Controlled Growth of Unidirectionally Aligned Hexagonal Boron

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Controlled Growth of Unidirectionally Aligned Hexagonal Boron Nitride Domains on Single Crystal Ni (111)/MgO Thin Films Junhua Meng, Bangming Ming, Xingwang Zhang, Menglei Gao, Likun Cheng, Zhigang Yin, Denggui Wang, Xingxing Li, Jingbi You, and Ruzhi Wang Cryst. Growth Des., Just Accepted Manuscript • DOI: 10.1021/acs.cgd.8b01542 • Publication Date (Web): 13 Dec 2018 Downloaded from http://pubs.acs.org on December 14, 2018

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Crystal Growth & Design

Controlled Growth of Unidirectionally Aligned Hexagonal Boron Nitride Domains on Single Crystal Ni (111)/MgO Thin Films Junhua Meng,†,‡ Bangming Ming,§ Xingwang Zhang,*,†,‡ Menglei Gao,†,‡ Likun Cheng,† Zhigang Yin,†,‡ Denggui Wang,†,‡ Xingxing Li,†,‡ Jingbi You,†,‡ and Ruzhi Wang§

†Key

Lab of Semiconductor Materials Science, Institute of Semiconductors, Chinese Academy

of Sciences, Beijing, 100083, P. R. China ‡College

of Materials Science and Opto-electronic Technology, University of Chinese Academy

of Sciences, Beijing 100049, P. R. China §College

of Materials Science and Engineering, Beijing University of Technology, Beijing

100124, P. R. China

KEYWORDS: hexagonal boron nitride, epitaxial growth, single-crystal domains, unidirectional orientation, ion beam sputtering deposition

ACS Paragon Plus Environment

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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

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ABSTRACT: Two-dimensional (2D) hexagonal boron nitride (h-BN) is considered as an ideal dielectric layer or substrate for other 2D heterostructure electronic devices. However, reported 2D h-BN films consist mostly of small sized and randomly oriented h-BN domains, resulting in a high density of grain boundaries during coalescence. Here, we report the growth of unidirectionally aligned h-BN domains in large area on the Ni (111) epitaxial thin film on MgO (111) substrate by ion beam sputtering deposition. It is found that the in-plane orientation of the underlying Ni thin film, which can be controlled by its deposition temperature, plays a key role on the h-BN domain alignment. Furthermore, density functional theory calculations are performed to determine the favorable configuration of the triangular shaped h-BN domains on Ni (111). This work provides a promising approach to prepare unidirectionally aligned h-BN domains in large area, and thus it is possible to achieve wafer-scale single crystal h-BN by stitching these h-BN domains.

INTRODUCTION Hexagonal boron nitride (h-BN), a unique insulating two-dimensional (2D) material, consists of alternating boron and nitride atoms arranged in a honeycomb lattice structure. Due to its outstanding properties such as atomically smooth surface, low density of dangling bonds/trapped charges, wide band gap, high thermal conductivity, and low dielectric constant,1-4 h-BN has been intensively investigated as dielectric or substrate layers for other 2D electronic and optoelectronic devices.5-9 Recently, deep ultraviolet photodetectors based on few-layer h-BN has also been fabricated.10-12 For both device applications and fundamental studies, synthesis of wafer-scale high-quality h-BN layer is extremely important. However, the large-area h-BN


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layers prepared on various transition metals usually exhibit polycrystalline nature with a typical average grain size of several microns or tens of microns.1-4 The limited size and random orientation of h-BN domains result in a high density of grain boundaries during coalescence, which significantly affects its electronic or mechanical properties and thus degrades the performance of devices. The difficulty of preparing high-quality large-area h-BN layers has hindered their widespread use. Similar to graphene, there are two possible approaches for the growth of large area h-BN single-crystals. The first approach is suppressing the nucleation density and growing a single domain as large as possible by optimizing the growth parameters or pretreating the substrate surface.13-23 Along this route, the lateral dimensions of single-crystal h-BN domains have increased from several micrometers to about 600 μm.22,23 However, further enlarging the grain size requires a long growth time and thus becomes more difficult. Another approach is growing unidirectionally aligned h-BN domains and then stitching them together to form uniform single crystal layer without grain boundaries, like the case of wafer-scale growth of single-crystal graphene.24 To this end, many efforts have been devoted to grow aligned h-BN domains on various substrates.25-32 For example, it has been demonstrated that the unidirectionally aligned hBN domains can be grown on Cu (102) or Cu (103) grains,27 however, the average grain size of Cu crystal facets was only about several hundreds of micrometers. Besides, h-BN domains were also synthesized on single crystal substrates such as Ge (110), Pt (111), Ir (111), Co (111) and Ni (111),28-32 nevertheless they frequently took at least two opposite orientations, inevitably leading to the formation of grain boundaries. The origin of the misoriented domains has been tentatively attributed to various factors, including commensurability, symmetry, interaction with the substrate, and also chemistry of the precursor.29,30 Among of them, Ni (111) is an ideal substrate


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for the epitaxial growth of h-BN layers because of its very small lattice mismatch (~0.4%) with h-BN.32,33 In a recent study, we have demonstrated the growth of two oppositely oriented h-BN domains on Ni (111) thin films on c-sapphire by ion beam sputtering deposition (IBSD).22 However, the factors of controlling the initial orientation of h-BN domains are not yet well understood, and much effort is still needed to achieve unidirectional alignment of h-BN domains on metal surfaces. Here, we report the growth of unidirectionally aligned h-BN domains over large area on the Ni (111) epitaxial thin film on MgO (111) substrate by IBSD. The lattice misfit between MgO (111) and Ni (111) was only 0.4% at a supercell matched condition, making MgO (111) a promising candidate for the epitaxial growth of Ni (111) thin films.34,35 More importantly, bulk MgO (111) has a three-fold symmetry similar to Ni (111) surface, resulting in a perfect Ni (111) epitaxial thin film under a suitable substrate temperature, whereas the Ni (111) film on a six-fold symmetrical sapphire (0001) surface exhibits two in-plane orientations (i.e., forming twin crystals). As a result, unidirectionally aligned h-BN single-crystal domains were successfully synthesized on the Ni (111) epitaxial films on MgO (111). Density functional theory (DFT) calculations are performed to determine the favorable configuration of the unidirectionally aligned h-BN domains. This work provides an effective approach for synthesizing h-BN domains with unique orientation in large area. RESULT AND DISCUSSION The growth process of h-BN on epitaxial Ni films is schematically illustrated in Figure S1. High-quality Ni thin films were initially deposited on MgO (111) or c-sapphire substrates by direct current magnetron sputtering, and then the Ni films were used as substrate for the growth of h-BN domains by IBSD (Details are provided in the Experimental Section). Figure 1a, b


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shows typical scanning electron microscopy (SEM) images of the h-BN domains grown on Ni/sapphire and Ni/MgO thin films, respectively. Well-defined triangular shaped h-BN domains were observed on both substrates, as reported previously, implying the single-crystal nature of the h-BN domains.15,22 Interestingly, it can be clearly seen that two oppositely oriented h-BN domains are observed on the Ni/sapphire substrate, while the h-BN domains grown on the Ni/MgO substrate exhibit a unique orientation over the whole scanned area. The atomic force microscopy (AFM) images in Figure S2 displays the as-grown h-BN domains on Ni substrate. Because of the relatively rough Ni surface, the thickness and roughness of the h-BN domain cannot be determined from the height profile of the AFM image (Figure S2a). The h-BN domain can be clearly distinguished from the surrounding Ni substrate in the phase image (Figure S2b), which reflects changes in the mechanical properties of the sample surface. To investigate the chemical composition of the sample, the auger electron spectroscopy (AES) spectra taken in the blue and cyan dotted areas in Figure 1b were recorded (Figure 1e). The B (KLL) and N (KLL) peaks measured on h-BN domains are centered at 177.3 and 386.4 eV, respectively, consistent with the values in previous reports.19 In addition, we can see that the h-BN-coated area (blue dot) shows a much weaker signal of O (KLL) than the uncoated area (cyan dot), demonstrating the high oxidation resistance of the h-BN layer.36 The B (KLL) and N (KLL) Auger electron maps of the h-BN domains on Ni/MgO substrate (Figure 1c, d) indicate that the distributions of both B and N elements are homogeneous in the domains. These results confirmed that the triangular shaped domains consist of B and N atoms. Besides, the B/N atomic ratio from XPS core-level spectra (Figure S3) was calculated to be 1.02, which is very close to the 1:1 stoichiometry of hBN. Raman and UV-vis absorption spectra were also collected after the samples were transferred on SiO2/Si and quartz substrates. As shown in Figure 1f, a characteristic Raman peak at about


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2 0

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0.00 0

200

400

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1200 1300 1400 1500 1600 -1 Raman Shift (cm )

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Figure 1. Characterization of h-BN domains. SEM images of the h-BN domains grown on (a) Ni (111)/sapphire and (b) Ni (111)/MgO substrates. (c,d) The corresponding N (KLL) and B (KLL) Auger electron maps of the h-BN domains on Ni (111)/MgO substrate. (e) AES spectra taken in the blue and cyan dotted areas in (b). (f) Raman spectrum of h-BN domains on a SiO2/Si substrate. (g) UV-vis spectrum of h-BN layer transferred on a transparent quartz substrate. The inset shows optical band gap (5.96 eV) analysis of h-BN from (g). 1371 cm-1 was observed with the full width at half maximum (FWHM) of ~20 cm-1, implying the high crystallinity of the h-BN. Figure 1g shows a typical UV-vis absorption spectrum of the hBN layer, in which the h-BN layer exhibits almost zero absorption in the visible-light range, except for a strong absorption peak at about 200 nm, showing a characteristic of 2D h-BN. By using the formula for a direct band semiconductor, the optical band gap of h-BN was estimated to be 5.96 eV (the inset of Figure 1g), which matches well with the previously reported values.15,25,37 To explore the origin of the orientation dependence of h-BN domains on the underlying substrates, x-ray diffraction (XRD) measurements were carried out for both Ni thin films


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deposited on c-sapphire and MgO (111) substrates after growing h-BN, respectively. The corresponding XRD patterns in a normal -2 configuration are presented in Figure 2a with each Bragg peak assigned to its identification. Besides the peaks from the substrates, only strong (111) and (222) diffraction peaks of the Ni films are clearly observed, suggesting that the Ni (111) planes are oriented parallel to the surface of c-sapphire or MgO (111) substrates. In order to extract the in-plane orientation of the Ni films with respect to the underlying substrates, XRD azimuthal scans were taken on the Ni (200), sapphire (116) and MgO (200) reflections, respectively. As shown in Figure 2b, the -scans of both sapphire (116) and Ni (200) exhibit six sharp peaks at 60° intervals, revealing the six-fold rotational symmetry about the sapphire [0001] or Ni [111] direction. In principle, three peaks at 120° intervals should be expected in the -scan of a perfect single-crystal Ni (111). Herein, the observed six-fold symmetry implies the presence of Ni (111) twin which is rotated by 60° with respect to its original orientation. Schematic of epitaxial relationship between Ni film and sapphire (0001) is shown in Figure S4a. The twin structure arises from different stacking sequences of Ni atom layers (ABC or ACB stacking fault) (Figure S4b).38 In sharp contrast, only three Ni (200) peaks are resolved occurring at azimuthal angles identical with those of MgO (200) reflections, implying that a perfect Ni (111) single-crystal thin film without twin structure is epitaxially grown on the MgO (111) substrate. The high in-plane orientation is explained by relatively perfect super cell matching between the Ni (111) film and the MgO (111) substrate. The in-plane lattice constants of MgO and Ni are 4.21 and 3.5238 Å, respectively, leading to a lattice misfit as high as ~16%. Nevertheless, once considering a (6×6) supercell of Ni to match a (5×5) supercell of MgO, the mismatch can be largely reduced to only 0.4%. Moreover, according to the above XRD results, the epitaxial relationship of Ni/MgO (111) can be determined to be Ni(111)[1-10]//MgO(111)[1-10], as


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Ni (200) MgO (200)

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Figure 2. Characterization of epitaxial Ni (111) thin films. (a) XRD θ-2θ scans of the Ni thin films on sapphire (0001) and MgO (111) substrates. Solid squares (■) and triangles () denote peaks of the MgO and sapphire substrates, respectively. XRD -scans of (b) the Ni (200) and the sapphire (116) reflections and (c) the Ni (200) and the MgO (200) reflections. (d,f) EBSD mapping and (e,g) corresponding {111} pole figure of the Ni thin film on (d,e) sapphire (0001) and (f,g) MgO (111) substrates. Inset of (d, f) shows the color codex of the Ni lattice direction. schematically illustrated in Figure S5a. In contrast to the six-fold symmetry of sapphire (0001), bulk MgO (111) is actually three-fold symmetrical due to its cubic ABC stacking structure along the [111] direction, as demonstrated in -scan of Figure 2c. The long-range forces (i.e. the crystal-lattice potential) arising from the second or third layer of the underlying substrate favor one of the two probable stacking sequences (ABC or ACB) during the growth of Ni layer,34 leading to the epitaxy of Ni (111) single-crystal film on MgO (111). (Figure S5b). To further identify the local crystal orientations of the Ni thin films on different substrates, electron back-scatter diffraction (EBSD) mapping was also performed. As shown in Figure 2d, f,


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the whole surface of Ni films on both c-sapphire and MgO (111) exhibits the same blue color, which indicates that the whole region is oriented in the [111] direction. In the {111} pole figures, six-fold symmetrical patterns at intervals of 60° are observed for the Ni film on c-sapphire (Figure 2e), indicating that the two kinds of Ni (111) grains are rotated by 60°, while it shows a three-fold symmetrical pattern for the Ni/MgO sample (Figure 2g). Both the XRD and EBSD studies on the Ni (111) films strongly suggest that a perfect Ni (111) single-crystal film is epitaxially grown on MgO (111) while the Ni (111) twin grains are formed on c-sapphire substrate. Based on the above results, we suppose that the underlying Ni twins on c-sapphire are responsible for the formation of two oppositely oriented h-BN domains. The unidirectionally aligned h-BN domains can be obtained on a perfect Ni (111) single-crystal film. Obviously, the crystalline quality of underlying Ni film is critical for the orientation of h-BN domains. In this work, we find that the substrate temperature plays a key role on the in-plane orientation of Ni (111) films. Figure 3a shows XRD -scans of the Ni/MgO thin films deposited at 500, 650 and 800 °C. A total of six Ni (200) peaks were clearly resolved for the 500 °C-grown Ni film. Three of them occur at azimuthal angles identical with those of MgO (200) (marked by square symbols, type I), while the other three weak peaks marked by star symbols (type II) are shifted 60° from the substrate peaks. Here, the observed six diffraction peaks indicate the existence of two sets of Ni crystallites (i.e., Ni (111) twin), similar to the case of Ni on c-sapphire substrate. As the substrate temperature increases to 650 °C, the three peaks from type II disappear and the -scan shows a dominant single-crystal structure, as discussed above. At 800 °C, a small amount of Ni (111) twin structure reappears, as evidenced by its weak diffraction peaks from type II. Temperature dependence of the in-plane orientation of Ni (111) film can be interpreted as follow.


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(a)

(b)

Intensity (a.u.)

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o

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Figure 3. Correlation between the orientation of h-BN domains and the underlying Ni film. (a) XRD -scans of the Ni (111)/MgO thin films deposited at various temperatures. Open squares () and stars () denote peaks from type I and type II Ni (111) twin, respectively. SEM images of the h-BN domains grown on the Ni (111)/MgO films which were deposited at (b) 500, (c) 650 and (d) 800 °C. The white arrows in (b), (c), and (d) point the 0° orientation of h-BN domains, while the white dashed circles in (d) indicate the oppositely oriented h-BN domains (i.e., 180° oriented h-BN domains). (e) Statistical distribution of the orientations of h-BN domains and the ratio of type I/type II diffraction peak in -scans for the samples shown (b)-(d). At low temperatures, both the energy and mobility of the adatoms are low, the Ni atoms will bind to both two possible positions in spite of their different energies, thus producing a twin


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structure. With the increase of temperature, the energy and mobility of the adatoms are high enough, and the adatoms can always find the lowest energy position, resulting in a single-crystal structure.34 At elevated temperatures, the surface diffusion rate is much enhanced and thus the deposited atoms have a few chances to be captured in the energetically metastable position, leading to weak diffraction peaks from twin crystals. To confirm the correlation between the orientation of h-BN domains and the underlying Ni film, the h-BN domains were synthesized on the Ni (111)/MgO films grown at different temperatures. Figure 3b-d shows SEM images of the h-BN domains grown on the Ni (111)/MgO films which were deposited at 500 °C, 650 °C and 800 °C, respectively. As expected, the h-BN triangles synthesized on the 500 °C-grown Ni film show two preferred orientations opposite to each other, and the numbers of two oppositely oriented h-BN domains are very close. On the contrary, almost all of the h-BN triangles on the 650 °C-grown Ni film are aligned along the same orientation over the whole scanned area, while the few oppositely oriented h-BN domains can be occasionally observed on the 800 °C-grown Ni film, as indicated by white dashed circles in Figure 3d. For easily comparison, both the statistical distribution of the orientations of h-BN domains and the ratio of type I/type II diffraction peak in -scans are shown in Figure 3e. Apparently, the proportion of the two oppositely oriented h-BN domains is highly consistent with the ratio of two sets of Ni twin crystal, further confirming that the h-BN domain alignment rely on the in-plane orientation of the underlying Ni film. According to the previous reports, both growth mechanisms, surface-mediated reaction and gas phase nucleation/growth, are responsible for the growth of h-BN on Ni, depending on the growth rate.15,22 The positions of B and N atoms with respect to the Ni-surface sites have been theoretically and experimentally investigated.22,39,40 The N-terminated zigzag edge is expected


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Formation Energy (eV)


-8.22

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-8.22

-8.24

Ni

-8.26 -8.26 top-hcp

B N

top-fcc

Figure 4. Structure models of two possible orientations of h-BN and calculated formation energy of top-hcp (-8.22 eV) and top-fcc (-8.26 eV) configurations. for the triangular shaped h-BN domains due to the lower edge energy,39 and N atoms preferentially sit on the top of Ni atoms because of the strong hybridization of N lone pairs with Ni atoms.40 Consequently, the arrangement of the adsorbed B atoms on the underlying Ni (111) substrate will determine the orientation of h-BN domains. Structure models of two possible orientations: N atoms sitting on the top of Ni atoms, and B atoms sitting either on fcc (N, B)=(top, fcc) or hcp (N, B)=(top, hcp) hollow sites, are shown in the inset of Figure 4, respectively. Herein, the unidirectionally aligned h-BN domains can be realized on the Ni (111) single-crystal film, indicating that there exists a preferred orientation of h-BN during nucleation. To determine which configuration (fcc or hcp) is dominant on a perfect single-crystal Ni (111), the formation energies of both fcc and hcp configuration were theoretically calculated by DFT methods as implemented in the Vienna Ab-initio Simulation Package (VASP). A primitive cell of 1×4×4 Ni (111) is constructed as substrate and the edges of h-BN nuclei in the models are passivated by hydrogen atoms. The formation energy is given by the formula: Ef=(EBN,Ni(111)ENi(111)-EBN)/n, where EBN,Ni(111), ENi(111) and EBN are the total energy of BN/Ni (111), Ni (111), and BN, and n is the number of BN units in BN supercell. The computational results reveal that


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the formation energies of hcp and fcc configuration are -8.22 and -8.26 eV, respectively, as shown in Figure 4. Clearly, the hcp configuration is energetically less favorable (40 meV/BN) than the fcc structure, implying the h-BN with fcc configuration is preferable on the perfect single-crystal Ni (111) substrae. Additionally, the average size of h-BN domain on the Ni/MgO substrate (less than 100 μm) is much smaller than that on the Ni/sapphire substrate (about 200 μm), which is probably related to the difference in surface morphology of underlying Ni films. Figure 5a, b shows the AFM height images of the 650 °C-grown Ni/sapphire and 650 °C-grown Ni/MgO film after annealing in vacuum at 1050 °C for 20 min. It can be clearly seen that the Ni film deposited on c-sapphire substrate has a rather smooth surface with a root-mean-square (RMS) roughness, Rq of 0.96 nm (Figure 5a). However, some ripples and corrugations could be observed on the 650 °C-grown Ni/MgO film with the higher Rq of 1.42 nm (Figure 5b). In particular, the 800 °C-grown Ni/MgO (a)

RMS Roughness of Ni (nm)

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Figure 5. AFM height images of (a) the 650 °C Ni (111)/sapphire, (b) the 650 °C Ni (111)/MgO, and (c) the 800 °C Ni (111)/MgO film after annealing in vacuum at 1050 °C for 20 min. (d) Average lateral size of h-BN domains and RMS roughness of the Ni film on different substrates.


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exhibits a rather rough surface with Rq of 2.75 nm, as shown in Figure 5c. The statistical distribution of the average lateral sizes of h-BN domains and the Rq values of Ni films are shown in Figure 5d. As reported previously, the rough surfaces (ripples and corrugations) are likely to act as nucleation seeds, leading to an enhanced rate of nucleation and thus the smaller size of hBN domain.13-15 It should be pointed out that the original surface of MgO (111) is rougher than that of sapphire (0001), as shown in Figure S6, which may be the main reason that cause a relatively rough surface of the Ni/MgO film. CONCLUSION In summary, we have demonstrated the controlled growth of unidirectionally aligned h-BN domains on the Ni (111)/MgO thin films over large area by IBSD. The formation of unidirectionally aligned h-BN domains is mainly attributed to the unique in-plane orientation of the epitaxial Ni (111) thin film on three-fold symmetrical MgO (111). The orientation of Ni (111) film on MgO (111) can be controlled by its deposition temperature. DFT calculations indicate that the fcc site is a favorable configuration of the triangular shaped h-BN domains on Ni (111). Furthermore, it is also important to keep the surface of Ni film smooth to maximize the h-BN domain size. This work provides an effective approach for synthesizing h-BN domains with unique orientation in large area, paving a promising route to prepare wafer-scale single crystal h-BN. EXPERIMENTAL SECTION Heteroepitaxy of Ni film. The Ni thin films were deposited by direct current sputtering on single crystal MgO (111) or c-sapphire substrates. Prior to the epitaxial growth, the substrates were successively cleaned for 10 min in each of the following solvents: acetone, isopropanol and ethanol for removing contaminants, and then loaded into the sputtering chamber. The sputtering


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chamber was pumped to a base pressure of 3×10−5 Pa, and subsequently filled with Ar to a working pressure of 1.0 Pa. The Ni thin films with a thickness of about 1.6 μm were grown at different temperatures for 60 min. Preparation of h-BN by IBSD. The pre-grown Ni/MgO or Ni/sapphire substrates were directly loaded into the IBSD chamber equipped with a Kaufmann ion source for the growth of h-BN. When the chamber was evacuated to the base pressure (2×10−5 Pa), the Ni thin films were annealed at 1050 °C for 10 min under 20 sccm H2 atmosphere to obtain a clean and smooth surface. Subsequently, the boron and nitrogen species were sputtered from a pure h-BN target by the Ar ion beam (1.0 keV, 0.2 mA/cm-2) at a constant pressure of 3×10−2 Pa. After the growth procedure, the ion source was shut down and the sample was cooled down to room temperature in pure Ar atmosphere. The thickness of h-BN was controlled by adjusting the ion beam density and the growth time. Finally, the h-BN was transferred on different substrates for characterization with the same method for h-BN on Ni foils, as previously reported.22 Characterization. The surface morphologies of samples were characterized by SEM (FEI Quanta-450) and AFM (NTMDT Solver P47, tapping mode). XRD spectra were recorded by Rigaku D/MAX-2500 system using Cu Kα as the X-ray source. Raman spectra were acquired with a confocal spectrometer (Renishaw Model in Via-Reflex) using a 532 nm laser as the excitation source. Optical absorption spectra of the samples were measured by a Varian Cary 5000 UV-Vis spectrophotometer in a double-beam mode. AES and EBSD measurements were performed on a scanning Auger nanoprobe (PHI 710, Ulvac-Phi). XPS measurements were carried out on an ESCALAB 250Xi instrument with a monochromated Al Kα source. Density Functional Theory Calculations. First-principles calculations were performed based on the framework of DFT as implemented in the VASP. The projector-augmented-wave


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(PAW) method was exploited to describe the interaction between ionic cores and valence electrons, while the generalized gradient approximation (GGA) with the function of Local Density Approximate (LDA) was employed to describe the exchange-correlation interactions. A supercell of 1×4×4 Ni (111) is constructed as substrate and the Van der Waals interaction had been considered between Ni (111) and h-BN. The plane wave expansion setting with an energy cut-off of 520 eV was used in the calculations. All the structures were relaxed, until the forces was less than 3×10−2 eV/Å and the energy was less than 3×10−4 eV between two consecutive self-consistent steps.

ASSOCIATED CONTENT Supporting Information The Supporting Information is available free of charge on the ACS Publications website. Schematic diagrams of growth strategy, epitaxial relationships and additional AFM, XPS characterizations. AUTHOR INFORMATION Corresponding Author *E-mail: [email protected] Notes The authors declare no competing financial interest. ACKNOWLEDGMENT This work was financially supported by the National Key Research and Development Program of China (No. 2018YFB0406503), the National Natural Science Foundation of China (No.


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61674137), the Beijing Natural Science Foundation (No. 4184101), and the Science and Technology Program of Beijing (No. Z181100004418007), and the China Postdoctoral Science Foundation (No. 2017M620873, 2018T110129).

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For Table of Contents Use Only Controlled Growth of Unidirectionally Aligned Hexagonal Boron Nitride Domains on Single Crystal Ni (111)/MgO Thin Films Junhua Meng, Bangming Ming, Xingwang Zhang,* Menglei Gao, Likun Cheng, Zhigang Yin, Denggui Wang, Xingxing Li, Jingbi You, and Ruzhi Wang

The unidirectionally aligned h-BN single-crystal domains over large area are epitaxially grown on the Ni (111)/MgO films by ion beam sputtering deposition. It is found that the in-plane orientation of the underlying Ni thin film, which can be controlled by its deposition temperature, plays a key role on the h-BN domain alignment.


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Controlled Growth of Unidirectionally Aligned Hexagonal Boron

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