constant In/(In + Ga) gas phase ratio of 59%. Sample D was grown at 780 °C with an increased gas phase ratio of 87%. As shown in figure 4(a), the PL peaks of A–D shift from 406 to 492 nm. We also tried lowering growth temperature further to 750 °C and obtained higher In compositional InGaN nanorods with good selectivity. However, PL intensity degraded due to possibly a high concentration of point defects under nonoptimal growth conditions. Below 750 °C we started to observe a loss of selectivity. Based on the PL results, In compositions in InGaN nanorods were calculated using the formula Eg = 0.67x + 3.42 ( 1 − x ) − bx ( 1 − x ), 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 and D are 8.5%, 11.7%, 17.7% and 20.5%, respectively. The sample 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° is attributed to the GaN template and the InGaN peak is located at ∼34.1°. A theoretical work predicts that, given the thickness and diameter of the InGaN nanorods, 90% of the 3
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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 nanorods with 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.
mismatched strain should be elastically relaxed [16]. By assuming that the InGaN nanorods are fully relaxed, the In composition in InGaN nanorods is calculated to be 12.3%. The discrepancy between In compositions calculated by XRD results and PL results is likely due to the In localization effect.
Microstructural properties were studied by cross-sectional TEM carried out on sample C which was grown at 780 °C. Figure 5(a) shows the low magnification TEM image taken with g = and zone axis of near [1–100] and no TDs were observed in this image. A TEM image taken with 4
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Figure 5. Cross-sectional TEM image with (a): g = and (b): g = .
g = is shown in figure 5(b) with pure edge- and mixed-type TDs visible. As shown in figure 5(b), the pure edge- or mixed-type TDs in the underlying GaN template are terminated by the SiO2 mask (labeled by blue arrows) or terminated at the interface of the nanorod/GaN template (labeled by the red arrow). Similar dislocation reduction has also been reported by Colby et al [31]. We also observed a few voids at the interface between some nanorods and the GaN template. These voids were attributed to the imperfect nucleation of nanorods grown on the processed GaN template because the GaN template was probably contaminated by the metal mask during nano-pattern processing. Mapping of the spatial distribution of elements in the InGaN nanorods was performed by EDS, as shown in figure 6. Figure 6(a) shows the cross-sectional high-angle annular-dark-field (HAADF) scanning TEM (STEM) mode. The interface between GaN buffer and InGaN nanorod is indicated by the white arrow in figure 6(a). EDS mapping was carried out and the distributions of N, Ga and In were shown figures 6(b) to (d), respectively. In figure 6(d), the InGaN nanorod shows a sharp interface with the GaN nano-buffer (delineated by the dashed line). The In signals on the GaN template, GaN nano-buffer, and SiO2 are attributed to the spattering of InGaN dusts during ion milling in the TEM sample preparation. The height of dislocation-free InGaN nanorods is about 150 nm. This observation is noteworthy given that the critical thickness for the onset of misfit dislocations in planar InGaN growth is typically less than 10 nm for an indium composition between 10 and 20% [32, 33]. EDS line scan was also carried out to exhibit the In distribution, as shown by the red array in figure 6(a), and the EDS count of In atoms is shown in figure 7. We observe that the EDS intensity of In atoms fluctuates, indicating the non-uniformity of In composition in InGaN nanorods. The In composition in InGaN nanorods shows an oscillatory behavior along the growth direction with an amplitude of about 2.6%. It is likely that the large difference in bond strength [34, 35] and bond length (strain) between GaN and InN, as well as the difference in surface kinetics including adsorption, diffusion, desorption, and incorporation, work together to create a
dynamic situation where the concentration of In (or Ga) adatoms oscillates periodically. Similar spontaneous compositional fluctuation along the longitudinal axis has been reported in AlGaN nanowires grown by MBE [36]. This phenomenon happens when one of the metal constituents has a desorptive tendency during growth. The surface concentration of this metal then undergoes a cyclic process of accumulating (not incorporating) and depleting (incorporating) that accounts for the observed cyclic fluctuations in composition, as illustrated in figure 9 of reference [36]. It should be noted that the observed fluctuation of In composition is only observable over the length scale of 100 nm and has not been reported in InGaN quantum well samples.
4. Conclusions In conclusion, ternary InGaN nanorods were prepared on dielectric-masked nano-holes with selective area MOCVD. To overcome the tendency for random nucleation of GaN at low temperatures, a pulsed growth procedure was introduced to enhance the diffusion length of Ga adatoms on SiO2, resulting in good selectivity at typical temperature ranges for InGaN. PL from the InGaN nanorods was tuned from near ultraviolet (400 nm) to blue-green (∼500 nm). Microstructural properties were characterized by TEM; TDs from the underlying GaN templates were terminated at the nanorod/template interface, resulting in dislocation-free nanorods. The nanorod geometry is useful in extending the critical thickness of InGaN heteroepitaxy on GaN. The composition profile of In along the growth direction was examined by EDS x-ray spectroscopic mapping and line scan. Oscillations of In composition along the growth direction were observed and are likely due to the kinetic competition of incorporation between In and Ga adatoms. These InGaN nanorods are expected to be useful as templates for growing higher In composition nano-LEDs. 5
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Figure 6. (a) STEM HAADF image of a InGaN nanorod. (b)–(d) The EDS mapping results of N, Ga and In, respectively.
Acknowledgments The authors gratefully thank Elison Matioli at Massachusetts Institute of Technology for patterning nano-substrates by interference lithography. This research was supported by the US Department of Energy, Office of Basic Energy Sciences, Division of Materials Sciences and Engineering under Award #DE-SC0001134. Facilities used were supported by the Yale Institute for Nanoscience and Quantum Engineering and NSF MRSEC DMR 1119826. References [1] Krames M R, Shchekin O B, Mueller-Mach R, Muller G O, Zhou L, Harber G and Craford M G 2007 J. Display Technol. 3 160
Figure 7. EDS count of In taken from the line scan as shown by the
red array in figure 6(a).
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