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Bottom-Up Nano-heteroepitaxy of Wafer-Scale Semipolar GaN on (001) Si

Jui-Wei Hus , Chien-Chia Chen , Ming-Jui Lee , Hsueh-Hsing Liu , Jen-Inn Chyi , Michael R. S. Huang , Chuan-Pu Liu , Tzu-Chiao Wei , Jr-Hau He , and Kun-Yu Lai *

J.-W. Hus, C.-C. Chen, M.-J. Lee, Prof. K.-Y. Lai Department of Optics and Photonics National Central University Chung-Li 320 , Taiwan E-mail: kylai@ncu.edu.tw Dr. H.-H. Liu, Prof. J.-I. Chyi Department of Electrical Engineering National Central University Chung-Li 320 , Taiwan Dr. M. R. S. Huang, Prof. C.-P. Liu Department of Materials Science and Engineering National Cheng Kung University Tainan 701 , Taiwan T.-C. Wei Institute of Photonics and Optoelectronics National Taiwan University Taipei 106 , Taiwan Prof. J.-H. He Computer, Electrical and Mathematical Sciences and Engineering (CEMSE) Division King Abdullah University of Science and Technology (KAUST) Thuwal 23955-6900 , Saudi Arabia

DOI: 10.1002/adma.201501538

either requires expensive facilities or renders the uniformity of inadequate area, let alone the surface damages that are essen-tially unavoidable in the destructive etching processes. Ideally, the problems can be unraveled by patterning the Si substrate with a bottom-up manner, but naturally controlling structural size/morphology with suffi cient precision and uniformity at the epitaxial interface are extremely diffi cult and rarely reported.

In addition to the unpleasant top-down fabrication, the growth of GaN-on-Si encounters another two challenges: i) To minimize the lattice mismatch with GaN, the Si substrate must be orientated along [111]. [ 3 ] However, (111) Si is not the sub-strate favored in CMOS manufacturing. ii) The huge strain of GaN-on-Si is generally engineered through a thick and slow-growing AlN-based buffer system, [ 3,9 ] which prolongs growth time and often results in severe chamber contamination caused by the parasitic reactions involving Al ions. [ 10 ] The long growth time and contaminated reactor demand increased chamber maintenance and spare parts, sacrifi cing the cost advantage of Si substrates.

In this study, we grow GaN-on-Si with a novel tactic involving bottom-up and Al-free nano-heteroepitaxy. The unique approach is realized by a self-assembled ZnO-nanorod buffer, eluding the costly and corrosive lithography/etching processes. ZnO shares the same crystalline symmetry and similar lattice constant (mismatch ≈ 1.9%) with GaN, [ 11 ] imparting smooth atomic transition to the GaN/ZnO inter-face. [ 12 ] Since the nanorods are synthesized via a templating procedure, being applicable to virtually any surface types, the growth of GaN and ZnO nanorods is performed on (001) Si, which is the substrate truly compatible with CMOS tech-nology. The templating procedure also grants ZnO nanorods the capability to reach wafer-scale uniformity. [ 13 ] GaN grown with this nano-heteroepitaxy technique exhibits pyramidal surface with semipolar {1011} facets, which can improve the internal quantum effi ciencies of III-nitride LEDs in light of the mitigated quantum-confi ned Stark effect. [ 14 ] These results open a new route for GaN-on-Si growth circumventing the technical hurdles that have been haunting researchers for years.

The self-organized nanorods sprout from the ZnO seed layer prepared by sputtering, allowing the nanostructure to grow on any substrates regardless of their mismatch in lattice con-stant or surface chemistry. [ 15 ] The oxide layer also avoids the notorious meltback etching effect often encountered by the growth of GaN on Si with metal–organic chemical vapor dep-osition (MOCVD), during which the reaction between Si and Ga ions destroys the substrate and the epilayer. [ 3 ] Single crystal-linity and the [0002] orientation of the nanorods are confi rmed

The unprecedented luminescence effi ciency of GaN-based solid state lighting has ushered in a new age of human life. [ 1,2 ] Compact blue/green light-emitting diodes (LEDs) nowadays are ubiquitous in the technologies of display and general lighting. In the pursuit of higher performance/cost ratios for III-nitride LEDs, Si has been a charming alternative substrate material to sapphire considering her low cost, high scalability, and the man-ufacturing compatible with complementary metal–oxide–semi-conductor (CMOS) technology. However, the enormous lattice and thermal mismatches between GaN and Si, [ 3 ] necessitating sophisticated growth techniques, have impeded the migration of LED manufacturing from sapphire to Si substrates.

To address the mismatch issues of GaN-on-Si, nano-hetero-epitaxy is one of the most promising approaches. The technique allows the strain energy in a heteroepitaxial layer to be released through lateral and vertical lattice deformation by performing the growth on the substrate patterned in nanoscale. [ 4,5 ] The ultra small contact area and 3D deformation dramatically extend the critical thickness of the epilayer and thereby defer the for-mation of crystal defects. [ 6–8 ] Despite the tempting merits, most of the reported nano-heteroepitaxy schemes fi nd limited prac-tical applications. This is because the patterning of substrate surface is usually attained via the so-called top-down fabrica-tion, in which advanced lithography and a subsequent etching procedure are carried out to form the nanoscale mesas. The top-down method capable of defi ning sub-micrometer patterns

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by X-ray diffraction (XRD) θ –2 θ scan (Figure S1, Supporting Information). One advantage of ZnO nanostructures is that a wide variety of geometric features can be controlled via spe-cifi c synthesis conditions. [ 16 ] To facilitate the nucleation of GaN on the nanostructured buffer, it is highly desirable that the nanorods are of uniform lengths and well-aligned along the [0001] growth direction. The goal can be achieved by annealing the ZnO seed layer with proper conditions. Annealing effect on the geometric features of ZnO nanorods can be seen with the scanning electron microscopy (SEM) images provided in Figure S2 in the Supporting Information. The thermally treated seed layer renders the nanorods with enlarged diameters, enhanced length uniformity, and improved vertical alignment. These differences are attributed to the increased grain sizes of the annealed ZnO, which are revealed by atomic force microscopy images (Figure S3, Supporting Information). The increased grain sizes enable the nanorods to evolve from the enlarged plateau of a grain, instead of the tilted surface at grain bounda-ries, and therefore give rise to the nanorods with enhanced uniformity. [ 17 ]

Since ZnO decomposes at the temperature above 600 °C and is particularly vulnerable in H + ambience, [ 11 ] the growth of GaN is initiated from a low-temperature nucleation layer to protect the nanorods as well as to facilitate lateral coalescence of the epilayer. Figure 1 a,b respectively shows the nucleation layer of AlN and GaN grown at 580 °C for 45 min in N 2 ambience. It

is clear that the nucleation of AlN results in disordered crystal orientations, while the GaN nucleation layer preserves the well alignment of the nanorods and displays the surface of a pyr-amid-like feature. One reason for the superior nucleation of GaN is the smaller lattice mismatch between GaN and ZnO, compared with that between AlN and ZnO. [ 11,18 ] Further, it can be seen that the nanorods covered by AlN are noticeably thinner than those covered by GaN, making the AlN/ZnO structure more vulnerable in the subsequent high-temperature growth. This result indicates that ZnO is partially decomposed during the nucleation of AlN, which is much slower than GaN because of the higher Al–N cohesive energy. [ 19 ] Figure 1 demonstrates that the growth of GaN-on-Si with ZnO-nanorod buffer can be achieved with an Al-free nucleation layer, avoiding the unde-sired chamber contamination by Al complexes.

Figure 2 a presents the substrate temperature plotted as a function of time during the growth of GaN. The 45 min nuclea-tion layer at 580 °C was carried out with a relatively high V/III ratio of 4266 to enhance lateral coalescence. [ 20 ] The substrate was then heated to 1000 °C for the growth of n-type GaN (n-GaN) and the undoped GaN capping layer. The InGaN/GaN multiple quantum wells (MQW) were attained at 785 °C. To evaluate the residual stress in the epilayer, ex situ Raman spec-troscopy was performed at different stages during the growth. Figure 2 b shows the spectra of GaN E 2 -high phonon mode, being highly sensitive to the in-plane strain of the crystal, [ 21 ] of the epilayers at the four stages designated in Figure 2 a.

It can be seen that the E 2 -high peak appears at 568.9 cm −1 after 20 min of the low-temperature nucleation (Stage 1), cor-responding to a stress-free state. [ 8 ] A minor peak at 549.3 cm −1 is also observed in the same spectra, which is ascribed to the signal from the ZnO nanorods. [ 22,23 ] The GaN phonon peak then shifts to lower wavenumbers with increased layer thickness, indicating the buildup of tensile stress as the lateral coalescence of GaN proceeds among the nanorods. With the entire layer structure completed at Stage 4, the signal peak shifts back to 565.1 cm −1 , showing a mitigated biaxial stress. The partial release of the tensile strain after the MQW growth is due to the com-pensation brought by the thin InGaN wells, which are compres-sively stressed on the GaN barriers. The biaxial stress, σ xx , can be estimated by the equation: [ 24 ] Δ ω = Kσ xx cm −1 GPa −1 , where Δ ω is the wavenumber shift of the E 2 -high peak from the stress-free peak and K is the linear stress coeffi cient. Using K = 4.3 reported by Tripathy et al., [ 25 ] σ xx of the epitaxial structure at Stage 4 is around 0.88 GPa. This estimated stress value is com-pared with that attained on a conventional planar Al-containing buffer layer, whose phonon peaks are shown in Figure 2 c. The fi gure presents the Raman spectrum of the epilayers (capping layer/MQW/n-GaN) identical to that in Figure 2 b, except that the buffer/substrate of AlGaN-superlattice/AlN/(111) Si is adopted as the supporting base. [ 26 ] Estimated with the same equation, the E 2 -high peak at 563.1 cm −1 in Figure 2 c gives the biaxial stress of 1.35 GPa, which is 53% larger than that achieved on the nanostructured buffer. It is noticed that the spectra meas-ured on the nanorod buffer exhibit increased secondary peaks and background noises, which can be attributed to the surface optical phonon modes brought by the scattering effect on the nanopyramid surface (Figure 2 d,e). [ 27 ] The results in Figure 2 b,c clearly indicate that the nano-heteroepitaxy employing ZnO

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Figure 1. Bird’s-eye view SEM images of the AlN a) and GaN b) nuclea-tion layers grown at 580 °C on ZnO nanorods/(001) Si substrates. The larger lattice mismatch between AlN and ZnO leads to randomly oriented crystals, whereas the well-aligned embryo pyramids are observed with GaN nucleation layer.

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nanorod buffer can effectively relieve the tensile strain of GaN-based layers grown on Si substrates.

Figure 2 d,e presents the cross-sectional and the top-view SEM images of the MQW structure grown with the tempera-tures described in Figure 2 a. It is seen that the nucleation of GaN among the ZnO nanorods left air voids buried in the buffer region (Figure 2 d). These voids are benefi cial for strain relieving in heteroepitaxial GaN. [ 8,28 ] The porous buffer also shows promise as the potential sacrifi cial layer for facile detach-ment of the light-absorbing Si substrate, facilitating the fabri-cation of thin-fi lm LEDs. [ 29 ] As displayed in the cross-sectional image, the GaN grown on the ZnO nanostructure exhibits rigid pyramidal feature with semipolar {1011} facets, which form a tilt angle of around 62° with the (0001) c -plane surface. The tilted facets are also reported with the GaN/ZnO core–shell structure grown with top-down methods. [ 12,30 ] It is worthwhile to mention that morphology of the GaN epilayer is determined by the lateral growth rate ( R L ) relative to the vertical growth rate ( R V ). Specifi cally, the ratio of R L / R V can be controlled by sub-strate temperature and reactor pressure in MOCVD growth. [ 31 ] Faceted surface texture is generally attained with low substrate temperature or high reactor pressure, either of which reduces atomic migration on the epitaxial surface. In our case, the semipolar {1011} nanopyramids can be reproduced with the temperature of 1000 °C and the pressure of 200 mbar.

Note that the heights of these nanopyramids are mostly within 300 nm. Surface roughness in this scale can be over-come by wafer bonding or electroplating processes in the fab-rication of thin-fi lm LEDs, during which the thickness of the bonding or electroplating metal layer is typically beyond a few

micrometers. [ 29,32 ] For the MQW grown on the semipolar facets, light extraction effi ciency is expected to be enhanced through the increased emission area and the reduced total internal refl ec-tion. [ 33 ] More importantly, the {1011} facets are of two proper-ties that are strongly desired by InGaN MQW: i) According to the theoretical analyses on surface orientation dependence of the polarization discontinuity at the InGaN/GaN interface, [ 34,35 ] the 62° tilt angle of {1011} plane is close to the one corresponding to the polarization crossover from positive to negative values. In other words, the total (spontaneous plus piezoelectric) polariza-tion discontinuity at the interface of InGaN/GaN can be signifi -cantly alleviated on the {1011} surface. InGaN MQW grown on this nearly polarization-free plane should exhibit much boosted radiative recombination effi ciencies owing to the increased overlap between electrons’ and holes’ wave functions. [ 14 ] ii) In addition to the greatly reduced polarization, the {1011} facets also render the indium incorporation effi ciency larger than those on polar (0001) and nonpolar {1011} surfaces. It has been found that the bonding sites on {1011} surface are ener-getically favorable for the incoming indium atoms because of the smaller chemical potential required, comparing to the cases on (0001) and {1011} surfaces. [ 36,37 ] For the growth of high-in-content MQW, which is vitally important for green emitters, the high indium incorporation effi ciency of a {1011} plane gains more budget in the trade-off between long emission wavelength and crystal qualities; that is to say, the growth temperature of MQW can be increased for superior qualities without sacrifi cing indium incorporation. [ 38 ]

It should be mentioned that the growth homogeneity of this bottom-up nano-heteroepitaxy technique hinges on the

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Figure 2. a) Substrate temperature plotted as a function of time during the growth of the MQW structure on ZnO nanorods/(001) Si. b) Raman spectra collected on the epilayers at Stages 1, 2, 3, and 4 designated in (a). c) Raman spectra of the MQW structure grown with the same parameters as in (b) on planar AlGaN-superlattice/AlN/(111) Si. d) Cross-sectional and e) top-view SEM images of the GaN nanopyramids grown on ZnO-nanorods/(001) Si. f) Photograph showing the 2 inch bare (001) Si substrate, the ZnO nanorods, and the GaN nanopyramids. The top row presents the sche-matics corresponding to each stage.

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thickness uniformity of the ZnO seed layer, which is critical to the morphology of the nanorods. [ 17 ] As the thickness of the seed layer is precisely controlled by the sputtering process, wafer-scale uniformity of the ZnO-nanorod buffer and the semipolar GaN epilayer can be achieved (Figure 2 f). In the fi gure, one can see clear Newton’s rings on the ZnO-nanorod buffer, which are created by the interference between the light beams refl ected at the fl at nanorod tops and the ZnO/Si interface. The dark appearance of the GaN epilayer seen in the fi gure is due to the low surface refl ectance on the nanopyramids, which induce strong light scattering/trapping among the inclined facets. Nanopyramids with slightly varied heights, as displayed in Figure 2 d, should lead to different surface refl ectance, thereby changing the brightness in some areas.

In order to acquire structural information of the nano-hetero epitaxial GaN, systematic characterization with transmis-sion electron microscopy (TEM) was performed ( Figure 3 a). The image displays the sections from the substrate to the high-temperature GaN (HT-GaN), whose boundaries are revealed by energy-dispersive X-ray spectroscopy (EDX) elemental map-ping of Si, Zn, and Ga (Figure 3 b–e). In Figure 3 b, it is found that the density of Ga is decreased in some areas of HT-GaN, particularly near the top. The result is mainly attributed to the sharpened apexes of the nanopyramids as well as the variation of space fi lling with GaN in the TEM sample. The dark area of low-temperature GaN (LT-GaN), seen in Figure 3 a, indicates

that most of the defects are confi ned in this region, and the defect densities are markedly reduced in the region of HT-GaN. The result is caused by the growth transformation from kinetic mode to thermodynamic mode when the substrate tempera-ture is increased from 580 to 1000 °C. The increased atomic diffusion length at 1000 °C allows Ga and N ions to be bonded at low-energy lattice sites and thus reduces the possibility of defect formation. [ 38 ] In Figure 3 e, it can be seen that the distri-bution of Ga is overlapped with the Zn region, indicating the ZnO–GaN core–shell structure formed at the temperature of 580 °C. The core–shell structure prevents the vulnerable ZnO from the bombardment of the hot H + ions during the growth of HT-GaN, preserving the epilayers built on the nanorod buffer.

To explore further details of the nanostructure, a high resolu-tion (HR) TEM image of the interface between LT- and HT-GaN is presented in Figure 4 a, showing the atomic stacking order along [0001], which agrees with the result of XRD characteri-zation (Figure S4, Supporting Information). The image reveals prevalent planar defects parallel to the (0001) basal plane within the LT-GaN layer. In this region, the corresponding streaks seen in the diffractogram from fast Fourier transfor-mation (FFT) and the diffraction pattern without extra spots (Figure S5, Supporting Information) suggest that the planar defects are stacking faults (SF). Contrarily, in the region away from the interface (dashed line in Figure 4 a) and into the

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Figure 3. a) Bright-fi eld TEM image and b) the corresponding EDX elemental mapping of the GaN nanopyramids grown on ZnO-nanorods/(001) Si. EDX elemental mapping of Si c); Zn d); and Ga e), indicating the ZnO–GaN core–shell structure formed in the nanorod buffer.

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HT-GaN, the crystallinity improves signifi cantly, as evidenced by the HR-TEM image and the corresponding diffractogram. The presence of the planar defects is suggested to dominate the strain relax due to the mismatch between ZnO and GaN, and results in the superior crystal quality of HT-GaN.

It is noted that the defect type observed in this material system is distinct from those commonly found in GaN, where the strain due to lattice mismatch is mostly released through threading dislocations. [ 9 ] The high density of SF indicates that, when the lateral diffusion is promoted by the high V/III ratio during the growth of LT-GaN, the lateral growth front sur-face would undergo coalescence by merging the ZnO–GaN core–shell nanorods with different in-plane crystallographic orientations, which tends to be accompanied by the formation of the planar defects. As atomic motion in the vertical growth direction is enhanced by the decreased V/III ratio (2400) in the HT-GaN, defect density can be signifi cantly reduced via

atomic rearrangement or a self-healing process. Similar obser-vation has been reported with the GaN grown on Si nanopillar arrays. [ 7 ]

Figure 4 b demonstrates the TEM image of the semi-polar {1011} InGaN/GaN MQW grown on the nanopyramid GaN surface, showing virtually no crystal defects. The spec-trum and a photograph of the photoluminescence (PL) meas-ured with semipolar MQW can be seen in Figure S6 in the Supporting Information. In the TEM image, the dark regions at the two vertical edges are due to the overlapping with adja-cent nanopyramids. A close observation leads to the fi nding that the barrier widths become thinner at the base region than those at the tip. The difference can be attributed to the decreased chemical reactants diffusing toward the lower part of nanopyramids, leading to the slower growth rate in comparison with that at the tip. [ 39 ] Although the crystal quality is dramati-cally improved in HT-GaN, planar defects similar to those pre-vailing in LT-GaN are occasionally observed in the MQW region (Figure S7, Supporting Information). The result suggests that lateral coalescence of GaN is continued at the temperature of 1000 °C despite the reduced V/III ratio and SF-like defects can still occur as misaligned crystal planes coalesce. The problem should be eased by optimizing the growth rates in vertical and lateral directions through properly selected V/III ratios, sub-strate temperature and reactor pressure for HT-GaN. [ 20,31 ]

In summary, wafer-scale GaN nanopyramids with semi-polar {1011} facets were grown on (001) Si substrates using ZnO-nanorod buffer layer. It is demonstrated that the Al-free and self-assembly nanostructure leads to much reduced tensile strain in the epilayer. The semipolar {1011} facets are not only nearly polarization free but also expected to offer the indium incorporation effi ciency larger than those on polar (0001) and nonpolar {1011} planes. TEM analyses reveal that planar SF-like defects are mostly found in the low-temperature nuclea-tion layer, and defect density is considerably decreased in the HT-GaN region. The bottom-up nano-heteroepitaxy method presented here promises the potential of high-brightness, low-cost III-nitride emitters whose fabrication is compatible with CMOS technology.

Experimental Section ZnO Nanorod Preparation : The synthesis of ZnO nanorods was

initiated from the 80 nm ZnO seed layer deposited by sputtering. After annealing the seed layer in air at 1000 °C for 1 h, the nanorods were grown with the precursors of zinc nitrate hexahydrate (Zn(NO 3 ) 2 ·6H 2 O) and hexamethylenetetramine (C 6 H 12 N 4 ). The hydrothermal process was carried out on the hot plate at 90 °C for 1 h.

GaN Growth : GaN layers were grown by MOCVD (AIXTRON 200/4 RF). Ammonia (NH 3 ), trimethylgallium (for Si-doped n-GaN and the undoped cap layer), triethylgallium (for MQW), and trimethylindium were used as the precursors. To prevent ZnO from chemical decomposition, the entire growth was carried out in N 2 ambience. The MQW structure consists of six repeat of intentionally undoped InGaN/GaN layers. V/III ratios for the 580 °C undoped GaN, 1000 °C n-GaN, 785 °C InGaN wells, and 785 °C GaN barriers were 4266, 2400, 9028, and 11 416, respectively. The reactor pressure was kept at 200 mbar.

Microscopy : SEM images were recorded with fi eld-emission Hitachi S-4300 at the acceleration voltage of 10 kV. TEM samples were prepared by focus ion beam using Ga ions at 30 kV. The analyses of TEM were

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Figure 4. a) HR-TEM image of the HT-GaN/LT-GaN interface and the corresponding FFT diffractograms, showing the SF-like planar defects in LT-GaN. b) Cross-sectional TEM image of the six-repeat semi-polar {10 11} InGaN/GaN MQW structure.

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performed with a JEOL 2100F system at 200 kV, equipped with an electron energy loss spectrometer and an EDX detector.

Spectrum : Raman spectra were collected using a confocal system with the excitation laser wavelength of 473 nm and the spot size of ≈0.5 µm. The intensity peaks at 517 and 520 cm −1 for (001) Si and (111) Si, respectively, were used as the reference for wave number calibration. PL spectrum was excited by a 325 nm He–Cd laser.

Supporting Information Supporting Information is available from the Wiley Online Library or from the author.

Acknowledgements This research was supported in part by Ministry of Science Technology Grants NSC 100-2218-E-008-015, NSC 101-2221-E-008-044, NSC 102-2221-E-008-074, and MOST 103-2221-E-008-046.

Received: March 31, 2015 Revised: May 19, 2015

Published online: July 15, 2015

[1] F. A. Ponce , D. P. Bour , Nature 1997 , 386 , 351 . [2] C.-C. Sun , Y.-Y. Chang , T.-H. Yang , T.-Y. Chung , C.-C. Chen ,

T.-X. Lee , D.-R. Li , C.-Y. Lu , Z.-Y. Ting , B. Glorieux , Y.-C. Chen , K.-Y. Lai , C.-Y. Liu , J. Solid State Lighting 2015 , 1 , 19 .

[3] A. Krost , A. Dadgar , Phys. Status Solidi A 2002 , 194 , 361 . [4] S. Luryi , E. Suhir , Appl. Phys. Lett. 1986 , 49 , 140 . [5] D. Zubia , S. D. Hersee , J. Appl. Phys. 1999 , 85 , 6492 . [6] R. Chen , T. T. D. Tran , K. W. Ng , W. S. Ko , L. C. Chuang ,

F. G. Sedgwick , C. Chang-Hasnain , Nat. Photonics 2011 , 5 , 170 . [7] S. D. Hersee , X. Y. Sun , X. Wang , M. N. Fairchild , J. Liang , J. Xu , J.

Appl. Phys. 2005 , 97 , 124308 . [8] G. T. Chen , J. I. Chyi , C. H. Chan , C. H. Hou , C. C. Chen ,

M. N. Chang , Appl. Phys. Lett. 2007 , 91 , 261910 . [9] D. Zhu , D. J. Wallis , C. J. Humphreys , Rep. Prog. Phys. 2013 , 76 , 106501 .

[10] G. B. Stringfellow , Organometallic Vapor-Phase Epitaxy , Academic Press , San Diego, CA, USA 1999 .

[11] N. Li , S. J. Wang , E. H. Park , Z. C. Feng , H. L. Tsai , J. R. Yang , I. Ferguson , J. Cryst. Growth 2009 , 311 , 4628 .

[12] J. Yoo , S. T. Picraux , G.-C. Yi , Proc. SPIE 2011 , 8106 , 81060N . [13] C. A. Lin , K. Y. Lai , W. C. Lien , J. H. He , Nanoscale 2012 , 4 , 6520 . [14] P. Waltereit , O. Brandt , A. Trampert , H. T. Grahn , J. Menniger ,

M. Ramsteiner , M. Reiche , K. H. Ploog , Nature 2000 , 406 , 865 . [15] Z. L. Wang , J. Phys.: Condens. Matter 2004 , 16 , R829 .

[16] R.-J. Chung , Z.-C. Lin , C.-A. Lin , K.-Y. Lai , Thin Solid Films 2014 , 570 , 504 .

[17] S. W. Chen , J. M. Wu , Acta Mater. 2011 , 59 , 841 . [18] M. E. Levinshtein , S. L. Rumyantsev , M. S. Shur , Properties of

Advanced Semiconductor Materials GaN, AlN, InN, BN, SiC, SiGe , John Wiley & Sons, Inc ., New York 2001 .

[19] M. Dauelsberg , D. Brien , H. Rauf , F. Reiher , J. Baumgartl , O. Häberlen , A. S. Segal , A. V. Lobanova , E. V. Yakovlev , R. A. Talalaev , J. Cryst. Growth 2014 , 393 , 103 .

[20] D. Kapolnek , S. Keller , R. Vetury , R. D. Underwood , P. Kozodoy , S. P. Den Baars , U. K. Mishra , Appl. Phys. Lett. 1997 , 71 , 1204 .

[21] T. Kozawa , T. Kachi , H. Kano , H. Nagase , N. Koide , K. Manabe , J. Appl. Phys. 1995 , 77 , 4389 .

[22] K. W. J. Wong , M. R. Field , J. Z. Ou , K. Latham , M. J. S. Spencer , I. Yarovsky , K. Kalantar-Zadeh , Nanotechnology 2012 , 23 , 015705 .

[23] B. Sieber , H. Liu , G. Piret , J. Laureyns , P. Roussel , B. Gelloz , S. Szunerits , R. Boukherroub , J. Phys. Chem. C 2009 , 113 , 13643 .

[24] J.-M. Wagner , F. Bechstedt , Appl. Phys. Lett. 2000 , 77 , 346 . [25] S. Tripathy , S. J. Chua , P. Chen , Z. L. Miao , J. Appl. Phys. 2002 , 92 ,

3503 . [26] C. Y. Liu , C. C. Lai , J. H. Liao , L. C. Cheng , H. H. Liu , C. C. Chang ,

G. Y. Lee , J.-I. Chyi , L. K. Yeh , J. H. He , T. Y. Chung , L. C. Huang , K. Y. Lai , Sol. Energy Mater. Sol. Cells 2013 , 117 , 54 .

[27] S.-E. Wu , S. Dhara , T.-H. Hsueh , Y.-F. Lai , C.-Y. Wang , C.-P. Liu , J. Raman Spectrosc. 2009 , 40 , 2044 .

[28] P. Frajtag , N. A. El-Masry , N. Nepal , S. M. Bedair , Appl. Phys. Lett. 2011 , 98 , 023115 .

[29] K. M. Lau , K. M. Wong , X. Zou , P. Chen , Opt. Express 2011 , 19 , A956 .

[30] C.-H. Lee , Y. J. Hong , Y.-J. Kim , J. Yoo , H. Baek , S.-R. Jeon , S.-J. Lee , G.-C. Yi , IEEE J. Sel. Top. Quant. Electron. 2011 , 17 , 966 .

[31] H. Miyake , A. Motogaito , K. Hiramatsu , Jpn. J. Appl. Phys. 1999 , 38 , L1000 .

[32] W. S. Wong , T. Sands , N. W. Cheung , M. Kneissl , D. P. Bour , P. Mei , L. T. Romano , N. M. Johnson , Appl. Phys. Lett. 2000 , 77 , 2822 .

[33] T. Fujii , Y. Gao , R. Sharma , E. L. Hu , S. P. DenBaars , S. Nakamura , Appl. Phys. Lett. 2004 , 84 , 855 .

[34] A. Strittmatter , J. E. Northrup , N. M. Johnson , M. V. Kisin , P. Spiberg , H. El-Ghoroury , A. Usikov , A. Syrkin , Phys. Status Solidi B 2011 , 248 , 561 .

[35] J. E. Northrup , Appl. Phys. Lett. 2009 , 95 , 133107 . [36] J. E. Northrup , L. T. Romano , J. Neugebauer , Appl. Phys. Lett. 1999 ,

74 , 2319 . [37] J. E. Northrup , J. Neugebauer , Phys. Rev. B 1999 , 60 , R8473 . [38] H. Morkoc , Handbook of Nitride Semiconductors and Devices ,

Wiley-VCH , Weinheim, Germany 2009 . [39] J.-R. Chang , S.-P. Chang , Y.-J. Li , Y.-J. Cheng , K.-P. Sou , J.-K. Huang ,

H.-C. Kuo , C.-Y. Chang , Appl. Phys. Lett. 2012 , 100 , 261103 .