AlN Nanostructures Fabricated on a Vicinal Sapphire - American

Article pubs.acs.org/crystal

AlN Nanostructures Fabricated on a Vicinal Sapphire (0001) Substrate Daeyong Eom, Jinwan Kim, Kyungjae Lee, Minhwan Jeon, Cheon Heo, Jaedo Pyeon, and Okhyun Nam* Advanced Photonics Research Center (APRC)/LED Technology Center, Department of Nano-Optical Engineering, Korea Polytechnic University (KPU), 237, Sangidaehak-ro, Siheung-si, Gyeonggi-do 429-793, Republic of Korea ABSTRACT: We report a novel and facile method for the fabrication of various AlN nanostructures with Al polarity using polarity control and selective etching without a mask or metal catalyst. To investigate the polarity transitions of the AlN layers obtained with different growth parameters, AlN layers were grown by high-temperature metalorganic chemical vapor deposition with varying growth temperatures and trimethylaluminum (TMAl) preflow rates. The growth of Alpolar AlN was clearly supported by a lower growth temperatures and higher TMAl preflow rates. Transmission electron microscopy showed that the threading dislocations (TDs) generated at the AlN−sapphire interface were bent toward the boundary of the N-polar grain because of the threedimensional growth mode of the mixed-polarity AlN layer. Finally, defect-free nanopillars, nanorods, nanofurrows, and nanowalls were fabricated by etching mixed-polarity AlN layers with an aqueous KOH solution.

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INTRODUCTION AlN layers are interesting materials for optical and electronic devices because of their wide band gap energy, high thermal conductivity, and high electrical resistivity.1,2 Although its lattice and thermal expansion coefficients are greatly different from those of AlN layers, c-plane sapphire (α-Al2O3) is commonly employed as a substrate for these layers because of its relatively low cost. However, in previous studies of AlN layers grown on sapphire substrates, the layers exhibited a low crystal quality because of the lattice and thermal expansion coefficient mismatches between the AlN layer and the sapphire, as well as the small surface diffusion length of the Al adatoms.3,4 A high density of threading dislocations (TDs) gives rise to a decrease in the internal quantum efficiency (IQE) of deepultraviolet (DUV) light-emitting diodes (LEDs)5,6 and the output power density of high-electron mobility transistors (HEMTs).7,8 Therefore, there have been several reports of improved methods for the growth of high-quality AlN layers on sapphire substrates for high-efficiency AlN-based applications. Fujimoto et al.9 and Brunner et al.10 reported on the growth of AlN layers by high-temperature metalorganic chemical vapor deposition (HT-MOCVD), which increases the surface diffusion length of the Al adatoms. Several methods have been employed to improve the crystal quality of the AlN layers, including epitaxial lateral overgrowth (ELO),11 patterned sapphire substrate (PSS) methods,12 the use of a nucleation layer,13 prior-to-growth flow schemes,14 and migrationenhanced epitaxy (MEE).15 With the development of these methods, the crystal quality of AlN layers has improved gradually. Another way to fabricate high-efficiency AlN-based devices is to use three-dimensional (3D) nanostructures, which © XXXX American Chemical Society

provide many advantages. These nanostructures can reduce the defect density,16 helping to relieve the strain.17 Moreover, their 3D shape provides an increased active volume of the optical devices with various planar geometries.18 Xiang J. reported a high-performance field-effect transistor (FET) using a onedimensional (1D) hole gas in an undoped Ge/Si core−shell radial nanostructure.19 Nevertheless, applications of AlN nanostructures have seldom been pursued, because of the difficulty of fabrication. The selective area growth (SAG) of AlN and high-Al content AlGaN is difficult because of problems with obtaining such area selectivity.20−22 Moreover, nanostructures fabricated using vapor−liquid−solid (VLS) methods with a catalyst typically exhibit catalyst contamination and many defects.23,24 In this study, we report the fabrication of various AlN nanostructures using a novel method without a mask or metal catalyst. The polarity transition of the AlN layers was investigated by controlling the growth temperature and trimethylaluminum (TMAl) preflow rate. Then, AlN nanostructures with Al polarity were produced by the chemical etching of AlN layers with a mixed polarity.

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EXPERIMENTAL SECTION

The AlN layers were grown on a c-plane (0001) vicinal sapphire substrate with an off-cut angle of 0.2° along the m-axis using HTMOCVD at varying growth temperatures between 1250 and 1350 °C. TMAl and ammonia (NH3) were used as the Al and N sources, Received: November 5, 2014 Revised: January 9, 2015

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DOI: 10.1021/cg5016274 Cryst. Growth Des. XXXX, XXX, XXX−XXX

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Crystal Growth & Design respectively, and the reactor pressure was maintained at 40 mbar during all processes. Prior to AlN growth, a TMAl preflow process was conducted at TMAl flow rates of 0, 170, and 340 sccm for 10 s at each of the growth temperatures. After the TMAl preflow process, the AlN layers were grown without interruption for 30 min in a H2 atmosphere. To investigate the surface morphology of the obtained AlN layers, scanning electron microscopy (SEM) (Hitachi S-4000) measurements were performed. The crystallographic polarity of the AlN layers was examined by transmission electron microscopy (TEM) (JEOL JEM2100F), and a wet etching process using a 2.5 wt % aqueous KOH solution at 80 °C for 10 s. Cross-sectional TEM specimens were prepared using the focused ion beam (FIB) (FEI Nova 600 Nanolab) technique. The surface morphology of the sapphire substrate was examined by atomic force microscopy (AFM) (Nanofocus Inc. nTracer) to identify the relationship between the surface of the vicinal sapphire substrate and the AlN layers.

A−I, which have significant differences in surface morphology. The surface morphology gradually improved as the TMAl preflow rate was increased and the growth temperature was decreased. According to the results of several previous experiments, the crystallographic polarity of the AlN layer should be controlled by the growth temperature of the nucleation layer and the preflow conditions prior to AlN growth.25,26 Therefore, it can be reasonably assumed that these morphological differences are related to the proportion of Al and N polarity in the AlN layers. To investigate the changes in the Al- and N-polarity proportions over the entire AlN layers, all samples were immersed in a 2.5 wt % aqueous KOH solution at 80 °C for 10 s. Figure 2 shows plan-view SEM images of all samples after the etching process. The samples were significantly altered after being etched. Specifically, the Al- and N-polar AlN were separated, because Al-polar AlN has an etching rate that is slower than that of N-polar AlN.27,28 The N-polar AlN is more vulnerable to reaction with OH− than the Al-polar AlN, because the bonding conditions between the {0001} and {0001} planes are different.29,30 This means that the area of residual AlN on the upper surface is Al-polar AlN. Thus, the KOH etching process explicitly revealed that the Al-polar AlN layer became predominant as the TMAl preflow rate increased and the growth temperature decreased. The residual Al-polar columnar structures on the upper surface after the KOH etching process in Figure 2d are known as inversion domains (IDs).31,32 It is clear that the Al-polar AlN domains are located in the middle of the 6-fold faceted N-polar AlN domains.33 Because Al-polar AlN has a growth rate that is higher than that of N-polar AlN,34 the Al-polar IDs formed mixed-polarity AlN layers with a rough surface morphology, even though fully Al-polar AlN layers have a smooth surface morphology. The surface morphology was improved by the coalescence of separate Al-polar IDs as the Alpolarity proportion increased. Figure 3 shows cross-sectional bright-field TEM images of samples D and F with the g vector parallel and perpendicular to the c-axis. In sample D, several inversion domain boundaries (IDBs) were observed (orange arrows) along the (0002) g

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RESULTS AND DISCUSSION Polarity Transition and Behavior of TDs in the AlN Layers. Table 1 shows the relationship between the growth Table 1. Growth Conditions and Corresponding Polarity Results of Grown AlN Layers growth conditions sample

temp (°C)

TMAl preflow rate (sccm)

polarity

A B C D E F G H I

1250 1250 1250 1300 1300 1300 1350 1350 1350

0 170 340 0 170 340 0 170 340

mixed mixed Al mixed mixed Al mixed mixed mixed

condition of the AlN layers and the crystallographic polarity results. Fully Al-polar AlN layers were found in samples C and F, while the others were mixed-polar AlN layers having Al and N polarity. Figure 1 shows plan-view SEM images of samples

Figure 1. SEM plan-view images of morphological changes in samples A−I, grown with different TMAl preflow rates and growth temperatures. B


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Crystal Growth & Design

Figure 2. SEM plan-view images of samples A−I after wet etching in an aqueous KOH solution for 10 s at 80 °C.

Figure 4. (a and b) TEM dark-field images of sample D under the two-beam diffraction condition with g = 1100. (c) Highly magnified MBDF image of the area indicated by the red box in panel a. Figure 3. TEM bright-field images of the same areas of samples D (a and b) and F (c and d) with g = 0002 (a and c) and 1100 (b and d).

To further investigate the behavior of TDs in sample D, darkfield TEM measurements were performed. Figure 4a shows a cross-sectional dark-field TEM image of sample D under a twobeam diffraction condition with g = [1100], and Figure 4b shows a magnified image of the area indicated by the red square in Figure 4a. The orange arrows in Figure 4a show where the edge-type TDs are gathered at the boundaries generated by the coalescence between the separate N-polar AlN domains. The TDs are finally terminated in the grooves between the neighboring N-polar AlN domains, as shown by the white arrows in Figure 4a. It is obvious from Figure 4b that there are the defect-free areas (white solid-line arrows) ∼200 nm from the AlN−sapphire interface, because the edge-type TDs bend sharply to the N-polar AlN boundaries (white dashed-line arrows). To more clearly observe the bending of all types of TDs in the vicinity of Al-polar IDs, multiple-beam dark-field (MBDF) microscopy was also performed, as shown in Figure 4c, which is a magnified image of the area indicated by the red

vector but not along the (1100) g vector. These findings are consistent with previous reports on AlN and GaN.32,35,36 Unlike sample D, sample F contained no IDBs in either the (0002) or (1100) g vector (see panels c and d of Figure 3, respectively). This clearly demonstrates that the generation of IDs is impeded by the TMAl preflow process, and thereby, the AlN layer is changed from N polarity to Al polarity. Here, edgetype dislocations with the Burgers vector b = (1/3) ⟨1120⟩ are visible for the (1100) g vector.36,37 In Figure 3b, although edgetype threading dislocations (TDs) are crowded throughout the interface between the AlN layer and sapphire, the TDs rapidly vanish as the thickness of the AlN layer increases beyond 300 nm (white dashed line). In sample F, which has no IDs, although the TDs diminish slightly as the thickness of the AlN layer increases, most TDs propagate to the surface of the AlN layer (Figure 3d). C


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Figure 5. (a−f) Schematic illustration of the TD bending process during the coalescence of nearby mixed-polarity AlN grains.

the surface and D is the surface diffusion coefficient given by the equation D = D0 exp(−E/kT), where E is the energy barrier for hopping and T is the growth temperature.42,43 If the surface diffusion length is larger than the terrace width of the vicinal sapphire, the adatoms will move from the terrace to the step edge and easily stick to the step edge, because the adatoms are highly stable at the step edge.44 It was previously reported that the migration length of Ga adatoms increased as the growth temperature increased,45 and it was also shown that the growth modes of AlxGa1−xN alloys were altered from growth on the terraces to step-flow growth at the step edges with increasing substrate temperatures.46 Therefore, it can be reasonably assumed that Al adatoms are able to diffuse to the step edge of the vicinal sapphire during a TMAl preflow process at high temperatures, resulting in step-flow-like Al adatom diffusion from the step edge. Consequently, the Al adatoms were dominantly absorbed at the step edge of the vicinal sapphire substrate, and the Al-polar AlN was predominantly grown at the step edge of the vicinal sapphire. In sample I, the TMAl preflow rate was so high that the migration length of the Al adatoms was instead decreased by collisions between Al adatoms on the sapphire surface. Therefore, sample I has a column shape rather than zigzag walls. To identify the relationship between the surface of the vicinal sapphire substrate and the AlN layer in sample H, atomic force microscopy (AFM) was performed. The sapphire was annealed at 1350 °C in ambient H2 for 30 s so that atomic steps and terraces of the sapphire could be clearly observed before the AlN layer was grown. Figure 6 shows an 8 μm × 8 μm AFM scan image of the annealed sapphire substrate (panel a) and an SEM image of sample H after 10 s of wet etching (panel b). The irregular zigzag-shaped step and terrace are clearly visible along the [1010]sap. direction, and the terraces are ∼600 nm wide with a step height of ∼5 nm, as shown in Figure 6a. In contrast, straight steps with a uniform terrace width of 600 nm and a step height of 0.8 nm were visible on the sapphire substrate annealed at 1100 °C in ambient H2 for 30 s (results not shown here). Nearly no roughening or faceting occurred during annealing of the sapphire substrate at 1100 °C, because

box in Figure 4a. The TDs generated by the large lattice mismatch between the AlN and sapphire were clearly bent (yellow arrows) toward the N-polar boundaries (white dashed line) at the middle of each of the two Al-polar IDs. Consequently, the TDs barely exist in the Al-polar IDs. The bending of TDs in mixed-polarity AlN is illustrated schematically in Figure 5a−f. In AlN layers grown at high temperatures, an AlOxN1−x interlayer is formed between the sapphire substrate and the AlN layer as soon as the growth of the AlN layer is initiated on the sapphire substrate, because the AlN layer chemically reacts with the sapphire because of the diffusion of oxygen from the sapphire (Figure 5a−c).25,38 Because the polarity of the AlN layer on the AlOxN1−x layer is inverted to N polarity, N-polar AlN is grown on the AlOxN1−x,25,39 and faceted island-shaped mixed-polarity AlN grows in a three-dimensional (3D) growth mode because of the faster growth rate of Al-polar AlN than of N-polar AlN (Figure 5d).34 During the 3D growth, because TDs tend to bend toward the surfaces of the faceted islands,40,41 bent TDs can gather and meet one another at the boundary of the N-polar AlN domains during the coalescence of the neighboring mixedpolar AlN islands (Figure 5e,f). Consequently, as the thickness of the mixed-polarity AlN layer increases, the regions between N-polar boundaries including the Al-polar IDs become defectfree areas. Growth Mechanism of Mixed-Polarity AlN and Fabrication of AlN Nanostructures. To fabricate AlN nanostructures, it is necessary to understand the different structures of the AlN layers. A comparison of panels a−i of Figure 2 indicates there are two distinct shapes of Al-polar AlN IDs in the mixed-polarity AlN. The first is rod-shaped Al-polar AlN IDs, as shown in Figure 2d. This crystallographic structure is conventional for a mixed-polarity AlN layer, where N-polar AlN surrounds Al-polar IDs. The other shape is periodic zigzag wall-shaped Al-polar IDs, which was clearly observed after the KOH etching process, as shown in Figure 2h. In general, the surface diffusion length of Al adatoms is enhanced at high temperatures. The surface diffusion length can be written as λ = (Dτ)1/2, where τ is the mean residence time of the adatoms at D


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Figure 6. (a) AFM image of α-Al2O3 annealed at 1350 °C in H2 and (b) SEM plan-view image of sample H after etching in an aqueous KOH solution for 10 s at 80 °C.

the temperature needed to create oxygen vacancies in alumina is >1200 °C.47,48 The effects of annealing sapphire surfaces at high temperatures were previously observed for various annealing conditions, and similar morphological changes were reported.48−50 In particular, it was reported that the step edges became highly corrugated and the terraces became narrower or wider as the annealing temperature increased, because of the strong faceting and pairing of terraces. The features and intervals of the exposed wall-shaped Al-polar IDs after the KOH etching process (Figure 6b) and those of the meandrous step edges of the sapphire surface (Figure 6a) are clearly very similar. Thus, a simplified model is suggested to explain the formation of the rod-shaped (sample D) and wall-shaped (sample H) Al-polar IDs in the mixed-polarity AlN layers, as illustrated in Figure 7a−e and Figure 7a,f−i, respectively. In the rod-shaped Al-polar IDs, the conventional straight step-terrace structure of the vicinal sapphire becomes corrugated when the temperature is increased to 1300 °C to grow the AlN layer (Figure 7a,b). The AlOxN1−x interlayer is formed between the sapphire substrate and AlN layer during growth because AlN chemically reacts with the sapphire because of the diffusion of oxygen from the sapphire at high temperatures. This causes the polarity of the AlN layer on the AlOxN1−x to invert to N polarity.25,39 Consequently, sample D consists of a mixedpolarity AlN layer with rod-shaped Al-polar IDs (Figure 7c−e). In contrast, wall-shaped Al-polar IDs (sample H) form at higher growth temperatures and when the TMAl preflow process is used. During the TMAl preflow process, the migration length of Al adatoms is sufficient for the majority of them to diffuse to the meandrous step edges at 1350 °C (Figure 7a,f). Away from the meandrous step edges, the AlOxN1−x interlayer is formed between the sapphire substrate and the AlN layer, as in sample D. As a result, N-polar AlN layers with small Al-polar IDs predominantly grow on the terraces of the sapphire (Figure 5g−i). On the other hand, the Al-polar AlN layer is dominantly formed at the meandrous step edge of the sapphire because the formation of the N-polar AlN is impeded by the Al adatoms. Although it was not shown, the off-cut angle of the vicinal sapphire substrate is another important factor for controlling the crystallographic polarity of the AlN layer with a zigzag wall shape. The intentional increase in the off-cut angle can decrease the terrace width of vicinal sapphire within the Al adatom diffusion length to enhance the step-flow TMA preflow process. Consequently, zigzag wall-shaped Al-polar IDs can be easily formed with a short period at relatively low temperatures.

Figure 7. Growth models for formation of rod- and wall-shaped Alpolar IDs in samples D and H, illustrated in consecutive growth steps (b−e and f−i, respectively).

In this study, AlN nanopillars, nanorods, nanofurrows, and nanowalls with Al polarity were successfully fabricated on a sapphire substrate by etching with KOH. Figure 8 shows bird’seye-view SEM images of the AlN nanostructures fabricated by KOH etching of sample D (a and b) and sample H (c and d) for 30 s (a and c) and 150 s (b and d), and the insets show schematic images of each AlN nanostructure. The nanopillars were ∼300 nm in diameter, and their density was ∼2 × 109 cm−2. The nanorods were ∼150 nm in diameter and had a density of ∼1 × 109 cm−2. The aspect ratios of the nanopillars and nanorods were ∼0.43 and ∼2.7, respectively. The width and height of the nanofurrows were ∼230 and ∼160 nm, respectively, and those of the nanowalls were ∼100 and ∼360 nm, respectively. The TEM measurements clearly established that the fabricated AlN nanostructures likely had very few TDs. E


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Figure 8. SEM bird’s-eye-view images of various nanostructures fabricated by KOH etching of sample D (a and b) and sample H (c and d) with etching times of 30 s (a and c) and 150 s (b and d).

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CONCLUSION Mixed-polarity AlN layers for the fabrication of nanostructures with Al polarity were grown on a vicinal c-plane sapphire substrate by HT-MOCVD. The crystallographic polarity of the AlN layers changed from N-polar to Al-polar as the growth temperature decreased and the TMAl preflow rate increased. Mixed-polarity AlN layers with either rod-shaped or wallshaped Al-polar IDs could be grown by controlling the growth temperature and TMAl preflow rate. Finally, AlN nanopillars, nanorods, nanofurrows, and nanowalls with Al polarity were successfully fabricated by KOH etching of mixed-polarity AlN layers. TEM analysis of the behavior of the TDs in the mixedpolarity AlN layer indicated that the fabricated Al-polar nanostructures had a very high crystal quality. Thus, the fabricated AlN nanostructures can potentially be used in highefficiency optical and electronic applications.

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AUTHOR INFORMATION

Corresponding Author

*Address: P-503, Korea Polytechnic University (KPU), 237, Sangidaehak-ro, Siheung-si, Gyeonggi-do 429-793, Republic of Korea. Telephone: +82-10-3329-0324. Fax: +82-31-8041-1917. E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the National Research Foundation of Korea (NRF) funded by the Korea government (MEST) (Grant 2012R1A2A2A01011702).

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AlN Nanostructures Fabricated on a Vicinal Sapphire - American

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