Growth Model of van der Waals Epitaxy of Films: A Case of AlN Films

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Growth model of van der Waals epitaxy of films: A case of AlN films on multilayer graphene/SiC Yu Xu, Bing Cao, Zongyao Li, Demin Cai, Yumin Zhang, Guoqiang Ren, Jianfeng Wang, Lin Shi, Chinhua Wang, and Ke Xu ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b14494 • Publication Date (Web): 28 Nov 2017 Downloaded from http://pubs.acs.org on December 2, 2017

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ACS Applied Materials & Interfaces

Growth model of van der Waals epitaxy of films: A case of AlN films on multilayer graphene/SiC Yu Xu,†,‡ ,§ Bing Cao,*,║,┴ Zongyao Li,‡ Demin Cai,‡ Yumin Zhang, †,‡ ,§ Guoqiang Ren, †,‡ Jianfeng Wang, †,‡ Lin Shi,† Chinhua Wang, ║,┴ and Ke Xu*,†,‡ †

Suzhou Institute of Nano-Tech and Nano-Bionics (SINANO), Chinese Academy of Sciences

(CAS), Suzhou 215123, P. R. China ‡

Suzhou Nanowin Science and Technology Co., Ltd., Suzhou 215123, P. R. China

§

University of Chinese Academy of Sciences, Beijing 100049, P. R. China

║

College of Physics, Optoelectronics and Energy and Collaborative Innovation Center of Suzhou

Nano Science and Technology, Soochow University, Suzhou 215006, P. R. China ┴

Key Lab of Advanced Optical Manufacturing Technologies of Jiangsu Province and Key Lab of

Modern Optical Technologies of Education Ministry of China, Soochow University, Suzhou 215006, P. R. China KEYWORDS: van der Waals epitaxy· growth model· orientation relationship· multilayer graphene· AlN films ABSTRACT: ‘Volmer–Weber’ island nucleation and step-flow growth model are the classical processes of the conventional epitaxy of films. However, growth model of van der Waals epitaxy of films is still not very well-documented. Here, we present an example of vdWE of AlN films on multilayer graphene/SiC by hydride vapor phase epitaxy at high temperature of 1100 oC, and

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reveal the orientation relationship of AlN, MLG, and SiC as (0001)[1-100]AlNǁ(0001)[1100]MLGǁ(0001)[11-20]SiC, which suggests that vdWE heterointerface is not an usual covalent bonding and no excessive strain during the growth process owing to the incommensurate inplane lattices. Remarkably, zigzag cracks are formed due to the anisotropy of strain after the films cooling down to room temperature, indicating that the growth model of vdWE is different from the one of conventional epitaxy. It is a layer-by-layer epitaxy, and planar substrate without miscut angle should be essential for obtaining single crystalline films. Additionally, the films can be transferred to foreign substrates by direct mechanical exfoliation without any stressor layer. An ultraviolet photosensor device illustrates an example of III-nitride heterogeneous integration application. Our work demonstrates an excellence step towards vdWE of varieties of compounds films on 2D materials for the applications of transferrable heterogeneous integration in future. 1. INTRODUCTION Van der Waals epitaxy (vdWE) is different from the conventional epitaxy due to the noncovalent bond between epitaxial material and substrate at the interface.1 As a result, although the orientation relationship of epitaxial material consist with the substrate, lattice match is not important for the vdWE, which is called incommensurate epitaxy; the in-plane lattice parameters of the 3D materials will be very close to the bulk value, results in no excessive strain and dislocation originated form the lattice mismatch. This interface phenomenon has been exposed by vdWE of 3D semiconductor on 2D materials, such as II-VI nanocrystals/nanowire on muscovite,2-4 II-VI nanowire/nanorod on graphene/graphite,5-10 and III-V nanocolumns/nanowire on graphene.11-15 However, all these previous studies focused on the nonplanar nanoarchitectures, the published data on utilization and characteristics of vdWE planar materials (films) is still


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sparse,16-20 and its generality, especially the growth model of vdWE films, is not very welldocumented to date. Graphene, a 2D material, has attracted a tremendous amount of attentions over the past few years due to its unique physical properties.21 The special honeycomb lattice structure makes it possible for growing c-plane III-nitride films and transferring onto foreign substrates.16 Although superiority of graphene is apparent, direct growth of III-nitride films on graphene faces is challenging due to the lack of chemical reactivity of the surface. Defects of graphene introduced by O2 plasma,16 ozone,19 fluorination treatment,22 and steps/terraces of multilayer graphene/SiC (MLG/SiC)23 were beneficial to nucleation due to potential enhanced chemical reactivity. As a result, the exact bonding at the nucleation heterointerface is still a matter of debate. Although single crystalline films without grain boundary can be obtained by using monolayer graphene,17, 24

it is not real vdWE, which is called ‘remote epitaxy’,24 whose heterointerface bond is covalent

bond. Therefore, MLG should be used for real vdWE, even though grain boundaries exist in films inevitable due to the polycrystallinity of MLG. Although the heterointerface of AlN and graphene has been reported by A. Kovács et al.,25 the lateral overgrown process may make mistakes to judge the orientation relationship. As 19.1o, 5, 23

0o, 26 and 30o 15 rotation of GaN or AlN lattice on graphene have been reported by now, there

is no clear evidence of the orientation relationship of AlN and graphene. In addition, as a low migration energy material, AlN films illuminate the growth model of vdWE films perfectly due to the lower selectivity of the nucleation, which were usually used as a low temperature buffer layer,15, 25, 27 few literatures reported direct vdWE of AlN films on graphene at high temperature. Furthermore, for ultraviolet light emitting diodes (UV LEDs), the AlN films are the crucially foundational materials for the device fabrication. However, the external quantum efficiency of


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UV LEDs (at wavelength of 200~320 nm) was mostly lower than 10% owing to the low light extraction efficiency (LEE).28 Few attempts have been adopted to improve the LEE by vertical structure LEDs due to the difficulty of releasing devices from the substrate and transferring to others.29 Therefore, vdWE of AlN films on MLG illuminates the incomparable superiority for the fabrication of transferable and vertical UV LEDs to realize the LEE improvement by UV light emission from N face. In this study, we perform the direct vdWE of AlN films on the MLG/SiC by hydride vapor phase epitaxy (HVPE) at high temperature of 1100 oC. The MLG was in situ treated at the NH3 atmosphere before growth of AlN films at high temperature of 1100 °C for enhancing the chemical reactivity of the surface. C-axis oriented AlN films were fabricated due to the honeycomb lattice of graphene. The orientation relationship of AlN, MLG and SiC was identified, which presented a generic atomic model to describe the vdWE growth configuration. Remarkably, the origination of fantastic zigzag cracks revealed that the growth model of vdWE was different from the one of conventional epitaxy, which plays an important role for understanding the growth model of vdWE of planar materials and making a good guidance for vdWE of other planar materials. Moreover, the AlN films can be transferred to foreign substrates by direct mechanical exfoliation without any stressor layer, and showed the practicability of transferable UV photosensor device. This study demonstrates a generality of vdWE films and lays a foundation of further insight into growth of AlN films on 2D materials for the applications of transferable, vertical UV LEDs. 2. RESULTS AND DISCUSSION Figure 1a depicts low magnification surface morphology of the AlN films grown on MLG/SiC substrate by HVPE. There are many cracks formed in AlN films after cooling down to


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room temperature. The overall direction of the cracks in AlN films is parallel to the [11-20]SiC direction and the average separation of the films is 60~140 µm. Further observation shows that the cracks exhibit wave shape. Therefore, AlN films display the morphology of several zigzag microbelts. Some of the zigzag microbelts rolled up from the MLG/SiC substrate. A representative SEM image of one of AlN zigzag microbelt is shown in Figure 1b. Corrugated edge corresponding to up and down sides of the microbelt can be clearly observed and the angle between two adjacent zigzag edges is about 120°. The width of the zigzag microbelt is about 70 µm. A representative high-magnification surface morphology of the microbelt is displayed in Figure 1c. Hexagonal pyramid nanohillocks randomly distribute on the surface, no nanocrack can be seen in microbelts. Figure 1d illuminates the cross-section SEM image of the AlN films after release from the substrate by direct mechanical exfoliation. The thickness of the films is about 670 nm with the height of the nanohillocks is about 30~100 nm. As the partial AlN microbelts rolled up from the MLG/SiC substrate, the AlN films can be transferred onto electroconductive tape by direct mechanical exfoliation without any stressor layer. It is interesting to determinate the released interface. Figure 2a~b demonstrates the high magnified AFM images of the surface morphology of MLG/SiC after AlN films release and the back surface of the AlN films after release. Clear stripes with the average spacing of about 100~200 nm can be seen at the surface of MLG/SiC substrate after AlN films release (Figure 2a), which are the steps and terraces of the MLG/SiC after graphitization before growth of the AlN films.30 The similar steps and terraces at back surface of the AlN films after release are shown in Figure 2b, which indicate that those features inherit from the surface of MLG/SiC substrate as the width of terraces is the same and there is no nanocrack in AlN films. Raman spectra were measured as shown in Figure 2c, the Raman peaks of AlN, SiC, and graphene can be clearly


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detected from the AlN/MLG/SiC. After release, SiC and graphene peaks exist without AlN peaks on the surface of MLG/SiC substrate, AlN and graphene peaks exist on the back surface of AlN films after release. The complete coverage of graphene on the back surface of AlN films was investigated by Raman mapping of D, G, and 2D peaks (shown in Figure S1 in Supporting Information), suggesting that the AlN films reliably nucleate on the graphene rather than SiC substrate and release together with graphene. The D peak (~1350 cm-1) enhances and D’ peak (~1620 cm-1) appears in AlN films after release, while only G and 2D peaks (~1580 cm-1 and ~2700 cm-1) exist on the MLG/SiC after release. This may be because of boundary-like defects formed from the defective graphene after release and defect-free few layers graphene reserved on the surface of SiC substrate. The blueshift of 2D peaks of the graphene in AlN films after release may be attributed to the different strain derived from SiC and AlN (the situation of graphene will be discussed later). It draws a conclusion that the separation of the films actually takes place within graphene layers, which act as a release layer for mechanical transfer owing to the weak vdW forces. It is similar to the GaN-based devices transferred using layered BN as a release layer,20 where BN was reserved at both of the separated surface after release. We investigate the crystalline orientation of the AlN films by XRD and plan-view TEM measurement. Figure 3a shows the θ-2θ configuration performed by power XRD for the AlN films grown on MLG/SiC substrate. Two dominated diffraction peaks at 36.0° and 76.2° can be detected, which are the wurtzite AlN (0002) and (0004) diffraction peaks. A very weak peak at 37.9° shows the (10-11) diffraction, indicating that c-plane dominated AlN films have been grown on the MLG/SiC. P. Gupta et al.27 reported preferential semipolar (10-11) orientation of AlN films grown on graphene with two steps at 800 °C and 1040 °C, which suggests that the nucleation temperature of AlN (~1100 °C) on MLG may be one of essential factors for obtaining


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c-axis oriented AlN films. Furthermore, high-resolution XRD 2θ/ω scans of the AlN films shows SiC (0004) and AlN (0002) diffraction peaks (Figure 3b), the full width at half maximum (FWHM) of the AlN (0002) diffraction peak was approximately 1°, illuminating that the AlN films are highly c-axis oriented crystal structure. This FWHM is similar with the GaN films (~0.8°) 31 grown on CVD graphene and much smaller than those (3-6°)

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exfoliated graphene. For the plan-view TEM sample, An AlN microbelt was transferred directly onto a copper grid with a 1 mm diameter through-hole for characterization (as shown in Figure 3c). A representative selected area electron diffraction (SAED) patterns image displays clearly six-fold symmetry as shown in Figure 3d. The (10-10), (20-20), (30-30), (11-20), (21-30), (2240), and (31-40) diffraction spots can be clearly seen in the SAED patterns, (10-11) diffraction spot was not observed due to the nanoscale local detection of SAED. It suggests that the AlN films are wurtzite structure with highly oriented in the c-axis direction. The honeycomb lattice of graphene makes it possible for growing c-plane AlN films. Therefore, both of XRD and SAED display the consistence of crystalline orientational relation of vdWE films and the substrate. Detailed microstructure of AlN films analyzed by TEM is shown in Figure 4. Cross-section bright field TEM image in Figure 4a reveals that AlN films were actually grown on MLG. Figure 4b depicts a high magnification TEM image of the interface region marked with a white rectangular box in Figure 4a, graphene can be observed at the interface of AlN and SiC obviously, and the steps of SiC along [11-20]SiC are evidently exhibited. The thickness of graphene seems different at the interface. In fact, owing to thermal mismatch of AlN and SiC, the AlN films shrank after cooling down to room temperature, resulted in malposition of interface. Furthermore, high-resolution TEM (HRTEM) illuminates that the top layers of graphene are defective, and only few layers of defect-free graphene exist on the surface of SiC


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substrate (Figure 4c). Nitrogen defects were introduced into the graphene lattice via NH3 treatment during the nucleation process, as reported by Z. Y. Al Balushi et al.23 Here, the higher temperature of 1100 °C for NH3 annealing may be the reason of more obvious defects formed on the surface of graphene. Just because of defective graphene, which enhanced chemical reactivity of the surface, nucleation of AlN was abundant to form films within such thin thickness of 670 nm. In addition, clearer images can be seen in Figure 4d-e, which are the magnified HRTEM images of the region marked with white rectangular boxes in Figure 4c. To examine the orientation relationship between AlN and SiC in detail, SAED patterns were investigated at the interface of AlN and SiC. Figure 4f-g are corresponding SAED patterns of AlN and SiC at the region recorded in Figure 4d-e. In view of the arrangement of spots, those reveal that the zone axes of SiC and AlN are [1-100]SiC and [11-20]AlN, which indicate that [1-100]AlN direction is rotated by ~30° from the [1-100]SiC direction. However, when AlN films were directly grown on SiC, the orientation relationship is characterized as [1-100]AlNǁ[1-100]SiC (see Figure S2 and Table S1 in Supporting Information). The 30° rotation of AlN towards SiC was derived from the rotation of graphene and SiC, which can be detected by low energy electron diffraction during the epitaxy of graphene on SiC.33, 34 Figure 4h-i shows the experiment and simulated convergent beam electron diffraction (CBED) patterns obtained from the AlN films. It reveals that the AlN films are Al-polar by vdWE on MLG, which is in close accordance with the simulated results at the similar sample thickness of about 84 nm. Therefore, the orientation relationship of AlN, MLG, and SiC can be determined as (0001)[1-100]AlNǁ(0001)[1-100]MLGǁ(0001)[11-20]SiC. Although defective graphene exists at the interface, the orientation relationship of AlN and MLG assures the presence of vdWE. The impact of graphene properties on GaN and AlN nucleation


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has been investigated by Z. Y. Al Balushi et al.,23 who revealed that defective graphene was more beneficial to nucleation. The in-plane lattice parameters of the AlN and graphene are 0.3112 nm and 0.2468 nm, respectively. Such orientation relationship results in a large lattice mismatch of 26.1%. As we all known, for conventional epitaxy with large lattice mismatch, such as ZnO/Al2O3 (-31.7%), GaN/Al2O3 (-33%), and AlN/Al2O3 (-34.6%), the orientation relationship is (0001)[1120]overlayerǁ(0001)[1-100]substrate. For vdWE, the given supercell may results in good coincidence with small misfit, such as (3 x 3) ZnO on (4 x 4) hBN (-2.5%),35 (3 x 3) GaN or ZnO on (4 x 4) graphene (-3.1%),26, 36 and (12 x 11) ZnO on (13 x 4) mica (0.02 and -1.11%),3 which lead to the orientation relationship of (0001)/(000-1)[1-100]overlayerǁ(0001)[1-100]substrate. Therefore, vdWE of AlN films on graphene exhibits the incommensurate in-plane lattices at the heterointerface. Here, we build a representative supercell with an AlN slab (4 x 4) on a bilayer graphene slab (5 x 5), which shows the lattice mismatch of 0.9%. A first principle calculation result is shown in Figure 5. The interaction energy (∆E=Etatal-EAlN-EMLG, where Etatal is the total energy per adatom of the AlN adatom layer on graphene, EAlN and EMLG are the energy of the isolated AlN adatom layer and bilayer graphene, respectively.) of the supercell is -1.08 eV. As a result, the total binding energy per atom is about 67.5 meV. To investigate the origination of the zigzag cracks in AlN films, the features of the zigzag cracks were characterized firstly. As Figure 3 discussed above, the direction of the zigzag edge is the [11-20]AlN (A direction) as the red line in Figure 3c parallel to the one in Figure 3d, namely the cross-section plane of the zigzag edge is (1-100) plane (M plane), which is the cleavage plane of the hexagonal crystal system material. Therefore, the overall direction of the cracks in AlN films is parallel to the [1-100]AlN direction. Here, Figure 6a shows the surface morphology of the


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AlN zigzag crack, the direction of which is marked with red arrows. Magnified surface morphologies at the crack regions marked with white rectangular boxes are shown in Figure 6b~e. Obvious steps of the SiC can be seen in the gap of the cracks, which point out the direction of the SiC marked with white arrows in the Figure 6b~e. This result conforms to the conclusion of the plane-view TEM analysis and confirms the orientation relationship of AlN, MLG, and SiC macroscopically. On the other hand, each side of cracks has corresponding half of hexagonal pyramid nanohillocks, which indicates that the cracks take place after cooling down to room temperature. In addition, microstructure characterization of vdWE AlN films on MLG/SiC was carried out further. As a substrate of 4° off axis miscut, the degree of the MLG/SiC step is 120o as shown in Figure 7 (also can be seen in Figure 4a). The lattice plane of the MLG is parallel to the surface of SiC. As a result, whether on the steps or terraces of MLG/SiC, the c-plane of the AlN is parallel to the surface of MLG/SiC, namely, the c-axis of AlN rotates 60° from the c-axis of SiC on the steps. Therefore, the high-angle grain boundaries (HAGBs) formed when the coalescence of AlN nuclei. (Here, Figure 7 is a joggled image, because the AlN films shifted from the original position due to thermal mismatch of AlN and SiC.) As a low migration energy material, AlN nuclei show random nucleation at the surface of graphene at a high temperature of 1100 oC (as shown in Figure S3 in Supporting Information). That is, HAGBs exist widespreadly in the AlN films grown on MLG/SiC with 4° off axis miscut. The weak (10-11) diffraction peak in XRD also verifies the results, as the interfacial angle of (0001) and (10-11) plane is 61.6°, which is similar with the included angle of c-axis of AlN on terraces and steps (60°). Figure 8a shows the schematic illuminations of the difference of growth model of conventional epitaxy and vdWE. In the case of conventional epitaxy, single crystal AlN


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nucleates preferentially at the steps of the SiC. It is the usual covalent bond formed at the heterointerface of the AlN and SiC due to the abundant dangling bonds at the steps of SiC. As a result, the orientation of the AlN nuclei is equal to the SiC to unify single crystal films (AlN caxis parallel to SiC c-axis), no matter the AlN nuclei are on terraces or steps. Namely, the miscut of AlN films is also 4 oC off axis according to the SiC substrate. Therefore, conventional epitaxy is a process containing nucleation on steps of substrate, coalescence of nuclei (island), and growth of films. It is a classical model termed ‘Volmer–Weber’ island nucleation and step-flow growth. The substrate with a miscut angle is beneficial to epitaxy, such as 0.2° sapphire substrate, 0.4° Ga-polar GaN substrate, 4° N-polar GaN substrate, 15° GaAs substrate, and 6° Ge substrate. In the case of vdWE, the in-plane lattice of AlN heterointerface is parallel to the surface of MLG, whether the nucleation of AlN takes place on the step or not. Although the beginning bonds of nucleation of vdWE maybe covalent bond, the dominating bonds at the vdWE heterointerface are noncovalent bond as epitaxial growth of the nuclei (island) follows the lattice plane of MLG. This is the interesting and tremendous difference of the growth model of vdWE and conventional epitaxy. We define the ‘layer-by-layer’ growth model of vdWE of the AlN films on MLG. In this growth model, steps of substrate lead to the maximal misorientation and result in nanocolumnar AlN films, and the HAGBs are created on the step edges of the MLG/SiC due to the coalescence of different nuclei (island) at terraces and steps. Therefore, substrate without a miscut angle is critical for vdWE of single crystalline materials on 2D materials, which is different from the conventional epitaxy. As the difference of growth model of conventional epitaxy and vdWE discussed above, the origination of the zigzag cracks of AlN films grown on MLG/SiC becomes distinct. Figure 8b displays the schematic illuminations of cracks of AlN films grown on MLG/SiC and SiC.


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Parallel cracks with only [11-20]SiC direction (also [11-20]AlN direction) can be found in the AlN films grown directly on SiC (as shown in Figure S4 in Supporting Information). M. Rudzinski et al.37, 38 grew GaN films on 4H-SiC with different substrate miscut angles (8° off axis, 3.4° off axis and on axis), and found that [11-20] direction cracks dominated as the miscut angle increased. They considered that the formation of [11-20] direction cracks can be explained by the presence of non-uniformly distributed strain due to the unpaired geometrical partial misfit dislocations (GPMDs) formed along the steps of GaN/SiC interface. For the AlN films grown on MLG/SiC, zigzag cracks can be observed along macroscopical [11-20]SiC direction (also [1100]AlN direction), which is similar to the one grown directly on SiC. The HAGBs are only created along [11-20]SiC direction due to the layer-by-layer growth model of vdWE. Therefore, the unpaired geometrical partial HAGBs may relax the strain along [11-20]SiC direction, resulting the cracks form along [11-20]SiC direction. Nevertheless, the orientation relationship of AlN, MLG and SiC is (0001)[1-100]AlNǁ(0001)[1-100]MLG ǁ(0001)[11-20]SiC and the (1-100)AlN plane is the cleavage plane of the hexagonal crystal system material. Creation of the zigzag cracks along macroscopic [11-20]SiC and microscopic (1-100)AlN plane actually exist, which attribute to the anisotropy of strain after the films cooling down to the room temperature, indicating that the growth model of vdWE is different from the one of conventional epitaxy of 3D martials on 3D substrates. The average cracks density of AlN/SiC is 7.5±2.5x103 cm-1, and the average gap of the cracks is 8±2 nm (supported in Figure S4 in Supporting Information), while the average cracks density of AlN films grown on MLG/SiC is 1.2±0.5x102 cm-1, and the average gap of the cracks is 1000±500 nm (proved in Figure 1a~b, 3c and 6). As the density of the cracks of AlN/MLG/SiC is much low than those of AlN/SiC and the growth temperature is the same, we


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can deduce that the strain of the AlN films during the growth process is different while the thermal mismatch (only related to the growth temperature) is the same (geometrical partial HAGBs or GPMDs relax the strain along only [11-20]SiC direction). Therefore, it suggests that vdWE shows no excessive strain during the growth process owing to the incommensurate inplane lattices at the heterointerface, this may be the direct evidence for the films grown on 2D materials. Although the lattice mismatch strain during the growth can be relaxed by vdWE, the strain originated from thermal mismatch is impossible to overcome. Figure 9 illuminates the Raman spectra of the AlN films before and after release. More clear intensity of A1 (TO), E2 (high), E1 (TO), and A1 (LO) peaks39 can be observed in AlN films after release, and we also calculate the residual strain of the AlN films by E2 (high) frequency shift40 (shown in Table S2 in Supporting Information). The strains of AlN films before and after release are the same (-0.09 GPa), which is far lower than AlN film (-0.66 GPa) grown directly on SiC without MLG at the same run. Therefore, we can conclude that MLG acts as a strain release layer during the AlN films cooling down to room temperature (another evidence is the average gap of cracks in AlN/MLG/SiC about 100 times larger than the one in AlN/SiC). Here, we make further efforts on discussing the strain which was come from thermal mismatch. J. Kim et al.41 reported the layer-resolved graphene transfer via engineered strain layers (ESL), and found the binding energies per atom (γ) between graphene and SiC was 106 meV, and the graphene can be released by ESL all over the SiC with

the condition of

γ(ESL)>γ(graphene-SiC). Since the γ(AlN-graphene) is 67.5 meV discussed above, which is larger than γ(graphene-graphene) (40~50 meV) and smaller than γ(graphene-SiC), the release interface can be taken place only in the interlayers of graphene due to AlN films act as an ESL.


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This is consistent with the results that we detected from the Raman spectra and mapping and may be the reason of no graphene remains on released GaN as monolayer graphene was used17. J. Kim et al.17 also reported 0.5~1 GPa tensile strain for releasing the GaN films form the graphene/SiC. We found the tensile strain of the AlN films grown directly on SiC was about 0.66 GPa, of which is the strain of thermal mismatch dominated. It deduces that the strain of AlN films (AlN/MLG/SiC) during cooling down to room temperature may be higher than 0.5 GPa, as the AlN films released from the graphene partially (as shown in Figure 1a). Therefore, the zigzag cracks formed inevitable due to the tremendous and geometrical partial tensile strain in [1120]AlN direction. From the discussed above, we can deduce the direction for obtaining crack-free, single crystalline AlN films by vdWE. The substrate should be on-axis without a miscut angle for the surface of MLG paralleling to the SiC surface after Si sublimation. The in-plane lattice orientation of MLG should be uniform as the substrate should be Si-terminated SiC. Then, crackfree, single crystalline AlN films will be obtained after vdWE. Here, we show the SEM and TEM images of the crack-free, single crystalline AlN films grown on MLG with Si-terminated on-axis 6H-SiC as a substrate (shown in Figure S5 in Supporting Information). Other characterizations are being performed and will be published in other article. Figure 10 illuminates the UV photosensor of the AlN/MLG heterojunction. Two Pt electrodes were prepared on the surface of AlN/MLG/SiC and MLG/SiC respectively, as depicted in Figure 10a. The I-V curves as shown in Figure 10b display the Schottky heterojunction characterization as the current is larger than zero at the voltage of 0 V, suggesting that it is Schottky contact of AlN and MLG. Clearly photocurrent effect can be seen as the UV light introduced. Notably, no obvious photoresponses is observed by indoor light or white


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radiation from a halogen lamp, indicating that the AlN/MLG Schottky heterojunction is sensitive only to UV light with the photon energy higher than the band gap of AlN. Figure 10c illustrates the time-dependent photocurrent response of the AlN/MLG Schottky heterojunction photosensor measured by turning on and off the UV light when a 15 V bias voltage is applied. Obvious responsiveness can be detected by the UV light turning on and off. One of each case of the timedependent photocurrent response is shown in the Figure 10d. Under illumination, the photocurrent responsiveness increases expeditiously in 0.8 s, and decreases dramatically to the initial value in 1.6 s when the UV light off. Such low rise and decay time indicates the excellent responsiveness of the AlN/MLG Schottky heterojunction photosensor. 3. CONCLUSIONS In conclusion, we have demonstrated a first step in the direction of growth model of vdWE of films using a case of AlN films on MLG/SiC at high temperature. The orientation relationship of AlN and MLG was (0001)[1-100]AlNǁ(0001)[1-100]MLG due to the incommensurate vdWE with a 0.9% in-plane lattice mismatch. Remarkably, the fantastic, interest zigzag cracks formed in the AlN films after cooling down to the room temperature, which reveals the ‘layer-by-layer’ growth model of the vdWE, which is different from the one of conventional epitaxy. Such characteristics only can be described distinctly by vdWE of films rather than nonplanar nanoarchitectures. It provides a guideline for obtaining crack-free, single crystalline planar materials by vdWE. Moreover, the residual strain of the AlN films is nearly free and far smaller than AlN films grown directly on SiC in virtue of the weak vdW forces, and the AlN films can be transferred to foreign substrates by directly mechanical exfoliation without any stressor layer, and shows the practicability of transferable UV photosensor device. 4. METHODS


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Growth of AlN films. MLG was grown on Si-terminated face 4H-SiC of (0001) plane off angle toward axis 4° and toward axis 0° by sublimation (TankeBlue Semiconductor Co., Ltd.). AlN films were fabricated by a home-made HVPE system at the pressure of 50 Torr. The MLG was pretreated in NH3 at 1100 °C for about 20 min to introduce the defects, and then the AlN films were grown for about 1 hour. A similar SiC substrate without MLG and a MLG/6H-SiC of (0001) plane without miscut angle were also used in the same run. Transfer of AlN and TEM sample preparation. The AlN films were transferred onto electroconductive tape by directly mechanical exfoliation without any stressor layer. For the cross-sectional TEM sample, a Pt protective layer was deposited onto the surface of the sample using the ion beams in a dual-beam focused ion beam (FIB) system, the FIB milling voltage was changed from 30 kV to 2 kV in order to minimize the beam-induced damage, a standard copper omniprobe grid was used for fix the FIB lamella. For the plan-view TEM sample, AlN films were transferred directly onto a copper grid with a 1 mm diameter through-hole for TEM characterization. Characterization. The surface morphologies of the AlN films were obtained using a field emission SEM (Hitachi; S-4800) and AFM (Veeco; Dimension 3100), The crystal quality was performed using power and high-resolution XRD (Bruker; D8 Advance and D8 Discover), the microstructural properties were investigated by TEM (FEI; Tecnai G2 F20 S-Twin), Raman spectra were recorded with a laser excitation at 532 nm (JY; LabRam HR 800). First principles calculation. The electronic structures of AlN/MLG heterostructure were studied by first principle calculation based on density function theory using ab initio simulation package. To model the AlN/MLG heterostructure, a slab contains bilayers of graphene followed by two layers AlN initiated with N atoms. The supercell structure is an AlN slab (5 x 5) on a


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bilayer graphene slab (4 x 4) with a vacuum of 15 Å introduced between neighboring slabs to avoid spurious interaction. Device fabrications and measurements. Two Pt electrodes were prepared on the surface of AlN/MLG/SiC and MLG/SiC respectively. The UV photosensor was carried out by Keithley 4200 semiconductor device analyzer system, A 25 W deuterium lamp (Ocean Optics DH-2000S-DUV) (190~400 nm) was used as the UV light source.

ASSOCIATED CONTENT The Supporting Information is available free of charge on the ACS Publications website at DOI: Raman mapping of AlN films after release (Figure S1), HRXRD and surface morphology of the AlN/SiC (Figure S2, S4 and Table S1), SEM images of the AlN nuclei grown on the MLG/SiC (Figure S3), Strain of AlN films (Table S2), SEM and TEM of the AlN/MLG/6H-SiC (Figure S5). AUTHOR INFORMATION Corresponding Author *E-mail: [email protected]; [email protected]. Notes The authors declare no competing financial interest. ACKNOWLEDGMENT This work was supported by National Key R&D Program of China (No. 2016YFA0201101), the China National Funds for Distinguished Young Scientists (No. 61325022), the National Natural Science

Foundation

of

China

(Nos.

61604170,

61574097,

61674164),


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the State Key Program of National Natural Science Foundation of China (No. 61734008), and the National Key Scientific Instrument and Equipment Development Project (No. 11327804).

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Table of Contents

Figure 1. (a) SEM image of AlN films grown on MLG/SiC. (b) Magnified SEM image of the AlN zigzag microbelt. (c) Surface morphology of AlN films. (d) The cross-section SEM image of AlN films. The white arrows represent the direction of SiC. Figure 2. Surface morphology of (a) MLG/SiC after AlN films release and (b) the back surface of AlN films after release. (c) Raman spectra of AlN/MLG/SiC and the surface at the region in (a~b). The white arrows represent the direction of SiC. Red, green, and balk inverted triangles in (c) correspond to AlN, SiC, and graphene, respectively. Figure 3. (a) Power XRD of AlN films grown on MLG/SiC. (b) High resolution XRD 2theta/omega scans. (c) Plan-view TEM image of AlN films after release. (d) SAED patterns of AlN films. Figure 4. (a) Cross-section TEM image of AlN films grown on MLG/SiC. (b) Magnified TEM image in (a). (c) HRTEM image of AlN/MLG/SiC interface. (d~e) HRTEM images at the region marked in (c). SAED patterns of (f) AlN and (g) SiC at the region of (d) and (e). (h) and (i) Experiment and simulated CBED patterns of the AlN films. The white arrows represent the direction of SiC. Figure 5. (a) Top and (b) side view of the first principle calculation of the AlN on graphene. The structure is an AlN slab (4 x 4) on a bilayer graphene slab (5 x 5). The gray, light blue, and navy blue balls are the C, N, and Al atoms, respectively. Figure 6. (a) Surface morphology of the AlN zigzag crack. (b-e) Surface morphology of the AlN zigzag cracks at the regions marked in (a). The white and red arrows represent the direction of SiC and AlN, respectively. Figure 7. HRTEM image of the AlN films grown on MLG/SiC at a step edge. The white and black arrows represent the direction of SiC and AlN, respectively. Figure 8. The schematic illuminations of (a) the difference of the growth model of conventional epitaxy and vdWE, (b) the difference of the cracks feature of AlN/MLG/SiC and AlN/SiC. The black and red arrows represent the direction of SiC and AlN, respectively. 4o is the miscut angle of SiC substrate. Figure 9. Raman spectra of AlN films grown on SiC, MLG/SiC and after release. Figure 10. (a) SEM image of the as-prepared AlN/MLG photosensor. (b) The I-V characteristics of the photosensor with and without UV irradiation. (c) The time response of the photosensor at a fixed bias voltage of 15 V. A deuterium lamp was used as the UV light source with the wavelength of 190 to 400 nm. (d) One of each case of the time-dependent photocurrent response.


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Figure 1. (a) SEM image of AlN films grown on MLG/SiC. (b) Magnified SEM image of the AlN zigzag microbelt. (c) Surface morphology of AlN films. (d) The cross-section SEM image of AlN films. The white arrows represent the direction of SiC.


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Figure 2. Surface morphology of (a) MLG/SiC after AlN films release and (b) the back surface of AlN films after release. (c) Raman spectra of AlN/MLG/SiC and the surface at the region in (a~b). The white arrows represent the direction of SiC. Red, green, and balk inverted triangles in (c) correspond to AlN, SiC, and graphene, respectively.


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Figure 3. (a) Power XRD of AlN films grown on MLG/SiC. (b) High resolution XRD 2theta/omega scans. (c) Plan-view TEM image of AlN films after release. (d) SAED patterns of AlN films.


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Figure 4. (a) Cross-section TEM image of AlN films grown on MLG/SiC. (b) Magnified TEM image in (a). (c) HRTEM image of AlN/MLG/SiC interface. (d~e) HRTEM images at the region marked in (c). SAED patterns of (f) AlN and (g) SiC at the region of (d) and (e). (h) and (i) Experiment and simulated CBED patterns of the AlN films. The white arrows represent the direction of SiC.


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Figure 5. (a) Top and (b) side view of the first principle calculation of the AlN on graphene. The structure is an AlN slab (4 x 4) on a bilayer graphene slab (5 x 5). The gray, light blue, and navy blue balls are the C, N, and Al atoms, respectively.


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Figure 6. (a) Surface morphology of the AlN zigzag crack. (b-e) Surface morphology of the AlN zigzag cracks at the regions marked in (a). The white and red arrows represent the direction of SiC and AlN, respectively.


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Figure 7. HRTEM image of the AlN films grown on MLG/SiC at a step edge. The white and black arrows represent the direction of SiC and AlN, respectively.


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Figure 8. The schematic illuminations of (a) the difference of the growth model of conventional epitaxy and vdWE, (b) the difference of the cracks feature of AlN/MLG/SiC and AlN/SiC. The black and red arrows represent the direction of SiC and AlN, respectively. 4o is the miscut angle of SiC substrate.

Figure 9. Raman spectra of AlN films grown on SiC, MLG/SiC and after release.


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Figure 10. (a) SEM image of the as-prepared AlN/MLG photosensor. (b) The I-V characteristics of the photosensor with and without UV irradiation. (c) The time response of the photosensor at a fixed bias voltage of 15 V. A deuterium lamp was used as the UV light source with the wavelength of 190 to 400 nm. (d) One of each case of the time-dependent photocurrent response.


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Growth Model of van der Waals Epitaxy of Films: A Case of AlN Films

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