InN hexagonal thin wurtzite disks were grown on γ-LiAlO2 by plasma-assisted molecular-beam epitaxy at low temperature (470oC). The ( 000 1 ¯ ) InN thin disk was established with the capture of N atoms by the β ¯ -dangling bonds of most-outside In atoms, and then the lateral over-growth of the In atoms were caught by the β ¯ -dangling bonds of the N atoms. From the analyses of high-resolution transmission electron microscopy, the lateral over-grown width was extended to three unit cells at [ 1 1 ¯ 00 ]InN direction for a unit step-layer, resulting in an oblique surface with 73o off c-axis.

III-Nitride semiconductor compounds were extensively investigated and applied to optoelectronic and spintronic devices, such as solar cells, sensors, high electron mobility transistors (HEMTs) and light emitting diodes (LEDs) in the past decades.1–5 Wide direct band-gap gallium nitride (GaN) and aluminum nitride (AlN) compounds, with energy gaps covering the spectrum from blue to ultraviolet range, are the dominant materials for solid state lighting devices and well-studied to date. Because of the improvement of InN films grown by molecular-beam epitaxy (MBE), the direct band-gap indium nitride (InN) compound has been demonstrated with the value of 0.64 eV rather than 1.9 eV.6,7 This important finding indicated that the full-color spectrum (red, green and blue) can be achieved by devices using the group of III-nitride compounds without any phosphor. Besides InN is a high potential material in optoelectronic applications due to its outstanding properties; such as, the smallest effective mass, the highest peak and saturation electron drift velocity, and the largest mobility in the nitride semiconductors.8,9 Currently, Liang et al. reported the growth of synthesized InN nanowires on selective-area of gold-patterned Si (100) substrates by a vapor-liquid-solid mechanism.10 Stoica et al. evaluated the optimum growth conditions for uniform and high-quality InN nanowires on Si (111) by plasma-assisted MBE (PA-MBE).11 However, the InN nanowires exhibited taper morphology with a large variation in the aspect ratio. Chang et al. grew non-tapered and nearly homogeneous InN nanowires on Si (111) with a constant wire diameter of ∼71 nm, using the PA-MBE with an in situ deposited thin (∼ 0.5 nm) In seeding layer. By varying the growth time, they also achieved non-tapered InN nanowires with lengths in the range of 0.1-5 μm.12 In our previous works, we have grown the high-quality self-assembling c-plane GaN ( 000 1 ¯ ) hexagonal microdisk with a diameter of 4 μm on LiAlO2 (LAO) substrates, which can be adjusted to optimize the quantum effect for nano-device applications.5 Besides, we developed a back process to fabricate an electrical contact for the GaN hexagonal microdisk on a transparent p-type GaN template.13 Therefore the InN microdisk provides an opportunity to fabricate the InGaN/GaN microdisk quantum well for the application of full-color micron LED without the sapphire substrate, which has a large lattice mismatch with InN and is mostly used for the bulk GaN-based quantum wells in commercial LEDs.1,2 In this paper, we will show the growth of InN ( 000 1 ¯ ) hexagonal thin wurtzite microdisk on LAO substrate by PA-MBE.

The sample was grown on high-quality 1x1 cm2 LAO (100) substrate by using low-temperature PA-MBE system (Veeco Applied-GEN 930) with standard effusion cells for In- and Ga-evaporation and an rf-plasma cell with 300 W for N2-plasma source. The LAO substrate was cut from the crystal ingot, which was fabricated by the traditional Czochralski pulling technique. Before mounting on a holder, the LAO substrate was cleaned with acetone (5 min), isopropanol (5 min) and de-ionized water in a while, and then dried with nitrogen gas immediately. After the chemical cleaning, a thermal treatment was introduced to the LAO substrate in MBE chamber before epitaxial growth. The LAO substrate was out-gassed at 850oC for 10 min. The temperature was defined by a thermal couple equipped at the backside of the substrate. Thereafter, the substrate temperature was decreased down to growth temperatures. The Ga wetting layer was performed on the LAO substrate for 5 min at 800oC, and then the two-step method (i.e. two different N/Ga flux ratio from 39.9 to 55.9, for 25 min and 30 min, at 800oC and 850oC, respectively) was used to fabricate GaN epi-film. The flux ratio was represented by Beam Equivalent Pressure (BEP) of evaporative III-group sources from standard effusion cell against that of N2 source from rf-plasma cell.14 Thereafter, the substrate temperature was ramped down to InGaN growth temperature at 650oC with a flux ratio In/Ga = 1.26 and the duration time for the growth of InGaN buffer layer was 30 min. Finally, the rf-plasma power was raised to 450 W to perform InN. The substrate temperature was ramped down to 470oC with a flux ratio N/In = 39.3 and the duration time for the growth of InN was 60 min.

The surface morphology of the sample was evaluated by the field emission scanning electron microscopy (FE-SEM, SII-3050). Figs. 1(a) and 1(b) showed the SEM images with a tilted angle and a top-view of the sample, respectively. The morphology of the sample exhibited that two-dimensional (2D) M-plane InN film and three-dimensional (3D) c-plane InN hexagonal thin disks were grown on the LAO substrate. Figs. 1(c) and 1(d) showed the SEM images with a tilted angle and a top-view of the InN thin disk shown in the center of Fig. 1(b), and the diameter of the InN disk was 0.74 μm. The micrographic images of the sample showed that the 2D M-plane InN film were developed along with the lateral orientation [ 11 2 ¯ 0 ]InN // [001]LAO, while the 3D c-plane InN hexagonal thin disks and nanopillars were grown atop an anionic hexagonal basal plane of LAO. The two-orientation growth of GaN nanopillars on LAO substrate has been reported in our previous papers.15,16 In this paper, we applied the two-orientation growth to the M-plane InN film and 3D c-plane InN hexagonal thin disks on LAO substrate with GaN and InGaN buffer layers. The crystal structure of the sample was characterized by the high-resolution X-ray diffraction (XRD) (Bede D1) measurements and showed in Fig. 2. The indium content of M-plane InxGa1-xN was evaluated by Vegards’ equation,17 giving the 2θ peak of X-ray diffraction pattern (31.69 ± 0.01)o to In0.20Ga0.80N. The 2θ peaks at (29.07 ± 0.01)o, (31.31 ± 0.01)o, (32.29 ± 0.01)o and (34.69 ± 0.01)o were corresponding to the X-ray diffraction patterns from M-plane InN ( 1 1 ¯ 00 ), c-plane InN ( 000 2 ¯ ), M-plane GaN ( 1 1 ¯ 00 ) and LAO (100), respectively. These peak positions for the X-ray diffraction patterns were obtained by software Quick Graph (Version 2.0) with the Asymmetric Double Sigmoidal linear curve fitting, and matched with those data of the standard wurtzite structure bulk InN (JCPDS file No. 50-1239). The results were reconfirmed by software OriginPro 8.0 with a Gaussian-function curve fitting. As compared to Bragg’s law (2dsinθ = nλ) and Cu Kα1 wavelength λ = 0.1540562 nm, the d-spacing between { 000 2 ¯ } planes of InN was evaluated to be d0002= 0.28216 nm. Compared with the value on JCPDS file, d0002= 0.28528 nm, the difference between wurtzite InN microdisk and bulk InN is 1.09%, indicating that the lattice constant of InN microdisk is smaller than that of bulk InN.

FIG. 1.

(a) The SEM image with a tilted angle of the sample, the scale bar is 1 μm. (b) The top-view SEM image of the sample, the scale bar is 1 μm. (c) Enlarged SEM image with a tilted angle of InN hexagonal thin disk, the scale bar is 0.3 μm. (d) Enlarged SEM image of InN hexagonal thin disk, the scale bar is 0.3 μm.

FIG. 1.

(a) The SEM image with a tilted angle of the sample, the scale bar is 1 μm. (b) The top-view SEM image of the sample, the scale bar is 1 μm. (c) Enlarged SEM image with a tilted angle of InN hexagonal thin disk, the scale bar is 0.3 μm. (d) Enlarged SEM image of InN hexagonal thin disk, the scale bar is 0.3 μm.

Close modal
FIG. 2.

X-ray 2Theta-Omega scan of the sample.

FIG. 2.

X-ray 2Theta-Omega scan of the sample.

Close modal

The microstructure of the sample was analyzed by field emission transmission electron microscope (FE-TEM) (Phillips, model Tecnai F-20) with an electron voltage of 200 kV. The cross-sectional TEM specimen of the sample was prepared by a dual-beam FIB system (Seiko Inc., SII-3050), on the cleavage plane along [ 1 1 ¯ 00 ] direction of the c-plane InN hexagonal thin disk. The FIB was performed with accelerated voltage of 30 kV to cut the samples roughly and then refined the specimen further by accelerated voltage of 5 kV. Fig. 3(a) showed the bright field image with [ 11 2 ¯ 0 ]InN // [001]LAO zone axis. It clearly exhibited that InN was well-formed on InGaN buffer layer and the InGaN buffer layer was well-established on GaN epi-layer as shown at the areas of DP4 – DP6. The thicknesses of M-plane InN, M-plane InGaN and M-plane GaN were measured to be about 265 nm, 51 nm and 137 nm, respectively. At the areas of DP3, a c-plane InN hexagonal thin disk was established. The height for the c-plane InN hexagonal thin disk from neck to top was about 188 nm. The selective area diffraction (SAD) patterns at the areas located at DP1 – DP6 were shown in Figs. 3(b)3(g). Fig. 3(b) clearly showed one rectangular (white) and one hexagonal (blue) diffraction patterns overlapped at the location of DP1, indicating that the c-plane wurtzite GaN and M-plane wurtzite GaN were initially nucleated atop the LAO substrate. We also found that the intensity of the hexagonal diffraction spots were brighter than the rectangular ones, indicating that the growth condition was suitable to 2D M-plane wurtzite GaN film at the beginning of nucleation. Fig. 3(c) simply showed one single rectangular diffraction pattern (red) at the location of DP3, indicating that the hexagonal thin disk was uniquely formed by the c-plane wurtzite InN. The d-spacing between { 000 1 ¯ } planes of InN hexagonal thin disk was measured to be dc = 0.5687 nm and the d-spacing between { 1 1 ¯ 00 } planes of InN hexagonal thin disk was dM = 0.3025 nm. Compared with the values on JCPDS file No. 50-1239 which are 0.5703 nm and 0.30647 nm, respectively, the difference between wurtzite InN microdisk and bulk InN for dc and dM are 0.28% and 1.24%, respectively, revealing that the lattice constant of wurtzite InN microdisk is smaller than that of bulk InN. The result is consistent with the X-ray diffraction patterns. At the neck area of the disk (location of DP2), the SAD patterns showed the overlapping diagram of two rectangles and two hexagons in Fig. 3(d), indicating that an M-plane InN (yellow hexagon) was formed in addition to the c-plane GaN (white rectangle), c-plane InN (red rectangle) and M-plane GaN (blue hexagon) at the neck area. The M-plane InN (yellow hexagon) can be checked by the SAD patterns performed at the locations of DP4 – DP6, as shown in Figs. 3(e)3(g). At the location of DP6, only one single hexagonal pattern (yellow) was detected for the M-plane wurtzite InN, as shown in Fig. 3(e). At the location of DP4, it showed one hexagonal (blue) diffraction pattern for M-plane GaN in Fig. 3(g), with some unclear spots which were produced from the diffraction of LAO substrate. At the location of DP5, two hexagonal patterns were overlapped in Fig. 3(f), indicating the formation of both M-plane GaN (blue hexagon) and InN (yellow hexagon) occurred. These two hexagons are identical to those shown in Fig. 3(d), indicating that the M-plane wurtzite InN and M-plane wurtzite GaN were grown in the same crystalline direction. From the analyses of SAD patterns, we found that the c-plane wurtzite nanocrystal was embedded between M-plane wurtzite GaN areas at the beginning of nucleation when GaN was grown on LAO substrate (e.g., at the location of DP1). The detailed microstructure of the neck area was investigated by high-resolution transmission electron microscopy.

FIG. 3.

TEM analyses of the InN hexagonal thin disk: (a) the bright field image with [ 11 2 ¯ 0 ]InN // [001]LAO zone axis. The selective area diffraction patterns taken at the points shown in (a) are presented in [(b) – (g)]. The scale bars are 2 (1/nm).

FIG. 3.

TEM analyses of the InN hexagonal thin disk: (a) the bright field image with [ 11 2 ¯ 0 ]InN // [001]LAO zone axis. The selective area diffraction patterns taken at the points shown in (a) are presented in [(b) – (g)]. The scale bars are 2 (1/nm).

Close modal

The high-resolution TEM images with the beam direction of [ 11 2 ¯ 0 ]InN // [001]LAO were performed at the areas HR1 – HR6 of the sample, as shown in Figs. 4(a). The interfaces between c-plane and M-plane wurtzite GaN, InGaN, and InN were shown in Fig. 4(b) and 4(c). The staking faults were found at the boundary between c-plane and M-plane GaN which release the dislocations between the misfit c-plane and M-plane wurtzite structures of GaN and InGaN. The c-plane wurtzite structure was followed up to the neck area (e.g., HR3) and formed a uniform c-plane InGaN pyramid-shaped structure, as shown in Fig. 4(d). Outside the pyramid-shaped structure, the wave-shaped structures were produced by the staking faults between the misfit c-plane wurtzite structures of InGaN and InN. The uniform c-plane InGaN pyramids can be observed in Figs. 1(a) and 1(b), which emitted a higher luminesce intensity like a quantum-dot structure. The wave-shaped structures became uniform at the area HR4, as shown in Fig. 4(e). The c-plane wurtzite structure was followed further to form the InN hexagonal microdisk structures. The symmetric hexagonal shape reveals the high-quality crystalline structure of the InN microdisk, as shown in Fig. 4(f). It is noted that the thinness of the hexagonal microdisk structure leads to an angle of 73o off the c-axis in Fig. 4(a), which is much greater than the angle of GaN microdisk,5 obtained from the lateral over-growth along the ( 1 1 ¯ 00 ) direction for each unit step-layer by one dM-spacing, θ = tan−1(dM/dc) = 28o. To establish the growth mechanism of the thin InN hexagonal microdisk, we demonstrated a ball-stick model for the self-assembled thin InN microdisk. The ball-stick model for the standard wurtzite InN (JCPDS file No. 50-1239) with a = b = 0.3537 nm, c = 0.5703 nm, and u = α ¯ / c = 3 / 8 was used to simulate the c-plane InN hexagonal microdisk in Fig. 5(a), where blue balls represented In atoms and red balls represented N atoms. In our previous paper, we showed that the GaN ( 000 1 ¯ ) microdisk with a tilted angle of 28o was established with the capture of N atoms by the β ¯ -dangling bonds of the most-outside Ga atoms for each unit step-layer during the GaN lateral over-growth.5 In the case of InN microdisk, when the growth temperature lowered to 470oC, the c-plane InN ( 000 1 ¯ ) hexagonal thin disk was built up with the capture of N atoms by the β ¯ -dangling bonds of the most-outside In atoms and then the lateral over-growth occurred; capture of In atoms by β ¯ -dangling bonds of N atoms, to form the thin microdisk. The lateral over-growth along the ( 1 1 ¯ 00 ) direction was extended to six dM-spacings for each unit step-layer (i.e., dc-spacing), resulting in the angle of 73o off the c-axis. Based on the ball-stick model, the laterally extensive width along the [ 1 1 ¯ 00 ]InN direction per unit step-layer was equal to 3 3 a . The edge was then tilted off the c-axis [ 000 1 ¯ ] direction by the angle, ϕ = tan 1 ( 3 3 a / c ) = 72 . 7 6 o , where 3 3 a is equal to 6dM, as shown in Fig. 5(b). We also calculated the angle from the measured SAD data at the InN hexagonal thin disk in Fig. 3(c), and obtained that the d-spacing between { 000 1 ¯ } planes was dc = 0.5687 nm and the d-spacing between { 1 1 ¯ 00 } planes was dM = 0.3025 nm, resulting in ϕ = tan−1(6dM/dc) = 72.60o, which was in good agreement with the model predicted. The angle between edge and growth direction can be examined directly by the high-resolution TEM image performed at HR6 to be about 73o, as shown in Fig. 5(c). In order to check the angle between edge and growth direction, we made the fast Fourier transform (FFT) and inversed fast Fourier transform (IFFT) from the high-resolution TEM image at HR6. Figs. 5(d) and 5(e) showed the FFT and IFFT patterns of the white-dot square area in Fig. 5(c), which reconfirmed that the oblique angle (ϕ ≈ 73o) of c-plane InN hexagonal thin disk was formed by the lateral over-growth of wurtzite structure.

FIG. 4.

TEM analyses of the InN hexagonal thin disk: (a) the bright field image with [ 11 2 ¯ 0 ]InN // [001]LAO zone axis. The high-resolution TEM images taken at the points shown in (a) are presented in [(b) – (f)]. The scale bars are 2 nm.

FIG. 4.

TEM analyses of the InN hexagonal thin disk: (a) the bright field image with [ 11 2 ¯ 0 ]InN // [001]LAO zone axis. The high-resolution TEM images taken at the points shown in (a) are presented in [(b) – (f)]. The scale bars are 2 nm.

Close modal
FIG. 5.

The ball-stick model for InN epilayer: (a) the chemical bonds of ( 000 1 ¯ ) surface, (b) the hexagonal thin disk. (c) The high-resolution TEM image taken at the point shown in 4(a), the scale bar is 2 nm. (d) The fast Fourier transform pattern from the white-dot square area in (c). (e) The inversed fast Fourier transform pattern from (d).

FIG. 5.

The ball-stick model for InN epilayer: (a) the chemical bonds of ( 000 1 ¯ ) surface, (b) the hexagonal thin disk. (c) The high-resolution TEM image taken at the point shown in 4(a), the scale bar is 2 nm. (d) The fast Fourier transform pattern from the white-dot square area in (c). (e) The inversed fast Fourier transform pattern from (d).

Close modal

We have grown InN hexagonal thin wurtzite disks on γ-LiAlO2 by plasma-assisted molecular-beam epitaxy. From the surface morphology and microstructure analyses, we found that c-plane wurtzite was established at the nucleation of GaN on LAO substrate and c-plane InN hexagonal thin disks were built up at low temperature (470oC) after the InGaN buffer layer. The c-plane InN ( 000 1 ¯ ) hexagonal thin disk was produced with the capture of N atoms by the β ¯ -dangling bonds of the most-outside In atoms, and then laterally over-grown along [ 1 1 ¯ 00 ] direction by 6 dM-spacings for a unit step-layer.

The authors would like to thank, Y. C. Lin, C. D. Tsai, and S. T. You for their assistance. The project was supported by the Ministry of Science and Technology of Taiwan and the Core Facilities Laboratory for Nanoscience and Nanotechnology in Kaohsiung and Pintung Area.

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