Sapphire via Growth

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High-quality AlN film grown on sputtered AlN/sapphire via growth-mode modification Chenguang He, Wei Zhao, Hualong Wu, Shan Zhang, Kang Zhang, Longfei He, Ningyang Liu, Zhitao Chen, and Bo Shen Cryst. Growth Des., Just Accepted Manuscript • DOI: 10.1021/acs.cgd.8b01045 • Publication Date (Web): 18 Oct 2018 Downloaded from http://pubs.acs.org on October 22, 2018

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Crystal Growth & Design

High-quality AlN film grown on sputtered AlN/sapphire via growth-mode modification Chenguang He*,†, Wei Zhao†, Hualong Wu†, Shan Zhang‡, Kang Zhang†, Longfei He†, Ningyang Liu†, Zhitao Chen*,†, Bo Shen§ † Guangdong Institute of Semiconductor Industrial Technology, Guangdong Academy of Sciences, Guangzhou 510650, China ‡ School of Physics & Electronic Engineering, Guangzhou University, Guangzhou 510006, China § State Key Laboratory of Artificial Microstructure and Mesoscopic Physics, School of Physics, Peking University, Beijing 100871, China

Heteroepitaxy of high-quality AlN film is the key to advance the prosperity of deepultraviolet (DUV) devices when a large-size and low-cost native substrate is unavailable. Here, we proposed a strategy to obtain high-quality AlN film by combining growth-mode modification with sputtered AlN buffer using metal-organic chemical vapor deposition (MOCVD). Compared with the MOCVD AlN buffer, the sputtered AlN buffer consists of smaller and more uniform grains with better c-axis orientation, leading to a better growthmode modification in the subsequent growth process. On one hand, the better c-axis orientation is well inherited by the upper AlN epilayer, resulting in a lower screw dislocation density. On the other hand, the better growth-mode modification significantly suppresses edge dislocations by producing high-density nanoscale voids and many 90˚ bended dislocations. Therefore, the total threading dislocation density of the AlN film grown on the sputtered AlN buffer is dramatically reduced to an extremely low value of 4.7 × 107 cm-2, which is 81.2% less than that of the AlN film grown on the MOCVD AlN buffer. This very simple yet effective technique demonstrates great potential for the massfabrication of low-cost and high-performance DUV devices. *E-mail: [email protected] (Chenguang He) Phone: (+86) 61086425-8007 *E-mail: [email protected] (Zhitao Chen) Phone: (+86) 61086421 1

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Crystal Growth & Design 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

High-quality AlN film grown on sputtered AlN/sapphire via growth-mode modification Chenguang He*,†, Wei Zhao†, Hualong Wu†, Shan Zhang‡, Kang Zhang†, Longfei He†, Ningyang Liu†, Zhitao Chen*,†, Bo Shen§ †

Guangdong Institute of Semiconductor Industrial Technology, Guangdong Academy of

Sciences, Guangzhou 510650, China ‡

School of Physics & Electronic Engineering, Guangzhou University, Guangzhou 510006,

China §

State Key Laboratory of Artificial Microstructure and Mesoscopic Physics, School of

Physics, Peking University, Beijing 100871, China *E-mail: [email protected] (Chenguang He) *E-mail: [email protected] (Zhitao Chen)

ABSTRACT: Heteroepitaxy of high-quality AlN film is the key to advance the prosperity of deep-ultraviolet (DUV) devices when a large-size and low-cost native substrate is unavailable. Here, we proposed a strategy to obtain high-quality AlN film by combining growth-mode modification with sputtered AlN buffer using metal-organic chemical vapor deposition (MOCVD). Compared with the MOCVD AlN buffer, the sputtered AlN buffer consists of smaller and more uniform grains with better c-axis orientation, leading to a better growth-mode modification in the subsequent growth process. On one hand, the better c-axis orientation is well inherited by the upper AlN epilayer, resulting in a lower screw 2


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dislocation density. On the other hand, the better growth-mode modification significantly suppresses edge dislocations by producing high-density nanoscale voids and many 90˚ bended dislocations. Therefore, the total threading dislocation density of the AlN film grown on the sputtered AlN buffer is dramatically reduced to an extremely low value of 4.7 × 107 cm-2, which is 81.2% less than that of the AlN film grown on the MOCVD AlN buffer. This very simple yet effective technique demonstrates great potential for the massfabrication of low-cost and high-performance DUV devices. INTRODUCTION Recently, there have been increasing demands for AlGaN-based deep-ultraviolet (DUV) devices owing to their applications in water purification, sterilization, high-density optical storage, weak ultraviolet signal detection, etc.1, 2 To fabricate such novel devices, bulk AlN should be the most ideal substrate because of a similar in-plane lattice constant and thermal expansion coefficient with those of high-Al-content AlGaN epilayers. To date, however, the commercially available AlN substrate has not been an appropriate candidate for practical applications due to the limitation of small size, high cost, and DUV absorption.1, 3

Thus fabrication of DUV devices still greatly relies on the large-scale, low-cost, and

DUV-transparent AlN/sapphire template. Unfortunately, the well-known two-step heteroepitaxy usually produces plenty of dislocations (109-1010 cm-2) and cracks due to the large lattice and thermal mismatch between AlN and sapphire, as well as the low surface migration of Al species.1, 4 These defects will extend into the AlGaN active region, severely

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deteriorating device performance.5, 6 To address these issues, various strategies have been proposed, such as micro/nanoscale epitaxial lateral overgrowth (ELOG),7-11 migration-enhanced epitaxy (MEE),12, 13 and high temperature annealing (HTA).14 Effects of sapphire off-cut15 and surface pretreatments16-19 on crystalline quality have also been intensively studied. However, at present, the lowest threading dislocation density (TDD) of AlN films grown on sapphire is still in the range of 3-5 × 108 cm-2,9, 11, 14 which is much higher than that of GaN films grown on sapphire (105-107 cm-2).20-22 Moreover, the requirements of multiplestep processing and specially designed reactors are also big obstacles to these techniques’ practical applications. The growth-mode modification (GMM) technique, which is very simple yet effective, has been widely adopted in the past few years.23-26 The essence of this technique is to enhance dislocation interaction via modifying growth modes. First, side facets inclined to (0001) plane are intentionally created by introducing three dimensional (3D) growth. Then, when the growth mode transits from 3D to two dimensional (2D), the side facets of the 3D islands expand laterally. Driven by the image force, dislocations tend to incline toward the side facets of the 3D islands and interact with each other through merging and forming half-loops. In this way, the TDD can be reduced to 1-2 × 109 cm-2.26 Furthermore, tensile stress can also be effectively relaxed by introducing 3D growth.25 Despite these advances, the GMM technique still has much room for further improvement to meet the demands of dislocation-sensitive DUV devices.5, 6

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In this work, we demonstrated an approach to achieve high-quality AlN film by combing the GMM technique with sputtered AlN buffer. A conventional AlN buffer grown by metal-organic chemical vapor deposition (MOCVD) was also adopted to fabricate AlN film for comparison. It is found the evolutions of growth modes are quite different for AlN films grown on different AlN buffers. Compared with the MOCVD AlN buffer, the sputtered AlN buffer consists of smaller and more uniform grains with better c-axis orientation, leading to a better growth-mode modification in the subsequent growth process. As a result, the total TDD of the AlN film grown on the sputtered AlN buffer is dramatically reduced to an extremely low value of 4.7 × 107 cm-2, which is 81.2% less than that of the AlN film grown on the MOCVD AlN buffer. The detailed evolutions of growth modes and crystalline qualities, as well as the corresponding evolution mechanisms are presented.

EXPERIMENTAL SECTION Sample A was fabricated on a 2-inch (0001) sapphire substrate by an AIXTRON closecoupled showerhead (CCS) MOCVD system. Trimethylaluminum (TMAl) and ammonia (NH3) were used as the Al and N sources, respectively. High-purity hydrgen (H2) was employed as the carrier gas. First, a 20-nm-thick AlN buffer was deposited at 900 ℃ after 2 s of TMAl preflow. Then, a 300-nm-thick high temperature AlN (HT AlN-1) was grown at 1230 ˚C, followed by a 350-nm-thick low temperature (LT) AlN interlayer grown at 935 ˚C. Finally, sample A was finished with 4900-nm-thick high temperature AlN (HT AlN-2)

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grown at 1235 ˚C. The V/III mole ratios of the AlN buffer layer, HT AlN-1 layer, LT AlN interlayer, and HT AlN-2 layer were 68, 2500, 2500, and 166, respectively. The reactor pressure during the growth process was 50 mbar. The gaps between showerhead and susceptor were 6, 7, 7, 11 mm, respectively. The epitaxy of sample B was initialed with a 20-nm-thick AlN buffer grown by a NAURA iTops A230 AlN sputter system. A 2-inch aluminum disk (99.999%) was used as sputtering target. The AlN buffer was sputtered at 650 ˚C by feeding 120 sccm N2, 30 sccm Ar, and 1 sccm O2. The chamber pressure was 6.7 × 10-3 mbar. Then the epitaxy of sample B was conducted by MOCVD using the identical growth conditions with those of sample A. The reflectance transients were recorded by a LayTec EpiTT in-situ monitoring system. The surface morphologies were characterized by a FINIAL FJ-3A optical microscope (OM), a HITACHI SU8220 scanning electron microscope (SEM), in combination with a Bruker Dimension Edge atomic force microscope (AFM). The X-ray rocking curves (XRCs) were measured by a Rigaku Smartlab 9kW high-resolution X-ray diffraction (HRXRD) using a wide open detector. Cross-sectional and plan-view images were taken by a FEI Tecnai Osiris TF-20 scanning transmission electron microscope (STEM).

RESULTS AND DISCUSSION First, evolution processes of growth modes were investigated by in-situ reflectance transients, as illustrated in Figure 1a and 1b. During the growth processes of the HT AlN1 layers, the reflectance intensity of sample A rises slightly, while that of sample B remains 6


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constant. It suggests that the growth modes of sample A and B are 3D-2D transition and 2D growth, respectively.19 Then, when the LT AlN layers are introduced, the reflectance intensities of sample A and B both drop rapidly regardless of their former states, implying 3D growth mode. In comparison, sample B has a little higher reflectance intensity, corresponding to a weaker 3D growth process. Afterwards, when the HT AlN-2 layer is grown, the reflectance intensity of sample A continues to drop in the first 1170 nm (3D growth) and then rises slowly (3D-2D transition). The reflectance intensity isn’t saturated until the end of the growth, signifying an incomplete coalescence. It is supported by the undesirable pits shown in Figure 1c.7 Vastly different from the reflectance intensity of sample A, that of sample B rises rapidly (3D-2D transition) and reaches its maximum when the thickness of the HT AlN-2 layer reaches 2600 nm. After that, the growth mode is governed by 2D growth, leading to a smooth surface without any pits or cracks, as presented in Figure 1d. Obviously, the growth modes of AlN films are significantly affected by the AlN buffer layers.

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Figure 1. Reflectance (405 nm) transients recorded for (a) sample A and (b) sample B. An oscillation period corresponds to a thickness of 90 nm. Final surface images of (c) sample A and (d) sample B taken with OM using a magnification of 100.

To further give some insights into the evolutions of the growth modes, growth interruptions were performed on sample A and B at four specific stages. Figure 2a and 2b display the buffer morphologies of sample A and B taken with SEM using a magnification of 300000. It can be seen that the MOCVD AlN buffer consists of large and inhomogeneous crystalline grains with an average size of 30 nm, causing a large root mean square (RMS) of 1.17 nm over 3 × 3 µm2.19 By contrast, the sputtered AlN buffer is composed of quite small and uniform grains. The average grain size is only 10 nm, resulting in a flatter surface with a much smaller RMS of 0.26 nm over 3 × 3 µm2, which is similar to the reported values of thin sputtered AlN buffers.27 Such different buffer morphologies arise from the kinetics discrepancy between MOCVD and sputtering deposition. For the MOCVD AlN buffer, the quasi-equilibrium growth enables the Al species to migrate to the potential energy minima in sufficient residence time. In this situation, the density of nucleation sites is relative low, and grains forming at the nucleation sites expand in both lateral and vertical direction, leading to the formation of large grains. For the sputtered AlN buffer, however, the non-equilibrium growth makes the Al species just nucleate at the place where they can arrive. The high-density nucleation sites give rise to the formation of small grains.

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Figure 2. Surface morphologies of (a) the MOCVD AlN buffer and (b) sputtered AlN buffer taken with SEM using a magnification of 300000. The MOCVD AlN buffer consists of large and inhomogeneous crystalline grains with an average size of 30 nm, while the sputtered AlN buffer is composed of small and uniform grains with an average size of 10 nm.

In the subsequent growth processes, the morphology evolutions of sample A and B are quite different. Figure 3a and 3e present the surface morphologies of sample A and B after the growth of the 300-nm-thick HT AlN-1 layers. The surface of sample A involves plenty of pits with a density of 1.1 × 109 cm-2 and a diameter of 60-90 nm, because the large 3D grains are hard to coalesce in 300 nm.28 However, no pits are observed on the surface of sample B due to the 2D growth on the flatter sputtered AlN buffer. Figure 3b and 3f confirm the dominant 3D island-like morphologies after the growth of the 350-nm-thick LT AlN layers, which is ascribed to the limited migration of Al species at a low growth temperature of 935 ˚C.25 Similar to the difference in buffer morphologies, the density of 3D islands in sample A is lower than that in sample B, and the average size of 3D islands in sample A 9



(100 nm) is larger than that in sample B (70 nm). It is evident that the 3D growth of sample A is enhanced by the uncoalesced surface. Accordingly, the surface of sample B is flatter than that of sample A, leading to a little higher reflectance intensity. The larger 3D-island spacing of sample A delays the coalescence of the islands. Hence, when the thickness of the HT AlN-2 layer reaches 1170 nm, the surface of sample A is still dominated by columnlike islands with large spacing (Figure 3c). However, the column-like islands in sample B have already merged into larger areas, with only a few deep voids left (Figure 3g). The faster coalescence of the smaller 3D islands in sample B explains why the reflectance intensity of sample B rises faster than that of sample A during this process.

Figure 3. AFM images of (a-d) sample A and (e-h) sample B at different growth stages.

Finally, when the whole growth process is accomplished, the surface of sample A is featured by disorderly distributed steps (Figure 3d) even though the RMS is only 0.19 nm 10


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over 3 × 3 µm2. Some steps are terminated by defects, which are indicated with white arrows. Unlike sample A, sample B achieves long parallel atomic steps without any terminations (Figure 3h). The RMS is as small as 0.14 nm over 3 × 3 µm2, and the step height is about 0.27 nm, corresponding to one monolayer of (0001) AlN (c/2=0.25 nm). Overall, compared with the AlN film grown on the rough MOCVD AlN buffer, that grown on the flatter sputtered AlN buffer demonstrates a weaker 3D growth, an earlier 3D-2D transition, and an adequate 2D growth, facilitating the achievement of a crack-free and atomically-flat surface. To investigate the effects of the growth modes on crystalline qualities, we performed HRXRD measurements. Figure 4a and 4b compare the FWHMs of symmetric (002) and asymmetric (102) reflections for sample A and B at different growth stages. It can be seen that the (0002) plane XRC FWHM of the sputtered AlN buffer (349 arcsec) is much smaller than that of the MOCVD AlN buffer (1382 arcsec). As a result, sample B has smaller FWHMs of (002) reflection across the whole growth process, indicating sample B always has lower screw-type dislocation density. The total TDDs of sample A and B at different growth stages can be estimated by using the empirical formula N = FWHM2/4.35|b|2, where b is the Burgers vector of the corresponding dislocation, as depicted in Figure 4c and 4d. Actually, this method is inaccurate when dislocations are not randomly distributed, wherein the grain diameter must be taken into account for correction.29 On the other hand, x-ray has a large penetration depth in AlN, considering the linear absorption coefficient is only 119 cm-1.30 It means the XRC FWHMs contain the information of the whole AlN epilayer. 11



Therefore, here, we just utilize it to see the rough variations of TDDs. For both sample A and B, the TDDs decrease during the process of 3D-2D transition and 2D growth, and increase during the 3D growth process. This variation trend is similar to that of (102) FWHMs, because dislocations with edge component are usually dominant in III-nitrides.31 The 3D growth process of sample A is so long that the TDD increases to a very large value, which is hard to be reduced in the subsequent 3D-2D transition process. In the case of sample B, however, the TDD only rises slightly during the shorter 3D growth process. Moreover, the TDD can be effectively reduced during the process of 3D-2D transition and 2D growth. Consequently, the final TDD of sample B is much lower than that of sample A.

Figure 4. FWHMs of symmetric (002) and asymmetric (102) reflections for (a) sample A and (b) sample B at different growth stages. The FWHMs of (102) reflections for MOCVD 12


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buffer and sputtered buffer are not measurable because they are very large. The corresponding TDDs of (c) sample A and (d) sample B roughly estimated by using the empirical formula.

To investigate the evolutions of crystalline qualities in more detail, we then performed cross-sectional bright-field (BF) STEM. Figure 5 shows the high-resolution transmission electron microscope (HRTEM) result. The selected area electron diffraction (SAED) patterns in Figure 5b and 5e determine that the MOCVD AlN buffer and sputtered AlN buffer are both hexagonal wurtzite structure, and the relative rotation angle between the AlN buffers and sapphire substrates is 30˚. The difference between them lies in the different atomic arrangements originating from the kinetics discrepancy mentioned above. At the initial stage of growth, the lattice arrangement of the MOCVD AlN buffer is disordered, producing a distinct amorphous zone (thickness < 2 nm) at the AlN/sapphire interface. However, the sputtered AlN buffer demonstrates highly-aligned atomic planes from the very beginning, creating a well-defined AlN/sapphire interface. In the subsequent growth process, the sputtered AlN buffer always exhibits better atomic arrangement, because the misorientations between the small and uniform grains are smaller. In consequence, the upper AlN epilayer of sample B also owns better c-axis orientation. This explains why sample B has much smaller FWHMs of (002) reflection across the whole growth process.

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Figure 5. Cross-sectional HRTEM images of (a) sample A and (d) sample B. (b) and (e) are the SAED patterns of the AlN/sapphire interfacial regions, manifesting the MOCVD AlN buffer and sputtered AlN buffer are both hexagonal wurtzite structure and the relative rotation angle between AlN buffers and sapphire substrates is 30˚. (c) and (f) are the enlarged HRTEM images of the dashed areas in (a) and (d). Sample B always exhibits better atomic arrangement, as illustrated in the schematic diagrams.

Figure 6 shows the cross-sectional BF STEM images of sample A and B, predominantly revealing strain contrast around dislocation lines. It means screw, edge, and mixed -type dislocations can be observed at the same time. At the beginning of the growth, for both sample A and B, high-density dislocations are generated from the AlN/sapphire interfaces due to a large mismatch between AlN and sapphire. Some dislocations are 14


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quickly annihilated with bundles or agglomerates at the initial stage of HT AlN-1 layers, benefiting from the close distances between neighboring dislocations. Then the surviving dislocations propagate upwards along [0001] direction. The TDDs demonstrate apparent decrease when the thicknesses of the HT AlN-2 layers reach certain values. Compared with sample A, sample B exhibits an earlier decrease in the TDD due to an advanced 3D-2D transition. For both sample A and B, dislocations are mainly suppressed through three channels. Some dislocations are reduced through merging and forming half-loops due to the small-angle inclinations, which is commonly observed in AlN film grown by GMM technique.25 Besides, dislocations are also effectively suppressed through two new channels: high-density self-organized voids and 90˚ dislocation bending in the areas away from the voids.

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Figure 6. Cross-sectional BF STEM images of (a) sample A and (d) sample B. (b) and (c) are the enlarged images of the red and green dashed areas in (a), respectively. (e) and (f) are the enlarged images of the red and green dashed areas in (d), respectively.

Generally, voids are hard to be observed in AlN films except for the case of ELOG. In this work, the formation of self-organized voids originates from the high-speed lateral growth rate and large depth/width ratio between 3D islands. During the growth processes of HT AlN-2 layers, the growth rate is as fast as 3 µm/h. The low reactor pressure (50 mbar), high growth temperature (1235 ˚C), and low V/III mole ratio (166) facilitate high-speed lateral growth, so the upper parts have coalesced before the lower parts complete the growth. For sample A, lots of voids with a length of 150-1300 nm are formed when the HT AlN-2 layer is grown. Some large voids extend to the top surface (Figure 6a), responsible for the pits appearing on the surface. For sample B, due to a faster coalescence of 3D islands, high-density nanoscale voids with a shorter length of 40-700 nm are formed and fully coalesce when the thickness of the HT AlN-2 layer reaches 1500 nm. Similar to the case of ELOG, the pre-existing dislocations can bend and terminate at the local free surfaces provided by these voids, following the principle of dislocation line energy minimization.9 Different types of dislocations achieve their energy minima at different bending angles, depending on their Burgers vectors and line directions.32 For edge-type dislocation, 44˚±7˚ and 90˚±2˚ off the c-axis are calculated to be two typical bending angles, corresponding to 16


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the local energy minimum and energetically most stable direction, respectively. When an edge-type dislocation bends in a few steps, it will pass through the local energy minimum (44˚±7˚) and reach the energetically most stable direction (90˚±2˚). In this work, those dislocations bending towards the voids are mainly edge-type, and the bending angles are about 50˚ and 90˚. It is in good accordance with the calculated result. Because of a higher island density, sample B has a higher void density of 1.7 × 1010 cm-2 in the (0001) plane, providing a higher probability for dislocation termination. During the coalescence processes of the voids, there exists obvious difference between sample A and B. As shown in Figure 7a, many dislocations are generated at the coalescence boundaries in sample A, while nearly no dislocations can be found at the coalescence boundaries in sample B. It is assigned to their difference in the spacing and orientations of the adjacent domains near the voids. As mentioned above, the sputtered AlN buffer provides small and uniform islands with excellent c-axis orientation, thus the domains around the voids in sample B have smaller spacing and better c-axis orientation. It means the misorientations between the adjacent domains are much smaller, so dislocations at the coalescence boundaries can be effectively suppressed.

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Figure 7. (a) Dislocation behaviors near the self-organized voids. For sample B, on one hand, the higher void density provides a higher probability for dislocation termination. On the other hand, the smaller domain spacing and better c-axis orientation effectively suppress dislocations at the coalescence boundaries of the voids. (b) Dislocation behaviors near the macrosteps. Some macrosteps with different heights are observed on the final surface of sample B over a scanning area of 15 × 15 µm2, which is responsible for the 90˚ dislocation bending. The schematic demonstrates that the bending angle is proportional to the ratio of lateral growth rate to vertical growth rate. The bended dislocation is easy to come in contact with other dislocations and block them.

The 90˚ dislocation bending in the areas away from the voids is found in both sample A and B, but the bending probability in sample B is much higher, as shown in Figure 6c and 6f. When the thicknesses of the HT AlN-2 layers exceed 1500 nm, the majority of the 18


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surviving dislocations in sample B exhibit a 90˚ bending, whereas only few dislocations in sample A behave like this. By changing diffraction conditions, these bended dislocations are confirmed to be pure-edge type. Different from the aforementioned 90˚ dislocation bending towards the voids, the 90˚ dislocation bending described here is likely associated with macrostep movement. In the case of 90˚ dislocation bending towards the voids, the presence of side facets is the prerequisite, which is similar to the case of ELOG.9, 33 Here, although the macrostep edge can also be regarded as a small side facet, the difference lies in that the bended dislocation is included by both the small side facet and the normal surface , as shown schematically in Figure 7b. When the macrostep advances laterally and meets with the dislocation in the normal surface, the dislocation will bend as the macrostep and surface advance, and the bending angles are proportional to the ratio of lateral growth rate to vertical growth rate.34 In the previous work, the bending angles is only 25˚, corresponding to a ratio of 1/2. 34 It is attributed to the poor lateral migration of Al species at 700 ˚C. In this work, the bending angle is about 90˚ (actually 80˚-90˚), requiring the growth rate ratio reaches 6 at least. During the growth processes of the HT AlN-2 layers, the lateral migration of Al species is promoted by the low reactor pressure (50 mbar), high growth temperature (1235 ˚C), and low V/III mole ratio (166). It is easy to achieve a high ratio exceeding 6. For sample A, the always-existing uncoalesced voids significantly hinder the lateral macrostep movements and weaken the bending effect. In consequence, many residual dislocations can still be observed near the top AlN surface. For sample B, those voids have 19



fully coalesced when the thickness of the HT AlN-2 layer reaches 1500 nm. The void-free surface is in favor of the lateral macrostep movements, accounting for the much higher bending probability in sample B. The 90˚ bended dislocations propagate laterally. Therefore, they are easy to come in contact with other dislocations and block them by merging or forming half-loops, as shown in Figure 7b.34 As a result, nearly no dislocations can reach the top AlN surface of sample B. To accurately evaluate the TDDs of sample A and B, we finally performed plan-view STEM. Figure 8a and 8e confirm that the selected areas of sample A and B are free from pits and cracks. So the strain contrast observed in Figure 8b-d and Figure 8f-h completely originates from dislocations. For sample A, dot-type dislocations vertical to the (0001) plane are dominant. In the case of sample B, there also exist some line-type dislocations in the (0001) plane, corresponding to the 90˚ bended dislocations. These results are in outstanding agreement with the dislocation behaviors observed by the cross-sectional STEM. According to the density of dot-type dislocations (white circles) in Figure 8b and 8f, the TDDs of sample A and B were calculated to be 2.5 × 108 cm-2 and 4.7 × 107 cm-2, respectively. A 81.2% reduction of TDD is realized by using the sputtered AlN buffer. The TDD of 4.7 × 107 cm-2 is an extremely low value for AlN films grown on sapphire.9, 11, 14

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Figure 8. Plan-view Z-contrast STEM images of (a) sample A and (e) sample B. Plan-view strain-contrast STEM images of (b) sample A and (f) sample B. The white circles correspond to dot-type dislocations. (c) and (d) are the enlarged images of the green and red dashed areas in (b). (g) and (h) are the enlarged images of the green and red dashed areas in (f).

CONCLUSIONS In summary, a 5.6-µm-thick AlN film with excellent surface morphology and crystalline quality is achieved by combining the GMM technique with sputtered AlN buffer. Compared with the MOCVD AlN buffer, the sputtered AlN buffer consists of smaller and more uniform grains with better c-axis orientation. On one hand, the better c-axis orientation is well inherited by the upper AlN epilayer of sample B, leading to a lower screw dislocation density. On the other hand, the sputtered AlN buffer results in a weaker 3D growth, an earlier 3D-2D transition, and an adequate 2D growth in the subsequent 21



growth process. During the processes of 3D-2D transition and 2D growth, the better growth-mode modification significantly suppresses edge dislocations in sample B by producing high-density nanoscale voids and many 90˚ bended dislocations, which are promoted by high-speed lateral growth. As a consequence, the total TDD of sample B is dramatically reduced to an extremely low value of 4.7 × 107 cm-2, which is 81.2% less than that of sample A. This very simple yet effective technique demonstrates great potential for the mass-fabrication of low-cost and high-performance DUV devices.

AUTHOR INFORMATION Corresponding Authors *E-mail: [email protected] (Chenguang He). *E-mail: [email protected] (Zhitao Chen). Notes The authors declare no competing financial interest.

ACKNOWLEDGMENTS This work is supported by the National Key Research and Development Program of China (No. 2017YFB0404100), National Natural Science Foundation of China (Nos. 61604045 and 61804034), Innovation-driven Development Capacity Construction Project of the Guangdong Academy of Sciences (Nos. 2017GDASCX-0845, 2017GDASCX-0410, and

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2018GDASCX-0952), the Science and Technology Program of Guangdong (Nos. 2015B010112002 and 2017A010103038), Natural Science Foundation of Guangdong Province (No. 2015A030310023), and Pearl River S&T Nova Program of Guangzhou (No. 201610010142). The authors thank Dr. Dong Boyu and Guo Bingliang from Beijing NAURA Microelectronics Equipment Co., Ltd. for providing sapphires with sputtered AlN buffers. The authors are also grateful to Dr. Lisheng Zhang and Mingxing Wang from Peking University for discussions.

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For Table of Contents Use Only

High-quality AlN film grown on sputtered AlN/sapphire via growth-mode modification Chenguang He*,†, Wei Zhao†, Hualong Wu†, Shan Zhang‡, Kang Zhang†, Longfei He†, Ningyang Liu†, Zhitao Chen*,†, Bo Shen§ †

Guangdong Institute of Semiconductor Industrial Technology, Guangdong Academy of Sciences, Guangzhou 510650, China ‡ School of Physics & Electronic Engineering, Guangzhou University, Guangzhou 510006, China § State Key Laboratory of Artificial Microstructure and Mesoscopic Physics, School of Physics, Peking University, Beijing 100871, China *E-mail: [email protected] (Chenguang He) *E-mail: [email protected] (Zhitao Chen)

We proposed a strategy to obtain high-quality AlN film by combining growth-mode modification with sputtered AlN buffer using metal-organic chemical vapor deposition (MOCVD). Compared with the AlN film grown on the MOCVD AlN buffer, that grown on the sputtered AlN buffer demonstrates an extremely low threading dislocation density. The detailed dislocation behaviors and the corresponding evolution mechanisms are revealed.

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Sapphire via Growth

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