Growth and low-energy electron microscopy characterizations of graphene and hexagonal boron nitride

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ScienceDirect Progress in Crystal Growth and Characterization of Materials 62 (2016) 155–176 www.elsevier.com/locate/pcrysgrow

Review

Growth and low-energy electron microscopy characterizations of graphene and hexagonal boron nitride H. Hibino a,b,*, S. Wang b, C.M. Orofeo b,1, H. Kageshima c b

a Kwansei Gakuin University, Sanda, Hyogo 669-1337, Japan NTT Basic Research Laboratories, NTT Corporation, Atsugi, Kanagawa 243-0198, Japan c Shimane University, Matsue, Shimane 690-8504, Japan

Available online 1 June 2016

Abstract Graphene and related two-dimensional (2D) materials are attracting huge attention due to their wide-range potential applications. Because large-scale, high-quality 2D crystals are prerequisites for many of the applications, crystal growth of 2D materials has been intensively studied. We have also been conducting research to understand the growth mechanism of 2D materials and have been developing growth techniques of high-quality materials based on the understandings, in which detailed structural characterizations using low-energy electron microscopy (LEEM) have played essential roles. In this paper, we explain the principles of obtaining various structural features using LEEM, and then we review the status of our current understanding on the growth of graphene and hexagonal boron nitride. © 2016 Elsevier Ltd. All rights reserved. Keywords: graphene; hexagonal boron nitride; two-dimensional materials; low-energy electron microscopy; chemical vapor deposition

1. Introduction 1.1. Graphene Graphene is a two-dimensional (2D) crystal of carbon atoms arranged in a honeycomb lattice. A three-dimensional (3D) stack of graphene sheets is graphite, which is familiar in our daily life. Graphene is also a building block of one-dimension-like carbon nanotubes and zero-dimensionlike fullerene. Theoretical investigations of graphene’s physical properties go back to 1947 [1], but its experimental history is relatively new. In 2004, graphene was first discovered among thin graphite flakes mechanically exfoliated * Corresponding author. Kwansei Gakuin University, Sanda, Hyogo 669-1337, Japan. Tel.: +81 79 565 9743; fax: +81 79 565 9729. E-mail address: [email protected] (H. Hibino). 1 Present address: Department of Physics, National University of Singapore, Singapore 119077, Singapore. http://dx.doi.org/10.1016/j.pcrysgrow.2016.04.008 0960-8974/© 2016 Elsevier Ltd. All rights reserved.

from bulk graphite [2]. Since then, owing to its novelty in science and expectations to a wide range of applications, graphene research has been expanding exponentially. Graphene contains two carbon atoms at A and B sublattices in a unit cell. Three out of four valence electrons of each carbon atom are used to form strong σ bonds with its three neighboring carbon atoms. The remaining pz orbital overlaps with the neighboring atoms to form the π bond. These π electrons govern the electrical properties of graphene. Their low-energy dynamics is described by the Dirac equation with zero mass. The π and π* bands show conical energy dispersions called the Dirac cones and make contact at the corners of the hexagonal Brillouin zone, K and K’ points. The contact point is called the Dirac point and coincides with the Fermi level in neutral graphene. Conduction electrons in graphene constitute a new class of 2D electron system, which behaves truly differently from the 2D electron gas in the heterostructures of 3D semiconductors such as Si or GaAs.

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Graphene has excellent properties in many aspects, such as mechanical, electrical, optical, and thermal properties [3]. Graphene is strong and flexible. The specific strength (tensile strength/density) of 48,000 kN m kg−1 [4] is more than 100 times larger than that of steel. The crystal structure does not break even after being stretched by 20% [5]. Electrically, the carrier mobility of graphene at room temperature, ~ 200,000 cm2 V−1 s−1 [6], is more than 100 times larger than that of Si. Optically, graphene is almost transparent and absorbs only 2.3% of light in the visible to infrared region [7]. However, this absorption coefficient is one to three orders of magnitude larger than those of normal semiconductor materials [8]. Graphene has a thermal conductivity of up to ~5300 W m−1 K−1 at room temperature [9], which is 10-fold or more of Cu and even larger than diamond. 1.2. Hexagonal boron nitride (hBN) The experimental discovery of graphene has renewed interest in atomic layers of the other layered materials such as insulating hexagonal boron nitride (hBN), semiconducting transition metal dichalcogenides (TMDCs), and semiconducting black phosphorous. These layered materials are bonded between layers via the van der Waals forces, which are much weaker than the covalent bonding within a layer. Therefore, they can be thinned to atomic layers by mechanical exfoliation. Single layer hBN has a honeycomb lattice structure, containing boron and nitrogen atoms at A and B sublattices. hBN is an insulator with band gap of 6 eV [10]. A major reason why hBN has attracted attention as a 2D material is its usefulness as substrates of graphene devices [11]. A flat inert hBN substrate prevents charged impurities that degrade the transport properties of graphene from being trapped at the graphene/hBN interface. Compared to the standard graphene device substrate of SiO2, much larger carrier mobility is obtained when hBN is used. Recently, in addition to individual 2D materials, van der Waals heterostructures fabricated by artificially stacking different 2D materials have attracted intense attention from expectations toward new physics and new functions [12]. Heterostructures are important components to enhance the performance and functionality of electronic and photonic devices. However, defects are easily introduced into the heterostructures of 3D materials when they are different in the crystal structure, bond type, lattice constant, and so on. Therefore, combinations of 3D materials available for fabricating high-quality heterostructures are limited. However, the weak bonding between 2D materials allows arbitrary heterostructures to be produced without defects by

mechanical transfer or crystal growth. Therefore, there are huge expectations that new physics and new applications are emerging from the van der Waals heterostructures, in which insulating hBN plays essential roles as dielectric layers and tunnel barriers [13]. 1.3. Applications of graphene Among a wide range of potential applications of graphene, flexible transparent electrodes, which can be used in touch screens, solar cells, light-emitting devices, and so on, are currently the most promising. Severallayer-thick graphene films are transparent, electrically conductive, and flexible. 30-inch-wide graphene films with a sheet resistance of ~30 Ω/sq and a light transmittance of 90% have already been demonstrated [14]. Compared to indium tin oxide (ITO), a typical transparent electrode material, graphene can sustain higher conductivity at the same light transmittance, and can offer flexibility unachievable with brittle ITO films. Graphene is also suitable for high speed electronic applications owing to its high carrier mobility at room temperature, and is anticipated as a transistor channel material in the post-silicon era. However, because intrinsic graphene has no band gap, it is not possible to turn off the current. So, research on the analog high-frequency devices is in progress [15–17]. At the same time, efforts to create a band gap in graphene are actively being made. For example, graphene nanoribbons (GNRs) with the width of ~1 nm have a band gap of an order of 1 eV due to the quantum confinement [18–20]. Until now, GNRs have been prepared by various methods [21–25]. In addition, a band gap can be created in graphene by breaking its symmetry. Application of a vertical electric field to bilayer graphene [26] and chemical modifications of graphene [27] are examples of the symmetry breaking. However, for both the GNRs and symmetry breaking, it is still quite challenging to induce a large enough band gap for digital electronics applications while maintaining graphene’s high mobility. Strong interactions with light of wide wavelengths from THz to visible range as well as high carrier mobility also make graphene promising for photonic devices such as optical detector and modulator [28]. In addition, because graphene has a high surface area, the physical properties can quickly respond to changes in the environment, which leads to high expectations for the sensor applications [29]. High surface area is also advantageous for the energy applications such as capacitors and batteries. Moreover, graphene is impermeable to atoms and molecules [30], and is therefore promising as a corrosion resistant coating. Lastly, graphene is also useful

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for structural, electrical, and thermal reinforcements due to its excellent mechanical strength, and electrical and thermal conductivities. 1.4. Production methods of graphene and hBN Low-cost production methods are essential toward industrial applications of 2D materials. Mechanical exfoliation from bulk materials using adhesive tapes can provide high-quality 2D sheets and their heterostructures. This method is therefore ideal for basic research, but in terms of productivity and size, it is not suitable for industrial applications. Solution-based exfoliation methods, which include, as for graphene, oxidative exfoliation by acid treatment (graphene oxide), exfoliation through intercalation with alkali metals, and direct exfoliation by sonication in organic solvents, can produce flakes of 2D materials in large quantities and at low cost [31], and are promising for applications such as battery, composite materials, and conductive ink. However, high-performance device applications require large-scale, high-quality 2D sheets. Therefore, methods of growing 2D materials on substrates are attracting intense interest. As for graphene, two types of growth methods have been investigated intensively. One is chemical vapor deposition (CVD) on metals, and the other is thermal decomposition of SiC. CVD is also promising as synthesis techniques for hBN and TMDCs and their CVD growth are being researched actively. During CVD growth of graphene, carbon-containing gas such as methane is decomposed with the catalytic action of metal substrates. CVD growth of graphene on metals, including, in a broad sense, segregation and precipitation of graphene films [32], has a long research history, and various metal substrates have been investigated [33–35]. It is known that graphene growth process differs greatly depending on the carbon solid solubility in metal [36]. Comparing the carbon solubility at 900 °C between Ni and Cu, two standard substrates for graphene CVD, the carbon solubility in Ni is ~0.9 at. %, whereas Cu has much smaller solubility of ~5 at. ppm [37]. On a substrate with low carbon solubility, such as Cu, the graphene growth processes are restricted on the surface. The catalytic activity is lost when the surface is covered with monolayer graphene, resulting in the self-termination of the growth at monolayer. The biggest advantage of using Cu for the substrate is that monolayer graphene is obtained with good reproducibility. On the other hand, on a substrate with high carbon solubility such as Ni and Co, decomposed carbon atoms diffuse into the bulk at the growth temperature, and graphene is grown by the segregation/precipitation of carbon atoms during cooling.

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Therefore, the number of graphene layers sensitively depends on the growth parameters such as the growth time and the cooling rate, and is thus difficult to control. Graphene growth by thermal decomposition of SiC is also not a new idea in the surface science field [38]. When heating the SiC substrate in a gas or a vacuum environment, silicon atoms selectively sublime, and the carbon atoms left on the surface spontaneously form graphene. SiC has many polytypes such as 6H, 4H, and 3C, but the influence of polymorphism on the graphene growth is limited. Meanwhile, the structure of the grown graphene is strongly dependent on the surface orientation of the SiC substrate. Two polar surfaces, SiC(0001) and SiC(000 1 ) , have been intensively studied for graphene growth. Bulk-terminated SiC(0001) and SiC(000 1 ) surfaces have Si and C atoms on top and are usually called Si and C faces, respectively. Basically, whereas few-layer graphene grows epitaxially on the Si face, multilayer graphene including in-plane and outof-plane rotational disorders grows on the C face. When the above two growth methods of graphene are compared, higher prices of SiC wafers than metal foils would limit the practical application of the thermal decomposition of SiC. CVD on inexpensive Cu foils enables to grow large-scale monolayer graphene reproducibly using relatively simple instruments like the tube furnace connected to the gas supply system. Therefore, CVD growth of graphene on the Cu foil is regarded as the most industrially promising graphene synthesis method based on its simplicity, scalability and cost efficiency. However, because metal substrates used in CVD growth are conducting, it is necessary to transfer CVD-graphene onto other insulating substrates for device applications [39]. Transfer of large-area graphene has been mainly achieved through the following steps: polymer coating as a temporary rigid support, etching the Cu substrate in an etchant, scooping the polymer/graphene on an insulating substrate, and removing the polymer. Graphene is easily damaged and contaminated during the transfer process. Because of this, damage- and contamination-free, as well as etch-free transfer methods are currently hot topics. On the other hand, because SiC is a wide-band-gap semiconductor, graphene electronic devices can be fabricated directly on SiC without transfer [15–17,40]. hBN has long been the subject in surface science. Already in the 1990s, monolayer hBN and graphene/ hBN heterostructure have been grown on single-crystal metals in ultrahigh vacuum (UHV) by depositing borazine (B3N3H6) [34,41]. Currently, however, methods of obtaining high-quality hBN films for van der Waals heterostructures are limited to the mechanical exfoliation from bulk hBN crystal synthesized by the

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high-temperature high-pressure method [10]. hBN plays important roles as substrates and tunnel barriers in exploring physical properties of 2D materials and developing their device applications. The growth technology of large-scale, high-quality hBN atomic layers is highly desirable. Recently, there are many research papers on CVD growth of hBN, in which solid ammonia borane (BN3-NH3) is more popular as a source material than borazine because it is safer and easier to handle [42]. Nitrogen has a low solid solubility in many metals. Therefore, from the analogy of graphene growth on Cu, CVD growth of hBN is expected to be self-limiting at monolayer. 2. Structural characterizations of 2D materials using low-energy electron microscopy Low-energy electron microscopy (LEEM) is a projection-type microscopy technique collecting lowenergy (typically 1–10 eV) electrons back-scattered from the samples for imaging [43]. Nanometer-scale spatial resolution and video-rate temporal resolution are compatible by LEEM. By changing the lens intensity and using selected-area apertures, low-energy electron diffraction (LEED) patterns in a ~1 μm area can also be obtained. LEEM is a powerful tool for investigating the growth mechanism of graphene and related 2D materials and optimizing their growth conditions. The advantages of LEEM are summarized in the following four aspects. First, LEEM can be used to in-situ observe deposition and segregation/precipitation of graphene on metal surfaces [44]. Second, selected-area LEED patterns can be used to determine the crystallographic orientation locally, and dark-field (DF) imaging using diffracted beams can provide crystallographic orientation maps of polycrystalline 2D materials [45]. Third, number of graphene or hBN layers can be digitally determined from the low-energy electron reflectivity spectra [46,47]. Last, the stacking structure in 2D materials can be determined by analyzing the energy dependence of the LEED intensities (I–V curves) using dynamical calculations [48]. These issues will be explained in detail in the following sections. 2.1. Dynamical observations of graphene segregation LEEM has provided important information about elementary processes in the CVD growth and segregation/ precipitation of graphene. On metal substrates with high carbon solubility like Ni, hydrocarbon molecules are decomposed into carbon atoms, and the carbon atoms diffuse on the surface or dissolve into the bulk [49]. On Cu, on

the other hand, active carbon species contributing to the graphene growth have not been fully clarified yet [49]. According to the in-situ LEEM observations of graphene growth on Ru and Ir, whose carbon solubilities are respectively about 0.1 and 0.03 at. % at 820 °C [50], after the supply of carbon atoms onto the substrate by ethylene or atomic carbon deposition, the coverage of carbon adatoms rises up to a few percent of monolayer, and graphene nucleates when the coverage exceeds a certain critical value [51,52]. On the other hand, the surface carbon concentration during atomic carbon deposition on Cu is always estimated to be less than 10−3 monolayer [53]. Regardless of the metal species, graphene often nucleates heterogeneously at the impurities and steps on the substrates [54]. Even on a single crystal substrate, graphene nuclei do not always have the same crystal orientation [55]. Thus, removal of the impurities and defects is essential to improve the crystallinity of graphene. Attachment and detachment of carbon atoms are repeated at the edges of the graphene islands on the substrates, and both are balanced at the equilibrium surface carbon concentration. When the actual surface carbon concentration is greater than equilibrium, the attachment exceeds the detachment, resulting in growth. Difference of the actual concentration from the equilibrium concentration is the supersaturation. The rate-limiting process of graphene growth depends on the metal species. It has been reported using LEEM that the attachment/detachment of carbon atoms at the graphene edge is a rate-limiting step of the graphene growth on Ru and Ir [51,52,56]. The graphene growth rate is proportional to a 4th to 5th power of the supersaturation. This means that 4–5 carbon atoms attach to or detach from the graphene edge as a unit, in which the binding of the graphene edge to the metal plays a critical role [57]. On the other hand, the surface diffusion of the carbon species, instead of the attachment/ detachment at the edge, is rate-determining in the graphene growth on Cu by atomic carbon deposition in UHV [53]. The diffusion-limited kinetics of graphene growth on Cu often causes dendritic-shaped graphene islands. CVD graphene growth on Cu is self-limiting at monolayer because the catalytic activity of the substrate is lost after full coverage of graphene. Therefore, it is rather difficult to grow few-layer graphene on Cu by CVD. Because few-layer graphene has various potential applications depending on the thickness, it is worth considering the possibility of growing few-layer graphene on metal substrates with high carbon solubility. For this purpose, it is critical to understand the segregation/precipitation mechanism of graphene on metal surfaces. It has been shown, based on the in-situ LEEM investigations on Ru, three states of carbon atoms are

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involved in the graphene segregation/precipitation [58]. The three states are carbon atoms in graphene, carbon atoms on the surface (adatoms), and carbon atoms in the bulk metal. With decreasing substrate temperature, the carbon adatom concentration in equilibrium with carbon atoms in the bulk increases because the carbon solid solubility decreases. When the adatom concentration in equilibrium with carbon atoms in the bulk is larger (lower) than the adatom concentration in equilibrium with graphene carbon atoms, graphene grows (shrinks). The adatom concentration in equilibrium with graphene is larger at higher temperatures. Therefore, during cooling, graphene appears at a critical temperature where these two carbon adatom concentrations, in equilibrium with carbon atoms in the bulk and graphene, are equal. The above LEEM results were obtained using bulk metal single crystal. Polycrystalline foils are more practical as the substrates for graphene segregation/ precipitation, and we investigated the growth processes on polycrystalline Ni foils in-situ using LEEM [44]. Based on the pioneering work on the segregation/precipitation of graphene/graphite on Ni [32], when the Ni substrate in which carbon atoms are dissolved beforehand is gradually cooled, monolayer graphene first appears and exists stably at a certain temperature range. On a Ni(111) substrate with the solid solution of ~0.26 at. % carbon, monolayer graphene is stable approximately between 1180 K and 1065 K, and graphite (multilayer graphene) is formed at lower temperatures. The grain boundaries in polycrystalline graphene are known to degrade the electronic and mechanical properties [59,60]. On polycrystalline metal substrates, the resulting graphene is inevitably polycrystalline because graphene islands nucleated on different metal grains should have different crystal orientations. However, if graphene grows from a single nucleus across the metal grain boundaries, a large single crystal can be obtained. Graphene tends to nucleate heterogeneously at the impurities and roughness on the substrate. In our experiment, therefore, to reduce the nucleation density of graphene, polycrystalline Ni foils were carefully polished up to the roughness below 5 nm. Multilayer graphene was first grown on this ultra-flat Ni substrate by CVD. The sample was heated in the LEEM instrument to high temperatures, until the multilayer graphene was completely dissolved into Ni, and was then cooled to form graphene. Fig. 1 is the result of the in-situ LEEM observations of graphene segregation on the ultra-flat Ni foil. The Ni grain boundaries in the Ni substrate were identified through the brightness difference between the grains of different surface orientations and/or linear contrasts due to the undulation, as shown by the broken lines

(a)

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(b) graphene

Ni surface

10 m t =0.27 s

t =0 s

(c)

(d) f g e t =0.50 s

(e)

(1,0)

t =0.73 s

(f)

(g)

(0,0)

Fig. 1. (a)–(d) LEEM images of polycrystalline Ni foil during segregation of monolayer graphene. The measurement times were indicated in the figures. The electron beam energy was 2.5 eV. The dashed lines mark the positions of the Ni grain boundaries. (e)–(g) LEED patterns measured at the corresponding positions in (d). The patterns were obtained after cooling until multilayer graphene was precipitated. The electron beam energy was 45 eV.

in Fig. 1(a). Graphene domains are distinguished from the Ni substrate by high contrast due to the large difference in the electron reflectivity. The graphene grew continuously beyond the Ni grain boundaries. To check the single crystallinity, selected-area LEED patterns were measured on the different Ni grains. However, the LEED patterns from monolayer graphene were too diffuse for this purpose. Therefore, the selected-area LEED patterns in Fig. 1(e–g) were obtained at the corresponding positions in Fig. 1(d), after cooling until multilayer graphene was precipitated. Because new graphene sheets always precipitate at the graphene/metal interface, the crystal orientation of the initial monolayer can be determined even after multilayer growth. The LEED patterns indicate that graphene keeps its single-crystal orientation when it grows across the Ni grain boundaries. By segregation/precipitation of graphene on polycrystalline Ni foils with limited roughness, graphene grew

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from a small number of nuclei, across the Ni grain boundaries, and to a large single-crystal grain at least ~100 μm in our experiments. We have also succeeded in preparing large bilayer graphene as wide as several tens micrometers. However, after the substrate temperature was further lowered to room temperature, multilayer graphene was formed. Therefore, controlling the amount of the precipitating carbon by the growth parameters such as metal species and thickness of the substrate is essential for growth of uniform few-layer graphene.

2.2. Mapping of crystallographic orientations in polycrystalline graphene CVD on Cu foils is a scalable, cost-efficient synthesis method of graphene, but normally provides polycrystalline graphene. For high-performance device applications, we need to control the grain structures in graphene, which requires a facile method of visualizing the grain boundaries. LEEM and LEED are useful for investigating the grain structures. Graphene grains with different in-plane orientations provide specular (0,0) spots at the same position in the LEED patterns. However, the diffracted beams from the two grains appear at different positions due to their relative rotation. Therefore, the rotation angle is determined by the selected-area LEED patterns. Besides, DF LEEM images using one of the diffracted spots selectively image the corresponding grain. From the DF LEEM images obtained using all the diffracted spots of an identical type, a map of the crystallographic orientation can be reconstructed. Similar orientation maps have been obtained using transmission electron microscopy (TEM) [61]. LEEM lacks atomic resolutions available with TEM, but has an advantage in the sample preparation. The plan-view TEM observations require graphene to be transferred onto TEM grids, but LEEM can image graphene on the conductive substrate without transfer. Nevertheless, LEEM is a UHV-based technique and normally requires several hours to remove the airborne adsorbates before the measurements. Therefore, we developed an easier in-air method of visualizing the grain boundaries based on Raman spectroscopy, one of the most standard methods of characterizing graphene [62]. Graphene can be identified by the G band (typically at ~1580 cm−1) and 2D (G’) band (typically at ~2700 cm−1) in the Raman spectrum. Their positions, shapes, and intensity ratio provide fruitful information about graphene, such as the carrier concentration, strain, and thickness. The Raman D band (typically at ~1350 cm−1) is originated from defects, and the

Fig. 2. (a) BF LEEM image of the isotope-labeled CVD-graphene transferred onto the epitaxial graphene/SiC substrate with a predefined mark. (b) LEED pattern obtained from the selected area indicated by the dashed circle in (a). (c) Orientation map reconstructed from the DF LEEM images of the different grains. (d) Raman 13C-2D band intensity map. The background corresponds to the SiC-related peak intensity. Reproduced from Ref. [62] with permission from the Royal Society of Chemistry.

intensity ratio between the G and the D bands can be used to estimate the distance between the defects. LEEM and LEED were used to verify the validity of the Raman-based visualization method of graphene grain boundaries, in which isotopic labeling is the key ingredient. CVD growth allows production of graphene consisting of 12C or 13C by feeding normal methane (12C, 99%) or 13C methane, respectively. 12C-graphene and 13Cgraphene are easily discriminated using Raman spectroscopy, because the Raman shift depends on the atomic mass. We first grew full coverage of graphene on a Cu foil using a mixture of normal methane and hydrogen. Then, we sequentially supplied hydrogen and a mixture of 13C methane and hydrogen to isotopically label the grain boundaries. The Raman maps of the G band intensities of 12C-graphene and 13C-graphene show that the 13Cgraphene forms a network-like structure, suggesting that the grain boundaries are selectively labeled by 13C. To check whether the network-like structure really corresponds to the grain boundaries, we measured LEEM images and Raman maps at the same areas. Fig. 2 shows the LEEM/LEED and Raman results for an isotopelabeled graphene sheet transferred onto the epitaxial graphene/SiC substrate with predefined marks. The brightfield (BF) LEEM image using the (0,0) beam [Fig. 2(a)] shows that the transferred sample is flat on a large scale, except some small wrinkles formed during the transfer process. The thick solid lines correspond to the edges of

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2.3. Microscopic thickness evaluations of graphene, hBN, and their heterostructures In this section, we demonstrate that the number of graphene layers can be digitally counted from the quantized oscillation in the energy dependence of the electron reflectivity (reflectivity spectrum) [46,47,63]. Fig. 3 shows BF LEEM images of the 4H-SiC(0001) sample graphitized by annealing at 1450 °C. The electron beam energies are (a) 2.5 and (b) 4.5 eV. These images show that the image intensities in different regions change with the energy in different manners. The energy dependence of the LEEM intensities (secular reflectivity of electrons in normal incidence) in areas A–H is shown in the energy window of (c) 0–10 eV and (d) 0–70 eV. Fig. 3(c and d) exhibits two oscillatory behaviors with long and short periods. The long-period oscillation is known to reflect the band structure of graphite from the very-lowenergy electron diffraction studies [64,65]. The inset of Fig. 3(d) shows an electronic band structure of graphite in the Γ-A direction calculated using a first-principles

(b)

(a) C A B

G D

E F H

1 m

Electron energy E-E F (eV) 0 10 20 30 40

50

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(c) H

Intensity (arb. units)

the predefined mark. Fig. 2(b) shows the LEED pattern from one of the grains, and includes two sets of the hexagonally arranged spots. The stronger spots correspond to the isotope-labeled CVD graphene on top (black arrow), while the weaker ones originate from the underlying epitaxial graphene grown on SiC (white arrow). This intensity difference is due to the electron attenuation by the top graphene. Grains with the same lattice orientation in the labeled CVD graphene are selectively visible in the DF LEEM image using the corresponding diffraction spot. The orientation map in Fig. 2(c) is reconstructed from the DF LEEM images of the different grains. Fig. 2(d) shows the Raman 13C-2D band intensity map obtained on the same sample. A Raman intensity map of the SiC-related peak intensity was used as the background, which enables us to identify the location of the mark. The linear contrast is related to the 13C-labeled region. The comparison between the DF LEEM image and the Raman map proves that all the grain boundaries are identified as the 13C labeled regions. The basic mechanism of this isotope labeling is preferential etching of 12C-graphene by hydrogen along the grain boundary and regrowth of 13C-graphene on the exposed Cu area. This visualization method of the grain boundaries is easy and fast because it is based on Raman spectroscopy measurements in air, and is compatible with other characterization methods because it is contaminationand damage-free. It could be useful for investigating the growth mechanism of graphene and effects of the grain boundaries on the graphene’s physical properties.

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10

20

30

40

50

60

70

Start voltage (V)

Fig. 3. (a)–(c) LEEM images of the 4H-SiC(0001) surface after annealing at 1450 °C. The electron beam energies were (a) 2.5 and (b) 4.5 eV. (c)–(d) Electron reflectivity spectra measured from the areas A–H in (a) and (b). The energy windows are (c) 0–10 eV and (d) 0–70 eV. Inset of (d) shows the electronic band structure of bulk graphite in the Γ-A direction calculated using the first principles method.

method, where the electron energy is measured from the Fermi level. It is clearly shown that the low-reflectivity windows correspond well to the dispersive unoccupied bands. The overall low reflectivity below 7 eV is due to the unoccupied electronic band in the Γ-A direction. Bulk graphite has a continuous electronic band in the 4–11 eV range above the Fermi level, which is well-known as the interlayer band [66–68]. This interlayer state has a charge density maximum between graphene layers and nearlyfree-electron-like dispersion parallel to them. It is also known that these states derive from the image-potential states associated with graphene [69]. The interlayer band splits into discrete energy levels in few-layer graphene. In freestanding n layer graphene, there are n-1 interlayer states. When the energy of the incident electrons coincides with one of the interlayer states, the electrons resonantly transmit through the graphene [46,47], resulting in n-1 minima in the reflectivity [70,71]. But it should be noted that an additional reflectivity minimum sometimes appears, arising from an interlayer state formed in the space between the graphene and the substrate. In

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(a) A

EF

(b) Precipitated graphene/Co

Intensity (arb. units)

any case, the number of graphene layers can be directly related to the number of minima in the reflectivity spectra. As explained later, epitaxial graphene grown on SiC(0001) has an interfacial carbon layer, normally called the buffer layer [72]. The buffer layer is a graphene sheet whose one third carbon atoms are bonded with Si atoms of the SiC surface and loses the graphene’s linear band dispersion [73,74]. Another graphene layer on top of the buffer layer has a graphene-like band structure. Two graphene sheets are needed to form epitaxial monolayer graphene on SiC(0001). There is no enough space to form an interlayer state at the buffer layer/SiC interface. Therefore, the number of epitaxial graphene layers (the number of graphene sheets minus 1) is equal to the number of minima in the reflectivity spectrum, indicating that areas A–H should be 1 to 8 layers of epitaxial graphene, respectively. This layer counting method is valid independently of the growth methods and also valid for hBN [75]. Fig. 4 compares reflectivity spectra of few-layer graphene and few-layer hBN grown on Co(0001) thin films. They were grown by segregation and by CVD, respectively. These curves clearly show quantized oscillations just above the vacuum level below which the electron beam is totally reflected. The number of graphene layers was determined by in situ observations of graphene segregation/ precipitation using LEEM, and the number of hBN layers was checked by cross-sectional TEM observations. The quantized oscillation in the reflectivity spectra of hBN suggests that hBN has an interlayer band. To confirm this, we calculated electronic structures of bulk hBN by the first principles method. The calculated electronic structures of graphite and bulk hBN [Fig. 4(a and c), respectively] confirm that graphite and bulk hBN have similar dispersive unoccupied bands in the energy windows corresponding to the reflectivity oscillations. Besides, the charge density map of hBN [inset of Fig. 4(d)], calculated at the bottom of the dispersive band, shows density maxima between the layers. Both monolayer graphene and monolayer hBN had no distinct minimum in the energy range of 0–5 eV. Similarly to epitaxial graphene grown on SiC, the number of minima in the reflectivity spectra is equal to the number of graphene or hBN sheets minus 1. No interlayer state exists at the interfaces of graphene/Co(0001) and hBN/ Co(0001), indicating strong interactions of the overlayers with the substrate. Although the reflectivity spectra of graphene and hBN show similar oscillatory behaviors at 0–6 eV, there is a clear difference at 7–13 eV. The reflectivity spectra of hBN have broad minima at this energy window, whereas those of graphene have broad maxima. The difference is easily

8 6 4 2 (c)

A

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Intensity (arb. units)

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5 4

3 2 0

5

10

15

20

Start voltage (V) Fig. 4. Calculated electronic band structures of (a) bulk graphite and (c) bulk hBN. Electron reflectivity spectra measured from (b) fewlayer graphene precipitated on a Co(0001) film and (d) few-layer hBN grown on a Co(0001) film by CVD.

understood from the electronic structures in Fig. 4(a and c). Only hBN has dispersive bands in this energy window. Graphene and hBN domains should be easily discriminated due to the difference in the reflectivity spectrum. The low-energy electron reflectivity from the heterostructures made of graphene and hBN also provides useful information about their stacking orders. We transferred CVD-grown graphene onto CVD-grown hBN to fabricate graphene/hBN heterostructures. Fig. 5(a) shows the BF LEEM image of the heterostructure. The CVD-grown graphene and hBN were mostly monolayer thick but contained local thicker regions. Therefore, the heterostructure is graphene/hBN bilayer in large areas, and is thicker than bilayer in some areas. The stacking orders in the thicker regions were determined by their shapes because the thicker regions of graphene were

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graphene hBN

(a)

(b)

A D C

1 m

structure of the graphene/hBN superlattice in Fig. 5(c) confirms the preservation of the interlayer band. On the other hand, these spectra exhibit distinct difference in the 5–15 eV range depending on the stacking orders. The dips indicate the inclusion of hBN and are deeper when the top graphene layers are thinner. Besides, the dip positions depend on the number of hBN layers. The LEEM reflectivity spectra provide useful information about how graphene and hBN are stacked in the heterostructures. 2.4. Determination of stacking structures in epitaxial bilayer graphene

EF

(c)

Intensity (arb. units)

A

B

163

Gr/Gr/BN/BN

Gr/Gr/BN

Gr/BN/BN Gr/BN 0

5

10 15 Energy (eV)

20

25

Fig. 5. (a) LEEM image of the heterostructure fabricated by transferring CVD graphene grown on a Cu foil onto CVD hBN grown on a Co(0001) thin film. (b) LEED pattern obtained from the area marked by the circle in (a). The electron beam energies were (a) 3.5 and (b) 45 eV. (c) Electron reflectivity spectra measured from the corresponding square areas in (a). Inset shows the electronic band structure of graphene/hBN superlattice.

normally polygonal and larger than irregular-shaped small bilayer hBN. In Fig. 5(a), the areas labeled by A-D are graphene/hBN, graphene/hBN/hBN, graphene/graphene/ hBN, and graphene/graphene/hBN/hBN, respectively. Fig. 5(b) shows a selected-area LEED pattern obtained from the graphene/hBN region encircled in Fig. 5(a). The hexagonally-arranged first-order diffraction spots from the top graphene and bottom hBN are clearly observed, and the graphene spots are much stronger than the hBN spots due to the electron attenuation in the graphene. The relative rotation angle between graphene and hBN can be determined as about 25°. The reflectivity spectra obtained from the heterostructures were plotted in Fig. 5(c). All the spectra have distinctive oscillations in the 0–6 eV range, and it is proven that the number of minima is equal to the total number of graphene and hBN layers minus 1. The calculated band

Graphene contains two carbon atoms at A and B sites in a unit cell. Graphite prefers Bernal stacking of the ABAB type, in which a carbon atom at a B (A) site resides just above a carbon atom at an A (B) site. However, multilayer graphene grown on the C-face SiC by thermal decomposition and grown on metals by CVD usually have out-of-plane rotational disorders. As for such turbostratic graphene, LEED can be used to determine their crystallographic orientations, and can be used to discriminate which is on top or beneath, as discussed in the previous sections. However, the surface sensitivity of slow electrons makes it difficult to determine the orientations of the deeper layers. On the other hand, graphene always grows epitaxially on the Si-face SiC. It has been shown using angle-resolved photoelectron emission spectroscopy that epitaxial bilayer graphene on the Si face is Bernal-stacked. However, Bernal-stacked bilayer graphene has three-fold symmetry unlike the six-fold symmetry of monolayer graphene. Fig. 6 illustrates A-on-B (AB) and B-on-A (BA) stacks schematically. AB and BA stacks are rotated by 180° with respect to each other. The LEED spots from the AB- and BA-stacked graphene appear at the same positions, but have different intensities in principle. Therefore, the AB and BA stacks should be identified using LEEM and LEED [48]. Fig. 7 shows BF and DF LEEM images of the SiC surface mainly covered with bilayer graphene. The DF LEEM images were obtained using the (1,0) spot. Bilayer graphene looks uniform in the BF image, but two types of domains are visible in the DF LEEM images. Among the six first-order diffraction spots at the corners of a hexagon, three of them at the corners of an equilateral triangle produce DF LEEM images with the same contrast and the other three produce images with the reversed contrast, indicating that the domains have threefold symmetry. We can therefore infer that the two types of domains correspond to AB and BA stacks. LEEM is also used to identify the stacking structure of the boundary between AB- and BA-stacked domains. We

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Fig. 6. Schematic illustrations of AB-stacked, BA-stacked, and slip-stacked bilayer graphene.

Fig. 7. (a) BF and (b)–(d) DF LEEM images of epitaxial graphene grown on 4H-SiC(0001). The DF LEEM images were obtained using the (1,0) beam. The electron beam energies were (a) 5.5, (b) 44.5, (c) 51.1, and (d) 58.1 eV, respectively.

found that the domain boundaries are visible at certain energies in both BF and DF LEEM images. Fig. 7(c) is the (1,0) DF LEEM image at 51.1 eV, showing that the domain boundaries are seen bright but are only partially visible. Among the six first-order diffraction beams, the opposite two beams produce the same DF LEEM images, which are different from those obtained by the other pairs. When these three different DF LEEM images are overlaid, most of the boundaries are visible. There are three types of domain boundaries, and each has two-fold symmetry. To determine the atomic structures inside the domains and at the boundaries, we calculated the LEED I–V curves

for graphene films with infinite thickness for various stacking structures and compared them with the experimental data. The calculations were done for AB, AA, and slip stacks. The slip stack is indicated between AB and BA stacks in Fig. 6 and is known to exist in highly oriented pyrolytic graphite [76]. AA and slip stacks were chosen to examine the boundary structure. We found that the (0,0) I–V curve calculated for AB-stacked graphite reproduces the experimental result fairly well even without any structural optimizations. In reality, our sample has only three graphene layers (including the buffer layer) on SiC. Therefore, the agreement between the experimental and the calculated LEED I–V curves indicates that the topmost couple of graphene layers determine the main features in the LEED I–V curve. This is consistent with the very short inelastic mean free paths of electrons in a solid (less than 1 nm) at the energies we investigated (40–150 eV) [77]. Next, we analyze the I–V curves of the (1,0) and (0,1) beams to verify the coexistence of AB- and BA-stacked domains. Between AB and BA stacks, the (1,0) and (0,1) intensities are reversed. As the AA stack has six-fold symmetry, its (1,0) and (0,1) intensities are the same. The slip stack has two orthogonal mirror symmetry planes, and therefore, among the six first-order diffraction beams, two beams are independent. Fig. 8 shows that the calculated I–V curves for AB stack reproduce the main features seen in the experimental curves. The (1,0) and (0,1) beams have largely different intensities at ~45 and ~60 eV, which agrees with the clear domain contrasts in Fig. 7(b and d) respectively obtained at 44.5 and 58.1 eV. Thus, we confirm that the domains observed in the DF LEEM images of bilayer graphene have different stacking orders, i.e., AB and BA stacks. AB and BA stacks can be converted by sliding one of the two layers relative to the other. The boundaries should

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Intensity (arb. units)

Slip stack

AA AB/BA

Exp.

30

40

50

60

70

80

Energy (eV) Fig. 8. Calculated and measured LEED I–V curves for the (1,0) and (0,1) beams.

have intermediate structures that appear during this sliding. AA and slip stacks are plausible highly symmetric boundary structures, but the two-fold symmetry determined experimentally is consistent with the slip stack but not with AA stack. The boundary is too narrow to measure the LEED I–V curves, but the boundaries are seen at high contrast at the peak energies of the calculated (1,0) and (0,1) I–V curves for the slip stack. The DF LEEM image in Fig. 7(c) was taken at the energy indicated by the arrow in Fig. 8. We also confirmed that, at the energies at which the boundaries make contrast in the BF LEEM images, the calculated LEED intensities for the slip stack and AB stack are largely different. Thus, we conclude that the slip stack appears at the boundaries of the stacking domains. This conclusion is supported by the first-principles calculation results showing that the slip stack is energetically more stable than AA stack [78]. 3. Epitaxial growth of graphene on SiC 3.1. Growth processes of epitaxial graphene on SiC According to the cross-sectional TEM investigations of nucleation and growth of graphene on SiC, graphene

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nucleates at the step region and grows in a layer manner on SiC(0001) [79], whereas pits are formed on the terrace of SiC(000 1 ) by the local sublimation of Si, and multilayer graphene nucleates therein [80]. Besides, while the buffer layer serving as a template for epitaxial graphene growth exists on SiC(0001), there is no buffer layer on SiC(000 1 ) [72]. This presence or absence of the buffer layer makes a big difference in the structure of the grown graphene. We first summarize structural changes on SiC(0001) during annealing in UHV. Starting from the √3 × √3 structure, in which 1/3 monolayer of Si atoms are arranged at the T4 sites on the bulk terminated SiC surface, annealing at about 1080 °C causes the appearance of the 6√3 × 6√3 structure [81]. When a graphene sheet is placed on the SiC(0001) surface with a relative rotation angle of 30°, the 6√3 × 6√3 SiC(0001) unit cell nearly matches with the 13 × 13 graphene unit cell. About 1/3 of C atoms in the graphene sheet make chemical bonds with the Si atoms on the substrate, resulting in a stable structure with the 6√3 × 6√3 periodicity. Due to the chemical bonding, the graphene sheet loses its intrinsic linear electronic band structure [82,83], and is usually called the buffer layer. Successive heat treatments above 1080 °C induce the gradual development of the 6√3 × 6√3 LEED pattern. The 6√3 × 6√3 LEED spots are visible even after the growth of epitaxial graphene, indicating that the buffer layer is preserved at the interface between the epitaxial graphene and the substrate. Multilayer graphene is formed after annealing at higher temperatures like 1400 °C [81]. The graphene growth by thermal decomposition of SiC is quite different from ordinary crystal growth in which materials are supplied from outside: (1) Carbon atoms for graphene growth are supplied by the thermal decomposition of the substrate, and therefore, the substrate is etched during the growth. The amount etched during monolayer graphene growth is nearly three SiC bilayers. For this reason, the SiC surface morphology greatly changes during the graphene growth. (2) New graphene sheets always grow at the interface [84–86]. As for SiC(0001), after the surface is covered with the buffer layer, a new graphene sheet (new buffer layer) is formed at the interface, whereby the old buffer layer loses the chemical bonds with the substrate and obtains the graphene’s physical properties. Graphene is thickened through repeated insertion of new buffer layers into the interface. Since the buffer layer plays as a template of the graphene growth, graphene on SiC(0001) always has epitaxial relationship with

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(a) ~SiC 0 0

500 nm

1

(b) ~buffer layer 0 1

0

2

2 500 nm

0

1

2

2 1

0

500 nm

(c) ~monolayer 2

3 2

3 500 nm

1

1 (d) ~bilayer

Fig. 9. Schematic illustrations of surface structures during graphene growth on the Si-face SiC in UHV [45]. Corresponding LEEM images and atomic force microscopy image [inset of (b)] are also shown. The numbers in the figures correspond to the number of epitaxial graphene layers. The preference for bilayer or trilayer steps on 4H- or 6HSiC(0001) surfaces is neglected for simplicity. © IOP Publishing. Reproduced with permission. All rights reserved.

the substrate. In addition, the top layer of graphene covers the entire surface of the substrate continuously, regardless of the presence of the substrate steps and the boundaries between the areas with different numbers of layers. (3) Graphene growth requires Si atoms to diffuse from the substrate to the surface through the graphene layers. However, graphene is basically impermeable to all atoms. Therefore, the growth rate decreases rapidly with the number of layers. Thereby, Si and C atoms are trapped at the interface between the buffer layer and the SiC substrate, which activates the planarization of the substrate. Based on the above peculiarities, we describe the growth processes of graphene on the Si-face SiC(0001) in more detail. Fig. 9 shows schematic illustrations of the surface morphology changes during graphene growth in

UHV along with the corresponding LEEM images. Starting from the √3 × √3 surface, the buffer layer first nucleates at the steps. Fig. 9(b) indicates that the surface considerably roughens during the buffer layer growth. The carbon atom density in graphene is three times as high as that in a SiC bilayer. Many Si and C atoms are released from the steps for graphene growth, resulting in active retreat of the steps. As shown by Hannon and Tromp, the buffer layer pins the steps [87]. Therefore, it normally happens that all the steps come to stop at the edges of the buffer layer before the buffer layer completely covers the surface. To convert the √3 × √3 areas surrounded by the buffer layer, craters are formed, and the surface severely roughens during the buffer layer growth. This high step density after the buffer layer growth leads to dense nucleation of monolayer graphene islands. High-density monolayer islands also facilitate nucleation of bilayer graphene. Thus, many bilayer graphene islands appear before a monolayer is completed [Fig. 9(c)]. Bilayer has an equal chance of taking AB or BA stack. Therefore, bilayer graphene grown in UHV contains highdensity stacking domains. Nevertheless, the regularity of the step/terrace structure is largely recovered after bilayer graphene growth [Fig. 9(d)]. The impermeability of graphene to Si atoms is an important factor in this morphology transition. Graphene growth requires the desorption of Si atoms from the surface. We measured the time evolution of the graphene thickness on vicinal SiC surfaces, and clarified that the growth rate quickly saturates with the thickness [88]. Thicker graphene is more resistive against the Si out-diffusion, which increases probabilities that the Si and C atoms detached from the substrate steps reattach back to the steps. Then, the rough SiC substrate becomes smoother to lower the interface energy. The annealing time dependence of the graphene thickness indicated that the average growth rate from 1.5 to 2 layers is roughly one order of magnitude smaller than that from 0 to 1.5 layers [88]. Thus, the step/terrace regularity is largely recovered during bilayer growth. Fig. 10 shows LEEM images of highly uniform epitaxial graphene grown on the Si-face SiC. By controlling the roughening and smoothening of the substrate during graphene growth in UHV with the annealing temperature and time, bilayer graphene can be obtained relatively uniformly [Fig. 10(b)]. If the SiC substrate is sufficiently flat after bilayer growth, trilayer graphene grows in a layer manner. Trilayer graphene in Fig. 10(c) was grown by annealing at relatively low temperatures to avoid pit nucleation. Inversely, to obtain a uniform monolayer, the surface roughness of the buffer layer needs to be reduced. This

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Fig. 10. LEEM images of highly uniform (a) monolayer, (b) bilayer, and (c) trilayer graphene grown on the Si-face SiC. (a) was grown in an Ar environment, and (b) and (b) were grown in UHV. The electron beam energies were (a) 4.0 (b) 5.0, and (c) 5.5 eV.

has been practically achieved in annealing the SiC substrates in an Ar atmosphere [89,90]. Ar atoms block the desorption of Si atoms from the surface, leading to an increase in the graphene growth temperature. Thus, the buffer layer grows while maintaining the surface flatness. Monolayer graphene in Fig. 10(a) was grown in an Ar environment, and is much smoother that monolayer graphene grown in UHV [Fig. 9(c)]. Annealing in Ar dramatically improves the uniformity of monolayer graphene. After the surface is uniformly covered by monolayer graphene, the graphene growth rate decreases substantially. The growth is effectively self-limiting at monolayer. Toward deeper understandings of the graphene growth mechanism on the Si-face SiC, we also theoretically investigated the stability and reactivity of the SiC steps during graphene growth by the first-principles calculation [91]. We showed that the stability and reactivity during the buffer layer growth depend on the temperature, Si pressure, and C coverage. At relatively high temperatures and low Si pressures, the Si desorption preferentially occurs on the terrace, resulting in the surface roughness evolution during the buffer layer growth. However, as decreasing the temperature or increasing the Si pressure, the Si desorption from the steps becomes more active. In the graphene growth by annealing SiC in Ar, the Ar atmosphere effectively enhances the Si pressure, which allows the buffer layer to grow while keeping the surface smooth. We also found that when the second

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graphene sheet grows at the buffer layer/SiC interface, the Si desorption and the C adsorption/aggregation preferentially occurs at the step edges. This promotes the recovery of the surface flatness during graphene growth. All these theoretical results are fully consistent with the above experimental results. Annealing in an Ar atmosphere has also provided important clues to understand the graphene growth mechanism. This is because the structural changes occur on a larger scale and become more visible. Finger and arrow features emerge along retreating single or double SiC bilayer steps during the growth of epitaxial monolayer graphene [92–94]. When single or double SiC bilayer steps retreat, the whole retreated area cannot be converted into graphene because the C atom density in graphene is nearly three times larger than that in SiC. This insufficient C atom supply can make straight steps unstable. In contrast, triple-layer steps move while keeping their shapes and form a uniform band of monolayer graphene on the swept area [93]. The arrow and finger features remind us of the importance of C surface diffusion in graphene growth. The arrow feature grows in a steady state because Si sublimation and graphene growth are spatially connected through the C diffusion [93]. Under the insufficient C supply, when the C incorporation at the graphene edge is asymmetric between the forward and the backward directions, finger features could be formed [94]. Finger features resemble the comb-shaped steps formed during Ga deposition on Si(111) [95] and could be a universal phenomenon. Meanwhile, SiC(000 1 ) does not have a buffer layer, which plays a role in determining the graphene crystal orientation on SiC(0001). The first graphene layer shows the intrinsic physical properties of graphene [82,83]. Because the first layer graphene on SiC(000 1 ) weakly interacts with the surface, compared to the buffer layer on SiC(0001), the crystal orientation of graphene nuclei is relatively random. Therefore, a graphene sheet after the coalescence of the islands becomes polycrystalline. Moreover, because new graphene formed at the interface does not always align to the existing graphene, the rotational disorder is likely to occur between layers. In addition to the nucleation of multilayer graphene in the etched pits, the grain boundaries of the polycrystalline graphene serve as efficient diffusion paths of Si atoms from the substrate to the surface [96], and therefore, the graphene layers are easy to thicken. 3.2. Atomic intercalation of epitaxial graphene on SiC Thermal decomposition of the Si-face SiC yields uniform monolayer and bilayer graphene. The carrier

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mobility of epitaxial monolayer graphene is comparable to that of exfoliated graphene on SiO2 at low temperatures [97]. At room temperature, however, epitaxial graphene has a lower mobility than exfoliated graphene. It has been reported that the stronger temperature dependence is due to carrier scatterings by a low-energy phonon mode, and that the interfacial layers including the buffer layer are responsible for this phonon mode [97,98]. Therefore, as a method of eliminating the buffer layer and decoupling graphene from the SiC substrate, atomic intercalation gathers intense attention. It has already been shown that various elements such as H [99], Li [100], O [101], F [102], Si [103], Ge [104], and Au [105] were intercalated into the interface between the graphene and the SiC substrate. The atomic intercalation greatly changes the electronic structures of the graphene sheets grown on SiC. Here, we consider hydrogen as a representative intercalant. When the buffer layer sample is heated in a hydrogen gas atmosphere typically at 700 °C [106], hydrogen atoms break the chemical bonds between the buffer layer and the SiC substrate, and terminate the dangling bonds of the SiC substrate [99]. Thus, the buffer layer is separated from the substrate, and acquires the electronic structure of graphene. Hydrogen intercalation creates so called quasi-freestanding monolayer graphene. Similarly, when epitaxial monolayer graphene is annealed in H2 at around 1050 °C [107], quasi-freestanding bilayer graphene is obtained [99]. Quasi-freestanding graphene also produces the oscillations in the low-electron reflectivity spectra, where the number of minima is equal to the number of graphene sheets [99]. This is because there is enough space to form an interlayer state at the interface between the bottom graphene and the SiC substrate, which suggests that graphene is effectively decoupled from the SiC substrate. The measured carrier mobility of quasi-freestanding monolayer graphene is less temperature dependent than that of epitaxial graphene, but their reported values are almost the same at room temperature [98]. It is found that the mobility of quasi-freestanding monolayer graphene is lowered by carrier scatterings by charged impurities. There is a possibility that all of the dangling bonds on the substrate are not terminated with hydrogen atoms, and the remaining dangling bonds act as the scattering centers [106,108,109]. On the other hand, quasi-freestanding bilayer graphene has higher mobility than epitaxial bilayer graphene [107,110,111]. The LEEM results show that stacking domains in quasi-freestanding bilayer graphene are larger than those in epitaxial bilayer graphene [110], indicating that the boundaries of the stacking domains could degrade the electronic transport properties. Whereas

the quasi-freestanding bilayer graphene was converted from epitaxial monolayer graphene grown in the Ar environment, epitaxial bilayer graphene was grown at lower temperatures in UHV. This difference in the growth temperature led to the different domain sizes. Atomic intercalation has great potential to control and enhance the electronic properties of graphene on SiC, but there is still a lot of room for improvement. Further optimization of intercalation conditions is desired. 4. CVD growth of graphene CVD typically yields polycrystalline graphene, and the higher density of the grain boundaries degrades the physical properties. Therefore, one of the challenges in graphene CVD is single-crystal growth. Two different approaches are mainly pursued toward this direction. The first one is to align all the graphene islands with a single crystal orientation. Rotational disorder does not appear even at a high density of the islands. However, there remains a possibility that, even though the islands have the same orientation, atomic level defects could be formed when they coalesce [112]. The second approach is to grow a single graphene island as large as possible. Because the distance between the nuclei limits the single-crystal grain size, it is critically important to reduce the nucleation density. In the first approach, heteroepitaxial thin films instead of bulk single crystals are usually used as substrates for CVD growth of graphene due to cost issues. Choice of the substrate orientation is crucial for obtaining a single orientation. Figs. 11 and 12 show LEEM images of graphene grown on heteroepitaxial substrates of Cu(111)/ MgO(111) and Cu(100)/MgO(100) at 1000 °C by CVD. While single-oriented graphene grows on Cu(111), graphene grown on Cu(100) is polycrystalline, mainly

Fig. 11. (a) BF and (b) DF LEEM images of graphene grown on a heteroepitaxial Cu(111) film by CVD [113]. The electron beam energy was 44.5 eV. Inset shows the LEED pattern obtained at 45 eV from the selected area indicated by the dashed circle in (a).

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Therefore, island rotation takes place at enough high temperatures, while the islands keep aligned at lower temperatures. Further increase in the growth temperature leads to polycrystalline graphene with wide orientation distributions, suggesting the enhanced thermal fluctuation of the Cu lattice close to the melting temperature of 1083 °C. The substrate orientation and growth conditions are important for controlling the crystallinity of graphene. In the second approach of single-crystal growth, mainly three strategies have been adopted for the reduction of nucleation density:

Fig. 12. (a) BF LEEM image and (b)–(d) DF LEEM images of graphene grown on a heteroepitaxial Cu(100) film by CVD [113]. The electron beam energy was 44.5 eV. The LEED patterns in (b) and (c) were obtained from the corresponding encircled areas. (b) and (c) were obtained using the first-order diffraction beams, and were overlapped in (d).

consisting of two types of grains rotated by 90° with each other [113]. Graphene has six-fold symmetry, and the topmost Cu atoms on Cu(111) and Cu(100) respectively have six-fold and four-fold symmetries. The LEED analyses of the epitaxial relationship between the Cu substrates and graphene indicated that the unit cells of Cu(111) and graphene are aligned to the same orientation. On Cu(100), on the other hand, the unit vectors of the graphene unit cell is parallel to either of the two orthogonal Cu(100) unit vectors. These two equivalent alignments are a natural consequence of the difference in the space symmetry between graphene and Cu(100). However, even on Cu(111), orientational disorders occur at high growth temperatures [114]. At relatively low temperatures of 930–1030 °C, graphene grows in the same orientation with Cu(111), but graphene grown above 1040 °C is rotated predominantly by ±3.4° from the aligned direction. There is a lattice mismatch of about 4% between Cu(111) and graphene, which would stabilize the long-range ordered structure. This orientation change might be explained by the size dependence of the stable structure. When the island size is small, aligned islands have smaller formation energy than the 3.4°rotated islands. But the formation energies are reversed when the size becomes larger than the critical size. The island rotation by 3.4° is an energetically activated process.

(1) Placing Cu substrates in confined geometries, such as Cu enclosure and Cu tube, during low-pressure CVD [115,116]. Cu atoms evaporated from the inner surface are redeposited on the inner surface, leading to a much smoother inner surface than the outer surface. Therefore, the nucleation density is greatly reduced on the inner surface. (2) Pre-treatments of Cu substrates before graphene CVD, such as the high-pressure annealing in H2 [117], and melting and subsequent resolidification of Cu films [118]. The pre-treatments are intended to reduce the roughness and remove the structural irregularities like impurities that act as heterogeneous nucleation sites. (3) Oxygen adsorption on Cu substrates, resulting in the passivation of the Cu surface active sites for graphene nucleation [119]. Under the control of the CVD growth parameters, the size of the single-crystal graphene has increased from ~10 μm in 2010, 0.5 mm in 2011 [115], 5 mm in 2013 [120], and to 1 cm in 2013 [119]. 1-cm-sized graphene grains were obtained on the oxygen-adsorbed Cu foils. It has also been shown that oxygen atoms adsorbed on Cu not only reduce the nucleation density, but also enhance the graphene growth. However, despite the growth enhancement effect, it took 12 hours to grow graphene islands 1 cm in size. Therefore, further enhancement of the growth rate is a critical issue. To solve this issue, we have also tried to establish a simple procedure of growing large-scale single-crystal graphene at a faster rate, and succeeded in growing 3-mmsized single crystals for 2 h [121]. This is one of the fastest growth rates ever reported. Commercial Cu foils were used as catalytic substrates without any chemical treatments for atmospheric-pressure CVD at 1075 °C. Preannealing in an Ar-only environment at the growth temperature and placing the foils on inert supports are the keys of our growth procedure. As the pre-annealing time increases, the nucleation density of graphene first

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decreases and then increases. During the pre-annealing in the non-reductive environment, oxygen atoms gradually evaporate from the surface, and the surface roughness is reduced simultaneously. Decreases in the oxygen concentration and surface roughness have two competing effects on the graphene nucleation, and therefore the nucleation density is minimized at a certain pre-annealing time. We also found that graphene grows faster on the inner surface facing the inert support than on the outer surface. This is probably due to a higher density of the active carbon species on the inner surface, which results from the reflection of the evaporated active carbon species by the support. Lastly in this section, we briefly mention CVD growth of bilayer graphene, which gathers attention because a tunable band gap can be induced in AB-stacked bilayer graphene by an external perpendicular electric field. Bilayer graphene growth requires not only the number of layers but also the stacking structure to be controlled, and is therefore much more difficult than monolayer growth. There are some examples of growing bilayer graphene on Cu foils using complex CVD procedures [122,123], but the most promising approach to date is CVD growth on Cu-Ni alloy substrates [124–126]. The catalytic activity and carbon solid solubility are controlled by the alloying so as to promote the bilayer growth. AB-stacked bilayer graphene with coverage of 98% on a 3-inch substrate has been obtained [126]. 5. CVD growth of hBN on heteroepitaxial Co films [75] Approaches toward single-crystal growth are common between graphene and hBN. Single alignment and reduction of the nucleation density are their essences. Here, we show our results on CVD growth of hBN on heteroepitaxial Co(0001) films sputtered on c-plane sapphire. hBN forms a commensurate 1 × 1 structure on Ni(111) due to the small lattice mismatch. Because the lattice mismatch between hBN and Co(0001) is similar to that between hBN and Ni(111), the hBN/Co(0001) system is also expected to form a commensurate 1 × 1 structure. Co(0001) is therefore suited for single alignment of hBN islands. Additionally, the heteroepitaxial Co(0001) films are much smoother than the commercial Co foils, which would lead to a fewer nucleation density. However, in contrast to graphene with six-fold symmetry, hBN has three-fold symmetry. When monolayer hBN is rotated by 180°, the B-N bond direction is reversed. Atomic defects would be formed at the boundary of the grains with the original and rotated orientations. So the coexistence of these grains should be avoided.

Fig. 13. (a)–(b) BF LEEM images of triangular hBN islands grown on a heteroepitaxial Co(0001) film by CVD. The electron beam energies were (a) 12.5 and (b) 70 eV. (c)–(e) LEED patterns from the three different areas indicated in (b). The electron beam energies were (c) 32 eV and (d)–(e) 45 eV. Schematic illustrations of (f) top(N)hcp(B) and (g) top(N)-fcc(B) structures of hBN on Co(0001).

Single-crystal growth of hBN is more difficult than that of graphene. We grew hBN on heteroepitaxial Co films at 1000 °C using ammonia borane as the precursor. Fig. 13(a and b) shows the LEEM images of the Co surface partially covered with hBN islands. Individual triangular hBN islands with sides longer than 10 μm are seen. Additionally, two types of triangles with opposite orientations are visible. Fig. 13(c–e) shows the LEED patterns of the three different positions in Fig. 13(b), as marked accordingly. The Co surface shows six first-order diffraction spots, out of which three are dominant [Fig. 13(c)]. The relatively weak intensity could be due to the oxidation of Co upon exposure to air. Similarly, Fig. 13(d) represents the LEED pattern recorded from one of the triangular hBN islands. Six first-order diffraction spots are identified. The oppositely directed triangular hBN island also showed a similar hexagonal LEED pattern [Fig. 13(e)]. The direction of the diffraction spots matches between the hBN islands and the Co substrate. The matched direction and absence of moiré diffraction spots in the hBN LEED patterns suggest a commensurate 1 × 1 structure. However, among the six first-order LEED spots from the hBN islands, three at the corner of the equilateral triangles are

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stronger, and the equilateral triangles constituted from the stronger spots are opposite between the two antiparallel islands. In the energy range we investigated, 40– 190 eV, LEED is extremely sensitive to the topmost surface structure. The emergence of the three bright diffraction spots could be ascribed to atomic structures of hBN. Thus, we naturally think that the B-N bond direction is reversed between the two triangular islands. The positions of the B and N atoms with respect to the underlying Co atoms can be inferred from the previous experiments and theoretical calculations. Similar studies of monolayer hBN on Ni(111) demonstrated that two hBN domains of different orientations coexist [127,128]. N atoms always locate on top of the topmost Ni atoms, but B atoms locate either on top of the second layer Ni atoms, top(N)-hcp(B), or above the hollow sites, top(N)-fcc(B) [129,130]. According to the theoretical calculations, two stable configurations of monolayer hBN, top(N)-hcp(B) and top(N)-fcc(B), have quite close formation energies on Ni(111) and Co(0001) surfaces. Besides, most reports have assigned the termination of the triangular-shaped hBN islands grown by CVD to nitrogen atoms [128,131,132]. Previous theoretical studies also pointed out that freestanding hBN favors the N-terminated zigzag edge when the edge is terminated by H atoms [133]. Fig. 13(f and g) represents models for the two triangles formed on the Co substrate. Fig. 14 shows BF and DF LEEM images of hBN with almost full coverage. The surface is covered with monolayer hBN uniformly except small patches of bilayer hBN. The two oppositely directed triangles grew in size and merged to cover the whole surface. Nitrogen has negligible solid solubility in Co [134], and the Co surface loses the catalytic activity when it is covered with hBN. Therefore, hBN growth is self-terminated at monolayer. The boundaries between the hBN grains can be clearly visualized in the DF LEEM image, Fig. 14(b and c). Comparison between the BF and the DF images indicates that bilayer hBN preferentially nucleates at the grain boundaries. The LEED patterns obtained from the bilayer hBN also indicated that the second layer grows on top of the first layer and is often rotated from the first layer. While we stated that hBN growth is self-limiting at monolayer, we showed the low-energy electron reflectivity of few-layer hBN in Fig. 4. The difference between the monolayer and the multilayer samples was the substrate preparation. The monolayer and multilayer hBN were grown on heteroepitaxial Co films sputtered on c-plane sapphire at 380 °C and 280 °C, respectively. The Co films prepared at lower temperatures exhibited rougher morphology. The larger roughness could lead to dense

(a)

171

4.5 eV

2

1 m (b)

58 eV

DF(1,0) (c)

DF(0,1)

58 eV

Fig. 14. (a) BF and (b)–(c) DF LEEM images of a hBN sheet grown on a heteroepitaxial Co(0001) film by CVD. The bright patches in (a) correspond to bilayer hBN.

heterogeneous nucleation, resulting in multilayer hBN films with a wide thickness distribution. The two inequivalent triangular islands never merge at atomic levels, consequently giving rise to line defects at the boundary [128]. Additional control is needed to achieve single-orientation monolayer hBN. Because the formation energies of the two stable configurations of hBN on Co(0001), top(N)-hcp(B) and top(N)-fcc(B), are almost the same [130], it seems quite difficult to grow one of the two orientations selectively. We also confirmed the nucleation of two antiparallel triangular hBN islands on the heteroepitaxial Ni(111) films. Recently, Orlando and co-workers succeeded in epitaxial growth of singleorientation monolayer hBN on Ir(111) by cyclic depositions of borazine at room temperatures followed by annealing at 1270 K [135]. They also showed that the borazine deposition at 1270 K resulted in monolayer hBN with grains of two orientations. This behavior can be explained by a small energy difference between the fcc(B) and the hcp(B) configurations. The energy difference allows the nuclei of the fcc(B) configuration to selectively emerge during the initial stages of the CVD process at lower temperatures. For the CVD growth at higher temperatures, however, the thermal energy exceeds the energy difference, resulting in the nucleation of both the fcc(B) and the hcp(B) configurations. The reduction of the hBN nucleation density is also in progress, and the reported maximum size of single hBN

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grains is rapidly increasing. The size of single-crystal grains grown inside the Cu enclosure and on Cu-Ni alloy reached up to ~70 μm [136] and ~130 μm [137] in edge length, respectively. However, the applications of hBN as dielectric layers highly demand multilayer hBN with desired number of layers. Therefore, in addition to improving the crystal quality of monolayer hBN, controlling the thickness of hBN is important.

6. Conclusions LEEM is a powerful tool for investigating the growth mechanism of 2D materials and their structural properties, such as thickness, orientation and stacking order. We have taken advantage of LEEM in deepening the understanding of the growth of 2D materials. Our research toward this direction is summarized as follows. LEEM can be used to microscopically determine the number of graphene layers and greatly contribute to clarifying the growth processes of epitaxial graphene on the Si face of SiC. During annealing in UHV, the surface morphology changes from an initial smooth SiC surface to a rough buffer-layer surface, and further to a smooth bilayer-graphene surface. The SiC substrates are etched by three bilayers during the buffer layer growth. The buffer layer restricts SiC to be etched underneath, leading to the etching of the SiC substrate from the area uncovered with the buffer layer. Thus, the substrate roughens substantially during the buffer layer growth. After the buffer layer growth, the impermeability of graphene to Si atoms plays an essential role in the recovery of the smooth surface. Owing to the understanding of the graphene growth processes, we succeeded in growing uniform bilayer graphene on a micrometer scale in UHV. On the other hand, uniform monolayer graphene can be grown in an Ar environment. Because Ar atoms block the evaporation of Si atoms, the graphene growth temperature is effectively raised. Therefore, the substrate remains flat during the buffer layer growth, and graphene grows nearly in the layer-by-layer mode. We also observed the graphene segregation/ precipitation on polycrystalline Ni foils in situ using LEEM. On the ultra-flat Ni foils, single-crystal graphene emerged from a small number of nucleation sites spread over the grain boundaries of the Ni foil, continuously in a carpet-like manner, and grew to a macroscopic size at least larger than 100 μm. However, the graphene segregation/precipitation on polycrystalline Ni foils always provided films with uncontrolled thickness. Controllability of the thickness should be improved by adjusting the segregation/precipitation parameters.

We have also conducted research toward singlecrystal growth of graphene and hBN by CVD. Millimetersized single-crystal graphene islands were grown on polycrystalline Cu foils by relatively simple CVD procedures. Pre-annealing of Cu foils in the Ar-only environment before atmospheric pressure CVD reduced the surface roughness while keeping active oxygen species, which greatly decreased the graphene nucleation density. In addition, the growth rate was enhanced by placing the Cu foils in the partially confined geometry. After the optimization of the growth conditions, isolated singlecrystal graphene islands as large as 3 mm were obtained with only 2-h growth time. Still, the growth rate is relatively slow, and there is a great need for breakthroughs in speeding up the growth rate. We used the heteroepitaxial Co films as substrates for CVD growth of hBN, aiming at the nucleation of epitaxial hBN islands at a lower density. hBN growth was self-terminated at monolayer due to low solid solubility of nitrogen in Co, but the grown hBN was polycrystalline, with grains of two antiparallel orientations. Control of the hBN orientation as well as the number of hBN layers are important for future applications of 2D heterostructure devices. Acknowledgements The authors thank Prof.Y. Homma, Dr. G. Odahara, Prof. S. Mizuno, Prof. H. Ago, Dr. Y. Ogawa, Prof. S. Tanaka, and Dr. S. Suzuki for their collaborations. References [1] P.R. Wallace, The band theory of graphite, Phys. Rev. 71 (1947) 622–634. [2] K.S. Novoselov, A.K. Geim, S.V. Morozov, D. Jiang, Y. Zhang, S.V. Dubonos, et al., Electric field effect in atomically thin carbon films, Science 306 (2004) 666–669. [3] A.C. Ferrari, F. Bonaccorso, V. Falko, K.S. Novoselov, S. Roche, P. Bøggild, et al., Science and technology roadmap for graphene, related two-dimensional crystals, and hybrid systems, Nanoscale 7 (2015) 4598–4810. [4] C. Lee, X.D. Wei, J.W. Kysar, J. Hone, Measurement of the elastic properties and intrinsic strength of monolayer graphene, Science 321 (2008) 385–388. [5] Y. Wang, R. Yang, Z. Shi, L. Zhang, D. Shi, E. Wang, et al., Super-elastic graphene ripples for flexible strain sensors, ACS Nano 5 (2011) 3645–3650. [6] J.H. Chen, C. Jang, S. Xiao, M. Ishigami, M.S. Fuhrer, Intrinsic and extrinsic performance limits of graphene devices on SiO2, Nat. Nanotechnol. 3 (2008) 206–209. [7] R.R. Nair, P. Blake, A.N. Grigorenko, K.S. Novoselov, T.J. Booth, T. Stauber, et al., Fine structure constant defines visual transparency of graphene, Science 320 (2008) 1308.

H. Hibino et al. / Progress in Crystal Growth and Characterization of Materials 62 (2016) 155–176 [8] F. Wang, Y. Zhang, C. Tian, C. Girit, A. Zettl, M. Crommie, et al., Gate-variable optical transitions in graphene, Science 320 (2008) 206–209. [9] A.A. Balandin, Thermal properties of graphene and nanostructured carbon materials, Nat. Mater. 10 (2011) 569– 581. [10] K. Watanabe, T. Taniguchi, H. Kanda, Direct-bandgap properties and evidence for ultraviolet lasing of hexagonal boron nitride single crystal, Nat. Mater. 3 (2004) 404–409. [11] C.R. Dean, A.F. Young, I. Meric, C. Lee, L. Wang, S. Sorgenfrei, et al., Boron nitride substrates for high-quality graphene electronics, Nat. Nanotechnol. 5 (2010) 722–726. [12] A.K. Geim, I.V. Grigorieva, Van der Waals heterostructures, Nature 499 (2013) 419–425. [13] L. Britnell, R.V. Gorbachev, R. Jalil, B.D. Belle, F. Schedin, A. Mishchenko, et al., Field-effect tunneling transistor based on vertical graphene heterostructures, Science 335 (2012) 947– 950. [14] S. Bae, H. Kim, Y. Lee, X. Xu, J.-S. Park, Y. Zheng, et al., Rollto-roll production of 30-inch graphene films for transparent electrodes, Nat. Nanotechnol. 5 (2010) 574–578. [15] Y.-M. Lin, C. Dimitrakopoulos, K.A. Jenkins, D.B. Farmer, H.Y. Chiu, A. Grill, et al., 100-GHz transistors from wafer-scale epitaxial graphene, Science 327 (2010) 662. [16] Y.-M. Lin, A. Valdes-Garcia, S.-J. Han, D.B. Farmer, I. Meric, Y. Sun, et al., Wafer-scale graphene integrated circuit, Science 332 (2010) 1294–1297. [17] Z. Guo, R. Dong, P.S. Chakraborty, N. Lourenco, J. Palmer, Y. Hu, et al., Record maximum oscillation frequency in C-face epitaxial graphene transistors, Nano Lett. 13 (2013) 942– 947. [18] K. Nakada, M. Fujita, G. Dresselhaus, M.S. Dresselhaus, Edge state in graphene ribbons: nanometer size effect and edge shape dependence, Phys. Rev. B 54 (1996) 17954–17961. [19] L. Yang, C.-H. Park, Y.-W. Son, M.L. Cohen, S.G. Louie, Quasiparticle energies and band gaps in graphene nanoribbons, Phys. Rev. Lett. 99 (2007) 186801. [20] F. Schwierz, Graphene transistors, Nat. Nanotechnol. 5 (2010) 487–496. [21] M.Y. Han, B. Özyilmaz, Y. Zhang, P. Kim, Energy band-gap engineering of graphene nanoribbons, Phys. Rev. Lett. 98 (2007) 206805. [22] X. Li, X. Wang, L. Zhang, S. Lee, H. Dai, Chemically derived, ultrasmooth graphene nanoribbon semiconductors, Science 319 (2008) 1229–1232. [23] J. Cai, P. Ruffieux, R. Jaafar, M. Bieri, T. Braun, S. Blankenburg, et al., Atomically precise bottom-up fabrication of graphene nanoribbons, Nature 466 (2010) 470–473. [24] L. Jiao, L. Zhang, X. Wang, G. Diankov, H. Dai, Narrow graphene nanoribbons from carbon nanotubes, Nature 458 (7240) (2009) 877–880. [25] D.V. Kosynkin, A.L. Higginbotham, A. Sinitskii, J.R. Lomeda, A. Dimiev, B.K. Price, et al., Longitudinal unzipping of carbon nanotubes to form graphene nanoribbons, Nature 458 (2009) 872– 876. [26] Y. Zhang, T.T. Tang, C. Girit, Z. Hao, M.C. Martin, A. Zettl, et al., Direct observation of a widely tunable bandgap in bilayer graphene, Nature 459 (2009) 820–823. [27] D.C. Elias, R.R. Nair, T.M.G. Mohiuddin, S.V. Mozorov, P. Blake, M.P. Halsall, et al., Control of graphene’s properties by reversible hydrogenation: evidence for graphane, Science 323 (2009) 610–613.

173

[28] F.H.L. Koppens, T. Mueller, P. Avouris, A.C. Ferrari, M.S. Vitiello, M. Polini, Photodetectors based on graphene, other twodimensional materials and hybrid systems, Nat. Nanotechnol. 9 (2014) 780–793. [29] F. Schedin, A.K. Geim, S.V. Morozov, E.W. Hill, P. Blake, M.I. Katnelson, et al., Detection of individual gas molecules adsorbed on graphene, Nat. Mater. 6 (2007) 652–655. [30] S. Hu, M. Lozada-Hidalgo, F.C. Wang, A. Mishchenko, F. Schedin, R.R. Nair, et al., Proton transport through one-atomthick crystals, Nature 516 (2014) 227–230. [31] W. Ren, H.M. Cheng, The global growth of graphene, Nat. Nanotechnol. 9 (2014) 726–730. [32] J.C. Shelton, H.R. Patil, J.M. Blakely, Equilibrium segregation of carbon to a nickel (111) surface: a surface phase transition, Surf. Sci. 43 (1974) 493–520. [33] A.E. Karu, M. Beer, Pyrolytic formation of highly crystalline graphite films, J. Appl. Phys. 37 (1966) 2179–2181. [34] C. Oshima, A. Nagashima, Ultra-thin epitaxial films of graphite and hexagonal boron nitride on solid surfaces, J. Phys. Condens. Matter 9 (1997) 1–20. [35] K.S. Kim, Y. Zhao, H. Jang, S.Y. Lee, J.M. Kim, K.S. Kim, et al., Large-scale pattern growth of graphene films for stretchable transparent electrodes, Nature 457 (2009) 706–710. [36] X. Li, W. Cai, L. Colombo, R.S. Ruoff, Evolution of graphene growth on Ni and Cu by carbon isotope labelling, Nano Lett. 9 (2009) 4268–4272. [37] G.A. Lopez, E.J. Mittemeijer, The solubility of C in solid Cu, Scr. Mater. 51 (2004) 1–5. [38] A.J. Van Bommel, J.E. Crombeen, A. Van Tooren, LEED and Auger electron observations of the SiC(0001) surface, Surf. Sci. 48 (1975) 463–472. [39] Q. Yu, J. Lian, S. Siriponglert, H. Li, Y.P. Chen, S.S. Pei, Graphene segregated on Ni surfaces and transferred to insulators, Appl. Phys. Lett. 93 (2008) 113103. [40] A. Tzalenchuk, S. Lara-Avila, A. Kalaboukhov, S. Paolillo, M. Syväjärvi, R. Yakimova, et al., Towards a quantum resistance standard based on epitaxial graphene, Nat. Nanotechnol. 5 (2010) 186–189. [41] A. Nagashima, Y. Gamou, M. Terai, M. Wakabayashi, C. Oshima, Electronic states of the heteroepitaxial double-layer system: graphite/monolayer hexagonal boron nitride/Ni(111), Phys. Rev. B 54 (1996) 13491–13494. [42] L. Song, L. Ci, H. Lu, P.B. Sorokin, C. Jin, J. Ni, et al., Large scale growth and characterization of atomic hexagonal boron nitride layers, Nano Lett. 10 (2010) 3209–3215. [43] E. Bauer, Surface Microscopy with Low Energy Electrons, Springer-Verlag, New York, 2014. [44] G. Odahara, H. Hibino, N. Nakayama, T. Shimbata, C. Oshima, S. Otani, et al., Macroscopic single-domain graphene growth on polycrystalline nickel surface, Appl. Phys. Express 5 (2012) 035501. [45] H. Hibino, S. Tanabe, S. Mizuno, H. Kageshima, Growth and electronic transport properties of epitaxial graphene on SiC, J. Phys. D Appl. Phys. 45 (2012) 154008. [46] H. Hibino, H. Kageshima, F. Maeda, M. Nagase, Y. Kobayashi, H. Yamaguchi, Microscopic thickness determination of thin graphite films formed on SiC from quantized oscillation in reflectivity of low-energy electrons, Phys. Rev. B 77 (2008) 075413. [47] H. Hibino, H. Kageshima, F. Maeda, M. Nagase, Y. Kobayashi, Y. Kobayashi, et al., Thickness determination of graphene layers formed on SiC using low-energy electron microscopy, e-J. Surf. Sci. Nanotechnol. 6 (2008) 107–110.

174

H. Hibino et al. / Progress in Crystal Growth and Characterization of Materials 62 (2016) 155–176

[48] H. Hibino, S. Mizuno, H. Kageshima, M. Nagase, H. Yamaguchi, Stacking domains of epitaxial few-layer graphene on SiC(0001), Phys. Rev. B 80 (2009) 085406. [49] R. Muñoz, C. Gómez-Aleixandre, Review of CVD synthesis of graphene, Chem. Vap. Deposition 19 (2013) 297–322. [50] W.J. Arnoult, R.B. McLellan, The solubility of carbon in rhodium ruthenium, iridium and rhenium, Scr. Metall. 6 (1972) 1013– 1018. [51] E. Loginova, N.C. Bartelt, P.J. Feibelman, K.F. McCarty, Evidence for graphene growth by C cluster attachment, New J. Phys. 10 (2008) 093026. [52] E. Loginova, N.C. Bartelt, P.J. Feibelman, K.F. McCarty, Factors influencing the growth of graphene on metals, New J. Phys. 11 (2009) 063046. [53] S. Nie, J.M. Wofford, N.C. Bartelt, O.D. Dubon, K.F. McCarty, Origin of the mosaicity in graphene grown on Cu(111), Phys. Rev. B 84 (2011) 155425. [54] J. Coraux, A.T. N’Diaye, M. Engler, C. Busse, D. Wall, N. Buckanie, et al., Growth of graphene on Ir(111), New J. Phys. 11 (2009) 023006. [55] K. Hayashi, A. Yamada, S. Sato, N. Yokoyama, Crystallographic characterization and control of domain structure within individual graphene islands, J. Phys. Chem. C 119 (2015) 4286– 4293. [56] P.C. Rogge, S. Nie, K.F. McCarty, N.C. Bartelt, O.D. Dubon, Orientation-dependent growth mechanisms of graphene islands on Ir(111), Nano Lett. 15 (2015) 170–175. [57] K.F. McCarty, N.C. Bartelt, Graphene growth on metal surfaces, MRS Bull. 37 (2012) 1158–1165. [58] K.F. McCarty, P.J. Feibelman, E. Loginova, N.C. Bartelt, Kinetics and thermodynamics of carbon segregation and graphene growth on Ru(0001), Carbon 47 (2009) 1806– 1813. [59] Q. Yu, L.A. Jauregui, W. Wu, R. Colby, J. Tian, Z. Su, et al., Control and characterization of individual grains and grain boundaries in graphene grown by chemical vapour deposition, Nat. Mater. 10 (2011) 443–449. [60] G.-H. Lee, R.C. Cooper, S.J. An, S. Lee, A. van der Zande, N. Petrone, et al., High-strength chemical-vapor-deposited graphene and grain boundaries, Science 340 (2013) 1073–1076. [61] P.Y. Huang, C.S. Ruiz-Vargas, A.M. van der Zande, W.S. Whitney, M.P. Levendorf, J.W. Kevek, et al., Grains and grain boundaries in single-layer graphene atomic patchwork quilts, Nature 469 (2011) 389–392. [62] S. Wang, S. Suzuki, H. Hibino, Raman spectroscopic investigation of polycrystalline structures of CVD-grown graphene by isotope labelling, Nanoscale 6 (2014) 13838– 13844. [63] T. Ohta, F. El Gabaly, A. Bostwick, J.L. McChesney, K.V. Emtsev, A.K. Schmid, et al., Morphology of graphene thin film growth on SiC(0001), New J. Phys. 10 (2008) 023034. [64] V.N. Strocov, P. Blaha, H.I. Starnberg, M. Rohlfing, R. Claessen, J.-M. Debever, et al., Three-dimensional unoccupied band structure of graphite: very-low-energy electron diffraction and band calculations, Phys. Rev. B 61 (2000) 4994–5001. [65] N. Barrett, E.E. Krasovskii, J.-M. Themlin, V.N. Strocov, Elastic scattering effects in the electron mean free path in a graphite overlayer studied by photoelectron spectroscopy and LEED, Phys. Rev. B 71 (2005) 035427. [66] M. Posternak, A. Baldereschi, A.J. Freeman, E. Wimmer, Prediction of electronic surface states in layered materials: graphite, Phys. Rev. Lett. 52 (1984) 863–866.

[67] N.A.W. Holzwarth, S.G. Louie, S. Rabii, X-ray form factors and the electronic structure of graphite, Phys. Rev. B 26 (1982) 5382– 5390. [68] T. Fauster, F.J. Himpsel, J.E. Fischer, E.W. Plummer, Threedimensional energy band in graphite and lithium-intercalated graphite, Phys. Rev. Lett. 51 (1983) 430–433. [69] V.M. Silkin, J. Zhao, F. Guinea, E.V. Chulkov, P.M. Echenique, H. Petek, Image potential states in graphene, Phys. Rev. B 80 (2009) 121408(R). [70] R.M. Feenstra, N. Srivastava, Q. Gao, M. Widom, B. Diaconescu, T. Ohta, et al., Low-energy electron reflectivity from graphene, Phys. Rev. B 87 (2013) 041406(R). [71] N. Srivastava, Q. Gao, M. Widom, R.M. Feenstra, S. Nie, K.F. McCarty, et al., Low-energy electron reflectivity of graphene on copper and other substrates, Phys. Rev. B 87 (2013) 245414. [72] K.V. Emtsev, F. Speck, T. Seyller, L. Ley, J.D. Riley, Interaction, growth, and ordering of epitaxial graphene on SiC{0001} surfaces: a comparative photoelectron spectroscopy study, Phys. Rev. B 77 (2008) 155303. [73] F. Varchon, P. Mallet, J.-Y. Veuillen, L. Magaud, Ripples in epitaxial graphene on the Si-terminated SiC(0001) surface, Phys. Rev. B 77 (2011) 235412. [74] S. Kim, J. Ihm, H.J. Choi, Y.W. Son, Origin of anomalous electronic structures of epitaxial graphene on silicon carbide, Phys. Rev. Lett. 100 (2008) 176802. [75] C.M. Orofeo, S. Suzuki, H. Kageshima, H. Hibino, Growth and low-energy electron microscopy characterization of monolayer hexagonal boron nitride on epitaxial cobalt, Nano Res. 6 (2013) 335–347. [76] J.M. Campanera, G. Savini, I. Suarez-Martinez, M.I. Heggie, Density functional calculations on the intricacies of Moiré patterns on graphite, Phys. Rev. B 75 (2007) 235449. [77] M.P. Seah, W.A. Dench, Quantitative electron spectroscopy of surfaces – a standard data base for electron inelastic mean free paths in solids, Surf. Interface Anal. 1 (1979) 2–11. [78] M. Aoki, H. Amawashi, Dependence of band structures on stacking and field in layered graphene, Solid State Commun. 142 (2007) 123–127. [79] W. Norimatsu, M. Kusunoki, Formation process of graphene on SiC(0001), Physica E Low Dimens. Syst. Nanostruct. 42 (2010) 691–694. [80] W. Norimatsu, J. Takada, M. Kusunoki, Formation mechanism of graphene layers on SiC(000 1 ) in a high-pressure argon atmosphere, Phys. Rev. B 84 (2011) 035424(R). [81] I. Forbeaux, J.-M. Themlin, J.-M. Debever, Heteroepitaxial graphite on interface formation through conduction-band electronic structure, Phys. Rev. B 58 (1998) 16396–16406. [82] A. Mattausch, O. Pankratov, Ab initio study of graphene on SiC, Phys. Rev. Lett. 99 (2007) 076802. [83] F. Varchon, R. Feng, J. Hass, X. Li, B.N. Nguyen, C. Naud, et al., Electronic structure of epitaxial graphene layers on SiC: effect of the substrate, Phys. Rev. Lett. 99 (2007) 126805. [84] H. Kageshima, H. Hibino, M. Nagase, H. Yamaguchi, Theoretical study of epitaxial graphene growth on SiC(0001) surfaces, Appl. Phys. Express 2 (2009) 065502. [85] P. Lauffer, P. Lauffer, K.V. Emtsev, R. Graupner, T. Sayler, L. Ley, et al., Atomic and electronic structure of few-layer graphene on SiC(0001) studied with scanning tunneling microscopy and spectroscopy, Phys. Rev. B 77 (2008) 155426. [86] J.B. Hannon, M. Copel, R.M. Tromp, Direct measurement of the growth mode of graphene on SiC(0001) and SiC(000 1 ), Phys. Rev. Lett. 107 (2011) 166101.

H. Hibino et al. / Progress in Crystal Growth and Characterization of Materials 62 (2016) 155–176 [87] J.B. Hannon, R.M. Tromp, Pit formation during graphene synthesis on SiC(0001): in situ electron microscopy, Phys. Rev. B 77 (2008) 241404(R). [88] S. Tanaka, K. Morita, H. Hibino, Anisotropic layer-by-layer growth of graphene on vicinal SiC(0001) surfaces, Phys. Rev. B 81 (2010) 041406(R). [89] C. Virojanadara, M. Syväjarvi, R. Yakimova, L.I. Johansson, A.A. Zakharov, T. Balasubramanian, Homogeneous large-area graphene layer growth on 6H-SiC(0001), Phys. Rev. B 78 (2008) 245403. [90] K.V. Emtsev, A. Bostwick, K. Horn, J. Jobst, G.L. Kellogg, L. Ley, et al., Towards wafer-size graphene layers by atmospheric pressure graphitization of silicon carbide, Nat. Mater. 8 (2009) 203–207. [91] H. Kageshima, H. Hibino, H. Yamaguchi, M. Nagase, Stability and reactivity of steps in the initial stage of graphene growth on the SiC(0001) surface, Phys. Rev. B 88 (2013) 235405. [92] M. Hupalo, E.H. Conrad, M.C. Tringides, Growth mechanism for epitaxial graphene on vicinal 6H-SiC(0001) surfaces: a scanning tunneling microscopy study, Phys. Rev. B 80 (2009) 041401(R). [93] T. Ohta, N.C. Bartelt, S. Nie, K. Thürmer, G.L. Kellogg, Role of carbon surface diffusion on the growth of epitaxial graphene on SiC, Phys. Rev. B 81 (2010) 121411(R). [94] S. Kondo, M. Sato, M. Uwaha, H. Hibino, Pattern formation of a step induced by a moving linear source, Phys. Rev. B 84 (2011) 045420. [95] H. Hibino, H. Kageshima, M. Uwaha, Instability of steps during Ga deposition on Si(111), Surf. Sci. 602 (2008) 2421–2426. [96] A. Drabin´ska, K. Grodecki, W. Strupin´ski, R. Boz˙ek, K.P. Korona, A. Wysmołek, et al., Growth kinetics of epitaxial graphene on SiC substrates, Phys. Rev. B 81 (2010) 245410. [97] S. Tanabe, Y. Sekine, H. Kageshima, M. Nagase, H. Hibino, Carrier transport mechanism in graphene on SiC(0001), Phys. Rev. B 84 (2011) 115458. [98] F. Speck, J. Jobst, F. Fromm, M. Ostler, D. Waldmann, M. Hundhausen, et al., The quasi-free-standing nature of graphene on H-saturated SiC(0001), Appl. Phys. Lett. 99 (2011) 122106. [99] C. Riedl, C. Coletti, T. Iwasaki, A.A. Zakharov, U. Starke, Quasifree-standing epitaxial graphene on SiC obtained by hydrogen intercalation, Phys. Rev. Lett. 103 (2009) 246804. [100] C. Virojanadara, A.A. Zakharov, S. Watcharinyanon, R. Yakimova, L.I. Johansson, A low-energy electron microscopy and x-ray photo-emission electron microscopy study of Li intercalated into graphene on SiC(0001), New J. Phys. 12 (2010) 125015. [101] S. Oida, F.R. McFeely, J.B. Hannon, R.M. Tromp, M. Copel, Z. Chen, et al., Decoupling graphene from SiC(0001) via oxidation, Phys. Rev. B 82 (2010) 041411(R). [102] A.L. Walter, K.J. Jeon, A. Bostwick, F. Speck, M. Ostler, T. Seyller, et al., Highly p-doped epitaxial graphene obtained by fluorine intercalation, Appl. Phys. Lett. 98 (2011) 184102. [103] C. Xia, S. Watcharinyanon, A.A. Zakharov, R. Yakimova, L. Hultman, L.I. Johansson, et al., Si intercalation/deintercalation of graphene on 6H-SiC(0001), Phys. Rev. B 85 (2012) 045418. [104] K.V. Emtsev, A.A. Zakharov, C. Coletti, S. Forti, U. Starke, Ambipolar doping in quasifree epitaxial graphene on SiC(0001) controlled by Ge intercalation, Phys. Rev. B 84 (2011) 125423. [105] I. Gierz, T. Suzuki, R.T. Weitz, D.S. Lee, B. Krauss, C. Riedl, et al., Electronic decoupling of an epitaxial graphene monolayer by gold intercalation, Phys. Rev. B 81 (2010) 235408. [106] S. Tanabe, M. Takamura, Y. Harada, H. Kageshima, H. Hibino, Effects of hydrogen intercalation on transport properties of

[107]

[108]

[109]

[110]

[111]

[112]

[113]

[114]

[115]

[116]

[117]

[118]

[119]

[120]

[121]

[122]

175

quasi-free-standing monolayer graphene, Jpn J. Appl. Phys. 53 (2014) 04EN01. J.A. Robinson, M. Hollander, M. LaBella, K.A. Trumbull, R. Cavalero, D.W. Snyder, Epitaxial graphene transistors: enhancing performance via hydrogen intercalation, Nano Lett. 11 (2011) 3875–3880. Y. Murata, T. Mashoff, M. Takamura, S. Tanabe, H. Hibino, F. Beltram, et al., Correlation between morphology and transport properties of quasi-free-standing monolayer graphene, Appl. Phys. Lett. 105 (2014) 221604. K. Yamasue, H. Fukidome, K. Funakubo, M. Suemitsu, Y. Cho, Interfacial charge states in graphene on SiC studied by noncontact scanning nonlinear dielectric potentiometry, Phys. Rev. Lett. 114 (2015) 226103. S. Tanabe, Y. Sekine, H. Kageshima, H. Hibino, Electrical characterization of bilayer graphene formed by hydrogen intercalation of monolayer graphene on SiC(0001), Jpn J. Appl. Phys. 51 (2012) 02BN02. K. Lee, S. Kim, M.S. Points, T.E. Beechem, T. Ohta, E. Tutuc, Magnetotransport properties of quasi-free-standing epitaxial graphene bilayer on SiC: evidence for Bernal stacking, Nano Lett. 11 (2011) 3624–3628. J. Lahiri, Y. Lin, P. Bozkurt, I.I. Oleynik, M. Batzill, An extended defect in graphene as a metallic wire, Nat. Nanotechnol. 5 (2010) 326. Y. Ogawa, B. Hu, C.M. Orofeo, M. Tsuji, K. Ikeda, S. Mizuno, et al., Domain structure and boundary in single-layer graphene grown on Cu(111) and Cu(100) films, J. Phys. Chem. Lett. 3 (2012) 219–226. H. Ago, Y. Ohta, H. Hibino, D. Yoshimura, R. Takizawa, Y. Uchida, et al., Growth dynamics of single-layer graphene on epitaxial Cu surfaces, Chem. Mater. 27 (2015) 5377–5385. X. Li, C.W. Magnuson, A. Venugopal, R.M. Tromp, J.B. Hannon, E.M. Vogel, et al., Large-area graphene single crystals grown by low-pressure chemical vapor deposition of methane on copper, J. Am. Chem. Soc. 133 (2011) 2816–2819. S. Chen, H. Ji, H. Chou, Q. Li, H. Li, J.W. Suk, et al., Millimetersize single-crystal graphene by suppressing evaporative loss of Cu during low pressure chemical vapor deposition, Adv. Mater. 25 (2013) 2062–2065. Z. Yan, J. Lin, Z. Peng, Z. Sun, Y. Zhu, L. Li, et al., Toward the synthesis of wafer-scale single-crystal graphene on copper foils, ACS Nano 6 (2012) 9110–9117. A. Mohsin, L. Liu, P.Z. Liu, W. Deng, I.N. Ivanov, G. Li, et al., Synthesis of millimeter-size hexagon-shaped graphene single crystals on resolidified copper, ACS Nano 7 (2013) 8924– 8931. Y. Hao, M.S. Bharathi, L. Wang, Y. Liu, H. Chen, S. Nie, et al., The role of surface oxygen in the growth of large singlecrystal graphene on copper, Science 342 (2013) 720– 723. H. Zhou, W.J. Yu, L. Liu, R. Cheng, Y. Chen, X. Huang, et al., Chemical vapour deposition growth of large single crystals of monolayer and bilayer graphene, Nat. Commun. 4 (2013) 2096. S. Wang, H. Hibino, S. Suzuki, H. Yamamoto, Atmospheric pressure chemical vapor deposition growth of millimeter-scale singlecrystalline graphene on the copper surface with native oxide layer, to be published, 2016. K. Yan, H. Peng, Y. Zhou, H. Li, Z. Liu, Formation of bilayer Bernal graphene: layer-by-layer epitaxy via chemical vapor deposition, Nano Lett. 11 (2011) 1106–1110.

176

H. Hibino et al. / Progress in Crystal Growth and Characterization of Materials 62 (2016) 155–176

[123] L. Liu, H. Zhou, R. Cheng, W.J. Yu, Y. Liu, Y. Chen, et al., Highyield chemical vapor deposition growth of high-quality largearea AB-stacked bilayer graphene, ACS Nano 6 (2012) 8241– 8249. [124] S. Chen, W. Cai, R.D. Piner, J.W. Suk, Y. Wu, Y. Ren, et al., Synthesis and characterization of large-area graphene and graphite films on commercial Cu-Ni alloy foils, Nano Lett. 11 (2011) 3519–3525. [125] N. Liu, L. Fu, B. Dai, K. Yan, X. Liu, R. Zhao, et al., Universal segregation growth approach to wafer-size graphene from non-noble metals, Nano Lett. 11 (2011) 297– 303. [126] W. Liu, S. Kraemer, D. Sarkar, H. Li, P.M. Ajayan, K. Banerjee, Controllable and rapid synthesis of high-quality and large-area Bernal stacked bilayer graphene using chemical vapor deposition, Chem. Mater. 26 (2014) 907–915. [127] W. Auwärter, H.U. Suter, H. Sachdev, T. Greber, Synthesis of one monolayer of hexagonal boron nitride on Ni(111) from B-trichloroborazine (ClBNH)3, Chem. Mater. 16 (2004) 343– 345. [128] W. Auwärter, M. Muntwiler, J. Osterwalder, T. Greber, Defect lines and two-domain structure of hexagonal boron nitride films on Ni(111), Surf. Sci. 545 (2003) L735–L740. [129] G.B. Grad, P. Blaha, K. Schwarz, W. Auwärter, T. Greber, Density functional theory investigation of the geometric and spintronic structure of h-BN/Ni(111) in view of photoemission and STM experiments, Phys. Rev. B 68 (2003) 085404.

[130] N. Joshi, P. Ghosh, Substrate-induced changes in the magnetic and electronic properties of hexagonal boron nitride, Phys. Rev. B 87 (2013) 235440. [131] Y. Liu, S. Bhowmick, B.I. Yakobson, BN white graphene with “colorful” edges: the energies and morphology, Nano Lett. 11 (2011) 3113–3116. [132] Y. Stehle, H.M. Meyer III, R.R. Unocic, M. Kidder, G. Polizos, P.G. Datskos, et al., Synthesis of hexagonal boron nitride monolayer: control of nucleation and crystal morphology, Chem. Mater. 27 (2015) 8041–8047. [133] R. Zhao, F. Li, Z. Liu, Z. Liu, F. Ding, The transition metal surface passivated edges of hexagonal boron nitride (h-BN) and the mechanism of h-BN’s chemical vapor deposition (CVD) growth, Phys. Chem. Chem. Phys. 17 (2015) 29327–29334. [134] T. Kilner, A.J. Dempsey, R.M. Pilliar, G.C. Weatherly, The effects of nitrogen additions to a cobalt-chromium surgical implant alloy, J. Mater. Sci. 22 (1987) 565–574. [135] F. Orlando, P. Lacovig, L. Omiciuolo, N.G. Apostol, R. Larciprete, A. Baraldi, et al., Epitaxial growth of a single-domain hexagonal boron nitride monolayer, ACS Nano 8 (2014) 12063– 12070. [136] X. Song, J. Gao, Y. Nie, T. Gao, J. Sun, D. Ma, et al., Chemical vapor deposition growth of large-scale hexagonal boron nitride with controllable orientation, Nano Res. 8 (2015) 3164–3176. [137] G. Lu, T. Wu, Q. Yuan, H. Wang, H. Wang, F. Ding, et al., Synthesis of large single-crystal hexagonal boron nitride grains on Cu-Ni alloy, Nat. Commun. 6 (2015) 6160.