ARTICLE pubs.acs.org/JPCC
Fluorinating Hexagonal Boron Nitride/Graphene Multilayers into Hybrid Diamondlike Nanofilms with Tunable Energy Gap Zhuhua Zhang,*,†,‡ Xiao Cheng Zeng,*,‡ and Wanlin Guo*,† †
Key Laboratory of Intelligent Nano Materials and Devices (Ministry of Education) and Institute of Nano Science, Nanjing University of Aeronautics and Astronautics, Nanjing 210016, China ‡ Department of Chemistry and Center for Materials & Nanoscience, University of Nebraska—Lincoln, Lincoln, Nebraska 68588
bS Supporting Information ABSTRACT: Using ab initio calculations and quantum molecular dynamics simulations, we demonstrate that a few layers of graphene sandwiched between hexagonal boron nitride (h-BN) layers can undergo spontaneous transformation into hybrid cubic BNdiamond (c-BN/Dmd) nanofilms upon fluorination. This spontaneous transformation stems from the remarkably higher stability of thin c-BN/Dmd nanofilm with sp3 hybridization over the precursor multilayer with sp2 hybridization and is promoted by strong selectivity of fluorination with the boron atoms of the coating BN layers. Upon increasing the total number of multilayers, however, the transformation is no longer spontaneous due to emergence of the energy barrier. Nevertheless, adding more h-BN layers to the hybrid nanofilm can assist the transformation into c-BN/Dmd nanofilms upon fluorination. The electronic properties of the c-BN/ Dmd nanofilms can be tuned by controlling the ratio of the BN component and film thickness, which can yield narrow-gap semiconductors for novel electronic applications. In addition, the energy gap in the nanofilms can be modulated linearly by applying external electric fields.
’ INTRODUCTION Diamond possesses a number of exceptional properties,1 such as the highest known hardness, the lowest coefficient of thermal expansion, and high thermal conductivity at room temperature, as well as high optical transparency in a wide wavelength range. However, developing diamond-based devices of widespread use is still not feasible, as some limitations remain in diamond’s synthesis and its electronic properties. First, most diamond films made by state-of-the-art chemical vapor deposition are polycrystalline, and graphite-like structures are usually present at the grain boundaries between the diamond crystallites. These structural flaws greatly degrade the properties of diamond films. Second, diamond can turn into a p-type semiconductor via boron substitution,2 but making an n-type doped diamond is experimentally still a challenge.3 On the other hand, cubic boron nitride (c-BN), with the same lattice structure as diamond, not only possesses desirable mechanical and thermal properties (e.g., its hardness and thermal conductivity are only second to those of diamond) but also is a superior electronic material than diamond due to its wider band gap, greater chemical inertness, and higher oxidation stability. Moreover, there exist well-developed processes to achieve p- and n-type doping in c-BN,2,4,5 contrary to diamond. Hence, to develop new materials with combined and superior mechanical, chemical, and electronic properties, it is desirable to make heterogeneous nanofilms that integrate diamond with c-BN. As the c-BN has a similar lattice constant as diamond, there is r 2011 American Chemical Society
currently widespread interest in growing c-BN nanofilm on diamond surfaces.49 Yet the interface between c-BN and diamond often contains amorphous or turbostratic BN layers, which actually makes the c-BN and diamond decouple both electronically and mechanically.10,11 Still, how to fabricate a hybrid c-BN/ diamond nanofilm with covalent interface and uniform crystallization remains an open question. Recent experimental and theoretical studies have shown that hydrogenation of a suspended graphene monolayer can transform it into a graphane layer with all the carbon atoms being sp3 hybridized.12,13 The same transition is also realized in graphene upon fluorination.1417 Especially, Zboril et al. demonstrate for the first time the transformation of fluorographene with diamond structure from graphene. These results implicate a possible path for making high-quality diamond nanofilms by functionalizing suspended graphene multilayers (MLs). Unfortunately, hydrogen or fluorine atoms on the ML graphene tend to form clusters, due to their preferential sticking into specific adsorbate structures.1820 This is certainly unfavorable to the making of diamond films because of the lack of carbon radicals for the interlayer bonding. Structurally similar to the graphene, hexagonal BN (h-BN) monolayers have also been fabricated experimentally.21,22 Received: July 27, 2011 Revised: September 18, 2011 Published: October 03, 2011 21678
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The Journal of Physical Chemistry C Contrary to graphene, adsorbates on the h-BN monolayer exhibit strong site selectivity due to the ionicity of BN bonds. This is particularly the case for fluorine atoms, which prefer to bond with boron atoms but cannot form stable bonds with nitrogen atoms in h-BN sheets.2326 On the basis of this unique feature and the well-developed experimental technologies of assembling h-BN/ graphene multilayers,2729 we propose a ML structure with a few layers of graphene sandwiched between two h-BN sheets (BNCBN). Energetics calculations prove that the BNC BN ML is increasingly less stable than the corresponding hybrid c-BN/Dmd nanofilms as the number of MLs decreases. Then using ab initio molecular dynamics (AIMD) simulations and calculations, we manifest that fluorination of thin BNCBN MLs is feasible to form hybrid c-BN/Dmd nanofilms wherein the crystallization orientation in both parts is along the [111] direction. The propensity of fluorination-induced transition is weakened with increasing the number of MLs due to the emergence of an energy barrier. However, increasing the ratio of h-BN layers in a BNCBN ML can significantly reduce the energy barrier or even remove it. Importantly, the fabricated nanofilms exhibit tunable electronic properties useful for nanoelectronics applications.
’ COMPUTATIONAL METHODS All the calculations are performed by using the Vienna Ab-initio Simulation Package (VASP).30,31 The projector-augmented wave method32 is employed for the core region, with a kinetic energy cutoff being set as 500 eV in the plane-wave expansion. It is known that semilocal generalized gradient approximation (GGA) fails to produce the interlayer distance in graphite or h-BN sheets, while the local density approximation (LDA) can give reasonable interlayer distances in these systems due to a delicate error cancellation between exchange and correlation. So we employ LDA for the exchange correlation potential. Actually, we have performed test calculations using GGA with the inclusion of van der Waals interactions and do not find any qualitative change compared to LDA calculated results. An initial lattice constant of 1.44 Å and 2 2 graphene supercell are used to calculate the total energy and energy barrier. The two-dimensional Brillouin zone integration is sampled by 50 special k-points for structural optimization and 60 k-points for electronic structure calculations. The minimum energy path (MEP) is mapped out using the climbing image nudged elastic band method.33,34 Both the atomic positions and supercell shape are allowed to relax until the force on each atom is less than 0.01 eV/Å. A vacuum region of 15 Å is set between two adjacent MLs. The AIMD simulations are carried out in the canonical ensemble. In the AIMD simulations, a 4 4 supercell is used and the time step is set to 0.5 fs. The external electric field is simulated by the planar dipole layer method as implemented in VASP. ’ RESULTS AND DISCUSSION In the BNCBN ML, the graphene layers and h-BN layers adopt an ABC (i.e., rhombohedral) stacking. In our models, this stacking is less stable than the AB stacking by smaller than 1 meV per unit cell. In the graphene/BN interfaces, we let the nitrogen atom be above a carbon atom, and the boron atom be located above the center of a carbon hexagon. This interfacial structure is found to be energetically the most favorable when the BN CBN ML is fluorinated. For convenience, we denote the ML containing m-layers of graphene and n-layers of BN sheets as
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Figure 1. (a) Schematic structural motifs for CmBN2-ML and the CmBN2-Dmd nanofilm. (b) Formation energies of CmBN2-ML and the CmBN2-Dmd nanofilm as a function of m. The formation energies for the nanofilms at different μF are presented by dash lines.
CmBNn-ML and the fluorinated ML as F-CmBNn-ML. Correspondingly, the hybrid c-BN/Dmd nanofilm is denoted as CmBNn-Dmd nanofilm (see Figure 1a). First, we study the case of n = 2, the minimum number for the h-BN layers in our model. To examine the feasibility of the proposed synthesis strategy, we calculate the formation energies for the CmBN2-ML and the CmBN2-Dmd nanofilm. The formation energy is defined as Ef ¼ Ecoh nC μC nB μB nN μN nF μF
ð1Þ
where Ecoh is the cohesive energy per atom in the ML or nanofilms, ni (i = B,N,C, F) is the number of i atoms in the supercell, and μi is the corresponding chemical potential. As the precursor materials consist wholly of graphene layers and BN sheets, we choose μC as the cohesive energy per atom of a single graphene sheet and μB + μN as the cohesive energy per BN pair of the h-BN sheet. The definition of formation energy allows a direct stability comparison between systems with different chemical compositions.35,36 Figure 1b presents formation energies of the CmBN2ML and CmBN2-Dmd nanofilm as functions of m at different μF. The red solid line corresponds to the condition that μF adopts the cohesive energy per atom in a fluorine molecule. It is shown that the formation energy for the CmBN2-Dmd nanofilm increases rapidly with increasing thickness, while that for the CmBN2-ML decreases very slowly. Within our calculated range, the CmBN2Dmd nanofilm has lower formation energy than the corresponding CmBN2-ML, particularly in the case of thinner nanofilms. This suggests that thin CmBN2-Dmd nanofilms are remarkably more stable than the precursor ML. Note that including the correction of van der Waals interaction energy will reduce the formation energy of the ML by 0.030.04 eV/atom, which does not essentially change our results. The difference in formation energy between CmBN2-ML and CmBN2-Dmd is also sensitive to the change in μF: higher μF gives larger formation energy difference. Here, the higher stability of a thinner CmBN2-Dmd nanofilm over the CmBN2-ML is due to the increased ratio of surface atoms, which greatly reduces the surface energy under fluorinated condition in light of the binding energy of FB bond being larger than that of the FF bond. Our finding qualitatively agrees with previous reports using a semiempirical quantum mechanical method, in which H-terminated nanosized diamond clusters are found to be more stable than polycyclic aromatics.37 Also, our results are consistent with a recent ab initio study on hydrogenated graphene, which yields the stability hierarchy: graphane > bilayer graphane > diamond.38 In view 21679
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Figure 2. (a) Total energy variation for the F-C1BN2-ML during the structural relaxation. Band structures for three snapshots, (b) A, (c) B, and (d) C, in the relaxation pathway, together with the charge density for the electronic states within the cyan shaded parts of the band structures, as well as the corresponding total charge densities.
of the higher stability, formation of CmBN2-Dmd nanofilm from thin CmBN2-ML is energetically very feasible. In what follows, we specially examine the transformation process of the CmBN2-ML upon fluorination. It is found that that the F-C1BN2-ML can transform spontaneously into a C1BN2Dmd nanofilm via structural relaxation (see movie 1 of the Supporting Information). The diamond nanofilm has identical bond structure to fluorine-terminated diamond (111) film but has a chemically more robust BN coating. The energy profile for the transition from the F-C1BN2-ML to C1BN2-Dmd is shown in Figure 2a, together with band structures and charge densities corresponding to electronic states around the Fermi level for three chosen images A, B, and C along the profile (Figure 2bd). In the image A, the two BN layers show zigzag buckling upon chemisorption of fluorine atoms on the boron atoms. Meanwhile, active dangling bond states are created around the Fermi level due to the breaking of π bonds and they distribute on nitrogen atoms. This behavior means that the high chemical reactivity of fluorine is transferred to the nitrogen atoms, which now becomes active radicals; see Figure 2b. Note that the system is nonmagnetic, contrary to a freestanding semifluorinated BN layer, which has a ferromagnetic ground state.39 Hereafter we will not consider spin-polarization in calculations. From the image B, the interlayer bonds are formed between the nitrogen atoms of the upper h-BN and the carbon atoms in a sublattice of the graphene (Figure 2c), attributed to a more favorable stacking between them. Then the radical sites are further transferred to the carbon atoms in another graphene sublattice. On the other side, the charge density of radicals on the lower BN layer starts to overlap with the radicals on graphene. As the relaxation goes on, these radical sites disappear upon forming NC bonds among them, and the system has transformed into a new semiconductor; see Figure 2d. This spontaneous transformation also occurs in other stacking modes of F-C1BN2-ML, e.g. the AB stacking. In the transformation demonstrated above, a question arises as to what drives the formation of interlayer bonds in the F-C1BN2ML. To clarify this issue, we first examine the interlayer interaction between the semifluorinated BN layers and the graphene layer by plotting the charge redistribution in the image A (see Figure 3a). It is shown that a dramatic charge transfer occurs from the graphene to both the semifluorinated BN layers, due to the strong chemical activity of the nitrogen radicals in the BN
Figure 3. (a) Charge redistribution upon formation of fluorinated BN and graphene interface in the image A (Figure 2b), ΔF = Fsystem FF‑BN FC, where FF‑BN, FC, and Fsystem, represent the total charge densities of the two fluorinated BN sheets, the graphene layer, and the combined structure of image A, respectively. The contour spacing is set to be 2 103 e/Å3. (b) Average local density of states on the nitrogen and carbon atoms in the image A, as well as those in an individual semifluorinated BN sheet and graphene.
layers. As a result, the BN layers are negatively charged while the graphene is positively charged, thus forming a strong Coulomb interaction between them. This Coulomb interaction becomes more pronounced with decreasing the grapheneBN distance and is the major driving force in reducing the distance between the semifluorinated BN layers and graphene. Moreover, the electron depletion in the graphene layer, to some extent, attenuates the bond strength therein, thereby making it easier to be buckled for interlayer bonding. To get more insight into the interlayer interaction, we plot the local density of states (LDOS) for the nitrogen and carbon atoms of the image A, together with the corresponding LDOS in an individual semifluorinated BN sheet and graphene for comparison (Figure 3b). Clearly, upon formation of the interface between the semifluorinated BN and graphene, the electron population of nitrogen atoms in the fluorinated BN sheet increases, consistent with the above charge redistribution. More importantly, there are some resonant peaks in the LDOS of the nitrogen and carbon atoms that are absent prior to the formation of the interface. These resonant peaks are indicative of the remarkable hybridization between the graphene and semifluorinated BN sheets and therefore hint that a weak 21680
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Figure 4. Calculated MEP from the (a) optimized F-C3BN2-ML, (b) F-C2BN3-ML, and (c) energy profile of C1BN4-ML by relaxation, to the corresponding hybrid diamond nanofilms. Insets illustrate the geometry structures of the initial state (I), transition state (T), and final state (F) in the MEP. The charge redistributions like that in Figure 3a are also given for the structures I in parts a and c.
interlayer bonding interaction has been formed in the image A, even though the interplanar distance is still at van der Waals distance. We thus conclude that it is the strong reactivity of nitrogen radials changing the interlayer interaction between the BN sheet and graphene and driving the bond formation at the interface. With increasing the number m in the CmBN2-ML, the hybrid diamond film cannot be achieved directly by structural relaxation. As an example, the relaxed structure of F-C3BN2-ML remains a ML structure (see inset I of Figure 4a). In the ML structure, the average interlayer spacing between the semifluorinated h-BN sheet and the adjacent graphene layer is 2.89 Å and the interlayer distance between the interior graphene layers is 3.2 Å, both reduced from 3.35 Å prior to fluorination. However, the total energy of the optimized F-C3BN2-ML is much higher than the C3BN2-Dmd by 4.68 eV per unit cell. So the relaxed F-C3BN2ML is metastable and an energy barrier must exist to hinder the transition from the F-C3BN2-ML to the hybrid diamond film. We perform nudged elastic band calculations to compute the MEP for the transition between the relaxed F-C3BN2-ML and C3BN2Dmd film, as shown in Figure 4a. The MEP shows an energy barrier of 0.18 eV. This is understandable, because the increased graphene layers will cost more energy to form zigzag buckling to achieve the sp3 hybridization. From the electronic point of view, the charge transfer from the graphene to the semifluorinated BN layers is greatly reduced from that in the F-C1BN2-ML because the electron depletion region distributes only around one side of the graphene layer (see Inset I, Figure 4c). As a result, the bond strength of graphene layers in the F-C3BN2-ML remains sufficiently strong to hinder buckling, as happened in the F-C1BN2ML. The buckling-induced barrier is also corroborated by the atomic structure of transition-state T displayed in the inset of Figure 4a, where the three graphene layers show slight zigzag deformation toward forming interlayer bonds. Once the barrier is overcome, the system would rapidly transform into a hybrid diamond film as the ground state. Our extensive calculations show that the energy barrier increases with increasing m; it disappears for m < 2 and is 0.02 eV for m = 2. Some intermediate states arise in the MEP for m > 3, resulting in two or more energy barriers. For example, there are three barriers and two intermediate states for the F-C6BN2-ML transforming into the
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C6BN2-Dmd film. Hence, the thicker F-CmBN2-ML is more difficult to transform into the hybrid diamond film. Besides MEP calculations, we also perform AIMD simulations to further demonstrate the formation of the hybrid diamond films. Since the energy barrier for the transformation of F-C3BN2-ML is merely 0.18 eV, we set a temperature of 300 K in the AIMD simulation. To assist the system approaching to the stable state, we also anneal the system to 20 K with a duration time of 5 ps. We find that the system transforms into the hybrid diamond film at ∼0.35 ps (see movie 2 of the Supporting Information). Once the diamond film is formed, it can be highly stable, even at 2000 K over 5 ps of AIMD simulation. However, we do not observe the transition for the thicker F-C6BN2-ML in a single AIMD trajectory. Instead, amorphous structures near the ML surfaces are formed and the system remains a layered structure. The latter becomes energetically more and more favorable as m increases. According to the AIMD simulation, the critical thickness for the formation of CmBN2-Dmd film is m = 4. On the other hand, we find that increasing the number n in the F-CmBNn-ML (for fixed m + n) can reduce the energy barrier. The MEP for the F-C2BN3-ML to C2BN3-Dmd is shown in Figure 4b, where the energy barrier is only 0.10 eV. In the initial structure (I) of the F-C2BN3-ML, the upper two h-BN layers form a c-BN-like structure, while the two graphene layers remain sp2 hybridized (see inset of Figure 4b). So the transition of F-C2BN3-ML is somewhat similar to that of the F-C2BN2-ML. When n is increased to 4, the energy barrier for F-C1BN4-ML even disappears and the formation of C1BN4-Dmd film is spontaneous (Figure 4c), as in the case of F-C1BN2-ML. We also alter the relative position of the graphene monolayer in the F-C1BN4-ML and find that the spontaneous transition is very robust. With further increasing n to 5, the F-BN5-ML can spontaneously transform into a c-BN (111) film, but with a layer of homogeneous NN bonds. This structure is 1.54 eV/cell more stable than a polymorph of the c-BN film wherein the two surfaces are terminated by fluorinated nitrogen and boron atoms, respectively, further underlining the strong reactive selectivity of fluorine atoms with boron atoms in BN structures, even on the surface of c-BN. Overall, these results suggest that the formation of interlayer bonds is more favorable in h-BN than in graphene under fluorination. This is due to the localized character of BN electronic states, which makes the depleted electrons in the h-BN layer mainly concentrated on the nitrogen atoms, as illustrated by the charge redistribution shown in the inset I of Figure 4c, unlike the case in graphene, where all the carbon atoms have electrons lost. Moreover, the BN bonds are inherently more flexible than the CC bonds, and therefore, the BN sheet is easier to buckle than graphene for forming interlayer bonds. Experimentally, we expect that the synthesis of c-BN nanofilms can be realized by heating the mixture of h-BN/graphene multilayers and XeF2, as done for synthesizing fluorographene.1416 It is exciting that a very recent experiment has synthesized a 2D hydroxylated diamond layer at room temperature by compressing graphene multilayer with the topmost layers covered with hydroxyl groups.40 This experimental advance further convinces us of the feasibility of the proposed diamond synthesis. However, we show that the transition to the hybrid diamond nanofilm is still less feasible in the F-CmBNn-ML for m + n > 8, even upon varying m. In this case, applying a vertical pressure onto the BN plane is necessary, which can reduce the interlayer distance among the ML to favor the interlayer bond formation. We have 21681
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The Journal of Physical Chemistry C estimated the pressure triggering the F-C6BN2-ML transformation to be 3.2 GPa. Once these hybrid nanofilms are synthesized, it is interesting to investigate the electronic properties of the hybrid diamond films. The band structures for the C3BN2-Dmd, C2BN3-Dmd, and C1BN4-Dmd films are shown in parts a, d, and g of Figures 5, respectively. All the hybrid diamond nanofilms are semiconducting with an indirect energy gap. Interestingly, we find that the energy gap of the nanofilms can be modulated widely upon changing the ratio of the BN component. The calculated energy gaps (based on LDA) of C3BN2-, C2BN3-, and C1BN4-Dmd films are 4.1, 2.3, and 0.64 eV, respectively. The conduction band minimum
Figure 5. (a) Band structure for the C3BN2-Dmd nanofilm. (b) Isosurface plot (0.005 e/Å3) of partial charge densities for the VBM (cyan) and CBM (magenta) of the C3BN2-Dmd nanofilm. (c) Plane-averaged electrostatic potential along the normal of the C3BN2-Dmd nanofilm. The cyan shades guide the trend of potential variation across the nanofilm. Corresponding cases for the C2BN3-Dmd (df) and C1BN4-Dmd (gi) nanofilms are shown, respectively. The dash rectangles mark the diamond layers within the nanofilms.
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(CBM) of the hybrid diamond nanofilms originates from the σ antibonding states in the carbon layer, while the valence band maximum (VBM) is from the 2p states of nitrogen and fluorine atoms, like a surface state (see Figure 5b,e,h). The hole and electron states therefore are well-separated, especially in the structurally asymmetrical C2BN3- and C1BN4-Dmd films. This charge character offers a potential application in developing solar cells. To understand the reduced gap in the CmBNn-Dmd nanofilms with increasing n, we plot the plane averaged electrostatic potentials along the normal of the three nanofilms, as shown in Figure 5c,f,i. It is found that the electrostatic potential shows a minimum within the region of the diamond section and a maximum at the film surface where VBM is located, indicating a local electric potential difference in these nanofilms. The built-in potential difference is due to different site potentials of boron and nitrogen atoms and increases with increasing the distance between CBM and VBM planes, which is identical to the thickness of the thicker c-BN section in the film. Upon this potential difference, the energy level of the CBM is shifted downward while that of the VBM is shifted upward, much like the Stark effect. Consequently, the energy gap of the CmBNn-Dmd film is reduced depending on the potential difference. The local potential difference is 2.2 eV in the C3BN2-Dmd film and is increased to 6.87 eV in the C1BN4Dmd film, thereby leading to a greatly reduced gap in the latter. On the basis of this analysis, the energy gap of a hybrid diamond film should also depend on the relative positions of the carbon layers in the film, and it will increase with moving the carbon layer away from the film surfaces. As an example, the energy gap of C1BN4-Dmd nanofilm is increased to 2.6 eV when the carbon layer is moved to the middle layer of the nanofilm, because the distance between the CBM and VBM planes now is shortened so that the built-in potential difference is reduced. Since the energy gap of the CmBNn-Dmd film is determined by the built-in potential difference, it should also strongly depend on the film thickness. The energy gap of the CmBNn-Dmd nanofilms of different thicknesses as a function of m is shown in Figure 6a. Generally, the energy gap decreases rapidly with increasing thickness at a given m, but at a given n, it decreases very slowly. This further quantifies the importance of the BN component in determining the gap of the hybrid nanofilms as it determines the value of built-in electric polarization. In this regard, the variable magnitude in the energy gap upon changing the ratio of BN component also increases with increasing the film thickness. Especially, the gaps in all C1BNn-Dmd films are closed when n > 4, although these films are sp3 bonded. We note that our LDA calculations underestimate the energy gap of the nanofilms, but
Figure 6. (a) Energy gap in the F-CmBNn nanofilm of different thicknesses as a function of m. (b) Electric-field-induced modulation on energy gap of the CmBNn-Dmd nanofilms with m + n = 5. The electric field is applied along the normal of the nanofilm, as shown in the inset, with the positive direction indicated by the big arrow. 21682
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The Journal of Physical Chemistry C the trend of gap variation with thickness should be reliable, because this phenomenon is akin to the widely reported gap modulation in other BN systems by external electric fields.4143 Motivated by the dependence of energy gap on the built-in potential difference, we expect that the energy gap of the CmBNnDmd films exhibits strong response to applied external electric field, which is able to modify the potential difference. We therefore apply an electric field along the normal of the film and define its positive direction pointing from the CBM plane to the VBM plane. Figure 6b presents the energy gaps of the CmBNn-Dmd films with m + n = 5 as functions of applied field strength. As expected, the energy gap of these nanofilms can be efficiently modulated by the electric field. The energy gaps in both the C2BN3-Dmd and C1BN4-Dmd films are reduced by negative fields but increased by positive fields, rendering a linear Stark effect, whereas the energy gap in the C3BN2-Dmd film linearly decreases with increasing field strength, regardless of the field direction. This difference is due to the structural symmetry of the C3BN2-Dmd film, contrasting with the C2BN3-Dmd and C1BN4Dmd films whereby a nonzero built-in electric polarization is developed across the film thickness. The mechanism for the fielddependent energy gap is just that the built-in potential difference is linearly modified by the applied electric field, which leads to linear band shifts in these nanofilms. To our knowledge, a fieldinduced linear gap modulation within such a wide range of field strength has been rarely reported. The linear gap modulation by a variable external electric field is paramount to precise control on the energy gap and will greatly facilitate the practical design of devices based on the hybrid nanofilms.
’ CONCLUSIONS Systematical ab initio calculations and AIMD simulations have demonstrated that hybrid diamond nanofilms can be synthesized by fluorinating an h-BN/graphene multilayer. The essential prerequisites for this transformation are the fluorination selectivity of the coating BN layers and associated remarkably higher stability of the hybrid diamond nanofilms than the BNCBN multilayers. This transformation is spontaneous for a sufficiently thin multilayer and will suffer an energy barrier with increasing the total number of multilayers. Interestingly, increasing the ratio of the h-BN layer in the initial multilayer can greatly reduce the energy barrier to promote the transformation of the hybrid diamond film. The electronic properties of the hybrid nanofilms can be controlled continuously from insulator to narrow-gap semiconductor by increasing the ratio of the BN component or film thickness. Moreover, the energy gap in the nanofilms can be efficiently modulated in a decent linear manner by applying a vertical electric field. As the proposed synthesis strategy is viable with current experimental technologies, the hybrid diamond nanofilms will offer a series of new features that are unavailable from pure diamond films and could find a wealth of new opportunities in nanoelectronics and photovoltaics applications. ’ ASSOCIATED CONTENT
bS
Supporting Information. Movies 1 and 2. This material is available free of charge via the Internet at http://pubs.acs.org.
’ AUTHOR INFORMATION Corresponding Author
*E-mail:
[email protected] (Z.Z);
[email protected] (X.C.Z.);
[email protected] (W.G.).
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’ ACKNOWLEDGMENT Z.Z. acknowledges the support by the National NSF (11172124), the NUAA Research Fund (4015-YAH10043), Jiangsu Province NSF (BK2011722) and National and Jiangsu Postdoctoral Research Foundation (20110490132, 1002015B). W.G. is supported by the 973 Program (2007CB936204), National NSF (10732040, 91023026) and Jiangsu Province NSF (BK2008042) of China. X.C.Z. is supported by ARL (W911NF1020099), NSF (DMR-0820521), and ONR(N0001409-1-0943), and by the University of Nebraska’s Holland Computing Center. ’ REFERENCES (1) May, P. W. Science 2008, 319, 1490–1491. (2) Kalish, R. Diamond Relat. Mater. 2001, 10, 1749–1755. (3) Lewis, N. S. Science 2007, 315, 798–801. (4) Mishima, O.; Tanaka, J.; Yamaoka, S.; Fukunaga, O. Science 1987, 238, 181–183. (5) Pryor, R. W. Appl. Phys. Lett. 1996, 68, 1802–1804. (6) Zhang, W.; Bello, I.; Lifshitz, Y.; Chan, K. M.; Meng, X.; Wu, Y.; Chan, C. Y.; Lee, S. T. Adv. Mater. 2004, 16, 1405–1408. (7) Zhang, X. W.; Boyen, H.-G.; Deyneka, N.; Ziemann, P.; Banhart, F.; Schreck, M. Nat. Mater. 2003, 2, 312–315. (8) Lux, B.; Kalss, W.; Haubner, R.; Taniguchi, T. Diamond Relat. Mater. 1999, 8, 415–422. (9) Chong, Y. M.; Ma, K. L.; Leung, K. M.; Chan, C. Y.; Ye, Q.; Bello, I.; Zhang, W.; Lee, S. T. Chem. Vap. Deposition 2006, 12, 33–38. (10) Zhang, W. J.; Meng, X. M.; Chan, C. Y.; Chan, K. M.; Wu, Y.; Bello, I.; Lee, S. T. J. Phys. Chem. B 2005, 109, 16005–16010. (11) Lifshitz, Y.; Kohler, T.; Frauenheim, T.; Guzmann, I.; Hoffman, A.; Zhang, R. Q.; Zhou, X. T.; Lee, S. T. Science 2002, 297, 1531–1533. (12) Sofo, J. O.; Chaudhari, A. S.; Barber, G. D. Phys. Rev. B 2007, 75, 153401. (13) Elias, D. C.; Nair, R. R.; Mohiuddin, T. M. G.; Morozov, S. V. Science 2009, 323, 610–613. (14) Nair, R. R.; Ren, W.; Jalil, R.; Riaz, I.; Kravets, V. G. Small 2010, 6, 2877–2884. (15) Jeon, K. J.; Lee, Z.; Pollak, E.; Moreschini, L.; Bostwick, A.; Park, C. M. ACS Nano 2011, 5, 1042–1046. (16) Zboril, R.; Karlicky, F.; Bourlinos, A. B.; Steriotis, T. A.; Stubos, A. K.; Georgakilas, V.; Safarova, K.; Jancík, D.; Trapalis, C.; Otyepka, M. Small 2010, 6, 2885–2891. (17) Robinson, J. T.; Burgess, J. S.; Junkermeier, C. E.; Badescu, S. C.; Reinecke, T. L.; Perkins, F. K.; Zalalutdniov, M. K.; Baldwin, J. W.; Culbertson, J. C.; Sheehan, P. E.; Snow, E. S. Nano Lett. 2010, 10, 3001–3005. (18) Leenaerts, O.; Partoens, B.; Peeters, F. M. Phys. Rev. B 2009, 80, 245422. (19) Hornekær, L.; Rauls, E.; Xu, W.; Sljivancanin, Z.; Otero, R.; Stensgaard, I.; Lægsgaard, E. Phys. Rev. Lett. 2006, 97, 186102. (20) Balog, R.; Jørgensen, B.; Wells, J.; Lægsgaard, E.; Hofmann, P.; Besenbacher, F.; Hornekær, L. J. Am. Chem. Soc. 2009, 131, 8744–8745. (21) Jin, C.; Lin, F.; Suenaga, K.; Iijima, S. Phys. Rev. Lett. 2009, 102, 195505. (22) Lin, Y.; Williams, T. V.; Connell, J. W. J. Phys. Chem. Lett. 2010, 1, 277–283. (23) Zeng, H.; Zhi, C.; Zhang, Z.; Wei, X.; Wang, X.; Guo, W.; Bando, Y.; Golberg, D. Nano Lett. 2010, 10, 5049–5055. (24) Zhang, Z. H.; Guo, W. L. J. Am. Chem. Soc. 2009, 131, 6874–6879. (25) Xiang, H. J.; Yang, J.; Hou, J. G.; Zhu, Q. Appl. Phys. Lett. 2005, 87, 243113. (26) Zhou, Z.; Zhao, J.; Chen, Z.; Schleyer, P. R. J. Phys. Chem. B 2006, 110, 25678–25685. (27) Dean, C. R.; Young, A. F.; Meric, I.; Lee, C.; Wang, L.; Sorgenfrei, S.; Watanabe, K.; Taniguchi, T.; Kim, P.; Shepard, K. L.; Hone, J. Nat. Nanotechnol. 2010, 5, 722–726. 21683
dx.doi.org/10.1021/jp207175u |J. Phys. Chem. C 2011, 115, 21678–21684
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(28) Usachov, D.; Adamchuk, V. K.; Haberer, D.; Gruneis, A.; Sachdev, H.; Preobrajenski, A. B.; Laubschat, C.; Vyalikh, D. V. Phys. Rev. B 2010, 82, 075415. (29) Xue, J.; Yamagishi, J. S.; Bulmash, D.; Jacquod, P.; Deshpande, A.; Watanabe, K.; Taniguchi, T.; Herrero, P. J.; LeRoy, B. J. Nat. Mater. 2011, 10, 282. (30) Kresse, G.; Hafner, J. Phys. Rev. B 1994, 49, 14251–14269. (31) Kresse, G.; Furthmuller, J. Phys. Rev. B 1996, 54, 11169–11186. (32) Blochl, P. E. Phys. Rev. B 1994, 50, 17953–17979. (33) Mills, G.; Jnsson, H.; Schenter, G. K. Surf. Sci. 1995, 324, 305–337. (34) Henkelman, G.; Uberuaga, B. P.; Jnsson, H. J. Chem. Phys. 2000, 113, 9901–9904. (35) Dumitric~a, T.; Hua, M.; Yakobson, B. I. Phys. Rev. B 2004, 70, 241303. (36) Barone, V.; Hod, O.; Scuseria, G. E. Nano Lett. 2006, 6, 2748– 2754. (37) Badziag, P.; Verwoerd, W. S.; Ellis, W. P.; Greiner, N. R. Nature 1990, 343, 244–245. (38) Leenaerts, O.; Partoens, B.; Peeters, F. M. Phys. Rev. B 2009, 80, 245422. (39) Zhou, J.; Wang, Q.; Sun, Q.; Jena, P. Phys. Rev. B 2010, 81, 085442. (40) Barboza, A. P. M.; Guimaraes, M. H. D.; Massote, D. V. P.; Campos, L. C. Adv. Mater. 2011, 23, 3014–3017. (41) Zhang, Z.; Guo, W. Phys. Rev. B 2008, 77, 075403. (42) Samarakoon, D. K.; Wang, X. Q. ACS Nano 2010, 4, 4126–4130. (43) Khoo, K. H.; Mazzoni, M. S. C.; Louie, S. G. Phys. Rev. B 2004, 69, 201401R.
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