Versatile Electronic and Magnetic Properties of SnSe2 Nanostructures

Article pubs.acs.org/JPCC

Versatile Electronic and Magnetic Properties of SnSe2 Nanostructures Induced by the Strain Yucheng Huang,* Chongyi Ling, Hai Liu, Sufan Wang, and Baoyou Geng* Center for Nano Science and Technology, College of Chemistry and Material Science, The Key Laboratory of Functional Molecular Solids, Ministry of Education, Anhui Normal University, Wuhu, 241000, People’s Republic of China ABSTRACT: First-principles calculations were employed to explore the electronic and magnetic properties of a twodimensional (2D) SnSe2 monolayer sheet and its derived onedimensional (1D) nanoribbons and nanotubes. The results unveiled that the semiconductor−metal or metal−semiconductor transition can be realized by subtly controlling the strain for all these nanostructures. Surprisingly, without introduction of impurities and the absence of transition metal atoms, a −10% compressive strain can induce magnetic behaviors in SnSe2 armchair nanoribbons and the emerged magnetic moment increases rapidly and linearly with the increase of strain. The magnetism is found to be stemmed from the nonmetallic anionic Se atom at the ribbon edge. The tunable electronic and magnetic properties can be well understood through the analysis of partial charge density distribution and partial density of states. It was found that the direction of applied strain is a determined factor that can affect the energy shift of Se p orbital, leading to different composition of the states near the Fermi level. Finally, the stabilities of these SnSe2 nanostructures were evaluated for the possibility of experimental realizations. We believe that our results will provide useful information for their potential applications in electromechanical nanodevices, which will stimulate further experimental and theoretical investigations in this field.

1. INTRODUCTION Due to the quantum confinement effect, low-dimensional nanostructures have shown unrivalled chemical and physical properties which have been the focus of research for several decades. As one of the most representative examples, the longrange π-conjugation of a two-dimensional (2D) sheet of graphene makes it exhibit extraordinary thermal, mechanical, and electrical properties. Through caving the 2D graphene, its derived one-dimensional (1D) graphene nanoribbons and nanotubes have been realized. It was proved that these nanostructures possess versatile electronic and magnetic properties closely associated with the low-dimensional quantum confinement effect and peculiar edge effect, and are promising for many applications, such as in catalysis, energy storage, spintronics, the fabrication of novel electronic devices, etc.1−9 Because of its graphene-like structure, recently interest in the layered chalcogenide materials (LCMs) is poised to extend well beyond the territory of graphene and become a new hotspot of research.10−22 Due to excellent optical and electrical properties, LCMs can be used in transistors,13−15 catalysis,16 optoelectronic devices,17−19 energy storage,20−22 and so on. Different from the graphene of zero band gap, most of the 2D LCMs sheets display a semiconducting feature with a moderate band gap, leading them to be modulated more easily. The electronic and magnetic properties modulation of LCMs like MoS216,18,23−25 has been proved to be an effective means to widen their potential applications. These novel materials and © 2014 American Chemical Society

the modulation inspiration included bring us new breakthroughs to explore other important and fantastic LCMs. As a member of the LCM family, SnSe2 is a IV−VI group semiconductor with a hexagonal crystal structure of the type CdI2.26 The covalently bonded layer of Sn−Se−Sn held by van der Waals forces endows its potential applications in the field of lithium ion batteries, photovoltaic devices, solar cells,27 as well as phase change memory.28 Our previous study showed that SnSe2 always displays a monotonous semiconducting property without magnetism, irrespective of 2D monolayer or 1D nanoribbons, which greatly hinders its wider applications.29 Thus, it is necessary to investigate the modulation of its electronic and magnetic properties. Recently, biaxial or uniaxial strain was proved to be a powerful strategy to tune the electronic and magnetic properties of 2D or 1D nanostructures of LCMs. For example, by using first-principles calculations, Ma et al.11 showed that VX2 (X = S, Se) monolayers are ferromagnetic and the magnetic moments can increase rapidly with the increase of isotropic strain from −5% to 5%; from the same group, by applying the strain, the half-fluorinated BN and GaN sheets exhibit intriguing magnetic transitions between ferromagnetism and antiferromagnetism.30 Zhou et al.12 found that uniaxial tensile strain can induce ferromagnetic character in layered Received: February 6, 2014 Revised: April 8, 2014 Published: April 9, 2014 9251

dx.doi.org/10.1021/jp5013158 | J. Phys. Chem. C 2014, 118, 9251−9260

The Journal of Physical Chemistry C

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a single layer of SnSe2 sheet, producing AC and ZZ nanotubes, respectively. For these two kinds of nanotubes, the k points setups are the same as that of nanoribbons. The energy cutoff and k points settings are proved to be sufficient for achieving converged results. The effect of strain can be applied to these nanostructures by isotropic biaxial or uniaxial strain. Here, the strain is defined as ε = Δa/a0, where the a0 and Δa + a0 are the lattice constants of the unstrained and strained supercells, respectively. Thus, the positive/negative value of ε indicates the tensile/compressive strain.

NbX2 (X = S, Se), and the magnetic moment is strongly dependent on the strain; the investigation of Lu et al.24 revealed that strain can effectively steer the electronic and magnetic properties of MoS2 monolayer, bilayer, nanoribbons, and nanotubes. Besides these theoretical predictions, large31,32 or small33,34 strains have been exerted on nanosheets in laboratories, indicating that the controllable electronic and magnetic properties can be realized in practice by applying the different levels of strain. Motivated by these studies, in this paper the effect of strain on the electronic and magnetic properties of 2D SnSe2 monolayer and its derived 1D nanoribbons and nanotubes was systemically explored by means of density functional theory (DFT) calculations. The aim of this study is, on one hand, to investigate the effect of strain on the electronic and magnetic properties of SnSe2 and the resulting potential applications and, on the other hand, to explore the effect of quantum confinement with dimensionality decreasing from 3 (bulk), 2 (sheet), to 1 (ribbon, tube). Our results showed that, without strain, SnSe2 nanostructures behave as a semiconductor or metal without any magnetism. However, when subjected to different levels of strain, a semiconductor−metal or metal− semiconductor transition always can be realized. Interestingly, accompanied by the occurrence of the semiconductor−metal transition, SnSe2 armchair nanoribbon presents magnetism and the magnetic moment strongly depends on the strain. The surprised magnetism induced by the strain indicates that it can be applied in practical electromechanical nanodevices. We hope that our interesting results will provide some help to further experimental and theoretical work in connection with SnSe2, especially in directing the way to the potential applications of SnSe2 nanostructures.

3. RESULTS AND DISCUSSION 3.1. 2D SnSe2 Monolayer. The model of the SnSe2 monolayer was cleaved from the SnSe2 (001) surface, which contains a Se−Sn−Se triple layer. As shown in Figure 1a, the

2. COMPUTATIONAL DETAILS Our first-principles DFT computations were performed by using the projector-augmented plane wave (PAW)35 method to model the ion−electron interaction as implemented in the Vienna ab initio simulation package (VASP).36,37 Over the computations, the generalized gradient approximation (GGA) in the Perdew −Burke−Ernzerhof (PBE) form38,39 and a 350 eV cutoff for plane-wave basis set were adopted. 5s25p2 electron states of the Sn atom and 4s24p4 electron states of the Se atom were considered as valence. The convergence threshold was 10−5 eV for energy and 10−2 eV/Å for force. A vacuum space of at least 12 Å is adopted to avoid the interaction between two periodical units. Both spin-unpolarized and spin-polarized calculations were performed to determine the ground state for all the structures. Further calculations of both ferromagnetic and antiferromagnetic couplings were employed to confirm the type of magnetism. Unless otherwise stated, these nanostuctures all present a nonmagnetic ground state. For the geometric and electronic structural calculations of the 2D structure, a supercell consisting of 3 × 3 unit cells of SnSe2 monolayer was employed. The Brillouin zone is represented by the set of 4 × 4 × 1 k points for the geometry optimizations, and a total of 40 k points for electronic structure calculations. For 1D nanoribbons, two kinds of terminated edges were constructed, i.e., armchair (AC) and zigzag (ZZ), both containing two unit cells in the supercell. Brillouin zone sampling for geometry optimizations used a 6 MonkhostPack40 k point grid. On the basis of the equilibrium structures, 21 k points were then used to compute the electronic band structures. 1D SnSe2 nanotubes were constructed by rolling up

Figure 1. (a) The top and side view of SnSe2 monolayer. (b) Sn−Se bond length and (c) energy gaps as a function of strain; (d) the electronic structures of SnSe2 monolayer under various strain. The red dashed lines in part d and blue arrows in part a represent the Fermi level and strain direction, respectively.

plane of Sn atoms is sandwiched between two planes of Se atoms, and each Sn atom is centered at an octahedron constructed of six Se atoms. To investigate the strain effect on the geometry and band structure, the biaxial strains from ε = −11% to 15% were applied. Figure 1b presents the variation of Sn−Se bond length in the SnSe2 monolayer under different levels of strain. At equilibrium state without strain, the Sn−Se bond length is calculated to be 2.74 Å, which is slightly longer than the one in the SnSe2 bulk. From ε = −11% to 15%, the Sn−Se bond length increases monotonously (Figure 1b). The largest 6% increase of the bonding lengths was found at a 15% tensile strain, as compared to the 10% increase of bonding lengths in graphene and the BN layer for the 10% strain. Over the whole strain process, the structure of the SnSe2 monolayer remains integrated without any Sn−Se bond breaking, implying a relatively large range of elastic limits. The corresponding variation of energy gaps with the strain is illustrated in Figure 1c. A band gap of 0.79 eV was calculated at equilibrium state, which is somewhat wider than the one (0.63 eV) of the SnSe2 bulk counterpart.29,41 The energy gap reduces monotonously with the increase of compressive strain. Different from the 9252



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Figure 2. The partial density of states (PDOS) of SnSe2 monolayer under different strain. The insets show the partial densities corresponding to some special points; the isosurface value is set to be 0.001 e/Å3. The red dashed line represents the Fermi level.

pattern change of the Sn−Se bond length, with the enhancement of tensile strain, the energy gap increases first, and reaches the highest value of 0.91 eV at ε = 4%, and then reduces linearly. Interestingly, when strain reaches a certain degree, e.g., compressive/tensile strain exceeds 11%/14% (Figure 1c), the energy gap reduces to zero with some electronic states across the Fermi level (Ef). In other words, both contractive and stretching strain can trigger the semiconductor−metal transition in the SnSe2 monolayer. In Figure 1d, the band structures of SnSe2 under selected strains are shown as schematic representations. It can be seen that without strain, the monolayer is more likely to possess a direct-band gap semiconducting character, whose valence band maximum (VBM) and conduction band minimum (CBM) are located at AH/MH and ML points, respectively (Figure 1d, the energy difference between AH and MH points is only 0.01 eV). Recalling the indirect semiconducting character of SnSe2 bulk,29,41 our calculation shows that as the thickness of layered SnSe2 decreases from the bulk toward the monolayer limit, the transition from an indirect to a direct band gap happens, indicating the emergence of the photoluminescence and the potential application to optical devices. A similar conclusion was also reached in the MoS2 system.42 Under compressive strain, the change in the band gaps can be related to the energy increase of the valence band and decrease of the conduction band (Figure 1d). Under tensile strain, the energy of the valence band still increases; however, the energy of the conduction band at the Γ point dramatically decreases, leading to the closure of band gaps ultimately and the occurrence of semiconductor−metal transition. To gain insight into the semiconductor−metal transition induced by the strain, we plotted the partial density of states (PDOS) as well as the partial charge density distribution at some special points in Figure 2. As the energy gap at ε = 4% is the largest, here the state of ε = 4% is regarded as the base point for the sake of discussion. An obvious shift of the Se pz

orbital in the valence band (indicated by the red solid line) in PDOS was observed, that is, compressive strain makes it shift away with respect to the Ef while tensile strain drives it closer and even across the Ef (the moving direction is denoted by the red arrow, Figure 2). The shift direction of Se px, py orbitals refers to the opposite scenario (Figure 2, not marked). The movement of pz or px, py orbitals can be ascribed to the variation of the Sn−Se bond length. From ε = 15% to −11%, the Sn−Se bond length continuously decreases (Figure 1b), therefore, the overlapping of the in-plane xy atomic orbitals is very significant, driving them to the high energy level. This is evidenced by the fact that the Se px, py orbitals gradually get close to the Ef (Figure 2). In the meantime, the Se pz orbital is drawn far away from the Ef because the out-of-plane atomic orbital receives less repulsion. When compressive/tensile strain reaches a certain extent, px, py/pz orbitals can cross over the position of Ef, and thus the transition from semiconductor to metal occurs. The partial charge density distribution analysis can give a vivid picture for the reason for p orbital displacements. As seen from the inset in Figure 2, at ε = 4%, the ML and ΓL points are mainly contributed by the coupling between Sn atom s and Se atom px, py orbitals, while the KH and ΓH originate from Se pz and py orbitals, respectively. With the increase of compressive strain, the Se atom px, py orbitals will gradually substitute the pz orbital to contribute to the KH point, in agreement with the analysis from PDOS. On the other hand, at the ML point, the charge becomes denser, meaning that their contribution to the Ef would increase. Indeed, this is supported by the PDOS analysis as a small peak mostly composed of an Sn s orbital gets close to the Ef (Figure 2). On the contrary, with the increase of stretching strain, the Se atom pz orbital will substitute the px and py orbitals to contribute to both ΓL and ΓH points, leading to reduction and even closure of the energy gap. The transition from semiconductor to metal was also found in the MoS2 monolayer through both tensile and compressive strains.23 9253



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3.2. 1D SnSe2 Nanostructures. Because of the quantum confinement effect, 1D nanostructures have engendered intense scientific interest.5−10,43−48 From a nanotechnology perspective, 1D structures offer a range of potential applications that are different from those provided by their 2D and 3D counterparts.49 Thus, here both the 1D-SnSe2 nanoribbons and nanotubes are investigated, which will be discussed in Sections 3.2.1 and 3.2.2, respectively. 3.2.1. SnSe2 Nanoribbons. From the parental 2D-SnSe2 single-layer crystal, two kinds of nanoribbons named zigzag (ZNR) and armchair (ANR) can be made. We first address the electronic and magnetic properties of ZNR. Figure 3a shows

Figure 4. The PDOS of SnSe2 8-ZNR and partial charge density distributions of ΓH and ΓL points under different strain. The isosurface value is set to be 0.001 e/Å3 and the red dashed line represents the Fermi level. Figure 3. (a) The top and side view of SnSe2 8-ZNR, and the direction of uniaxial strain (blue arrows). (b) The band structures of SnSe2 8ZNR under various strain; (c) energy gaps as a function of strain. The red dashed line represents the Fermi level.

the other hand, the composition of the ΓL point changes little and the charges assembling at the inner site tend to transfer to edges. Thus, the charge localization to delocalization would account for the semiconductor−metal transition, which is consistent with the studies regarding SnSe nanoribbons50 and carbon nanotubes.51 On the contrary, when tensile strain was applied, the composition of the ΓL point changes dramatically while the ΓH point is almost unvaried. As seen from PDOS in Figure 4, the main contribution to the ΓL point transforms from Sn s and Se p orbitals to Sn s and Se px. Recalling the results on the monolayer, we can conclude that the composition of the states near the Fermi level strongly depends on the direction of the applied strain: for the SnSe2 monolayer, the direction of biaxial strain is along the xy plane, then the changes of energy gap originate from whether the electronic states near the Ef are contributed by Se atom px, py orbitals or the pz orbital; for SnSe2 8-ZNR, the direction of uniaxial strain is along the y-axis, and thus the shift in energy gap depends on which orbital contribution is dominant, Se atom p orbitals or py. Therefore, we have a reason to believe that one can artificially manipulate the electronic properties through controlling the direction of strain, and further through tuning different levels of strain, to realize the semiconductor−metal or metal−semiconductor transition. Next we turn our attention to the ANRs. Here 10-ANR was selected (Figure 5a) as the prototype. As shown in Figure 5b, the unstrained 10-ANR is an indirect semiconductor. Compared with ZNRs whose energy gap is closed at ε =

top and cross-section view of the structure of SnSe2 8-ZNR, with the uniaxial strain applied in the range of ε = −10% to 10% along its growth direction. The corresponding energy band structures are illustrated in Figure 3b. Obviously, 8-ZNR presents a direct-energy gap semiconducting character at the equilibrium state. Under tensile strain, the direct semiconducting character is still kept and the energy gap decreases with the strain increasing. Similarly, compression can also shift the ΓH point upward and the ΓL point downward, leading to a decrease in the energy gap. Furthermore, the energy gap can be closed when ε reaches −8%, indicating the semiconductor− metal transition can also be realized by applying compressive strain on SnSe2 ZNRs. In Figure 4, the PDOS and partial charge density distribution at ΓH and ΓL points were respectively plotted to understand the band gap engineering through controlling uniaxial strain. At the equilibrium state, the ΓH point is mainly contributed by the Se atom p orbital and the charges are almost completely localized on the edges, while the ΓL point originates from the coupling between the Se atom p and Sn atom s orbitals, and the charges assemble at the inner site. With the increase of compressive strain, the main contribution to the ΓH point changes to the Se atom py orbital, together with the charge delocalization (gradual transfer to the inner site, Figure 4). On 9254



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comparable width. Stain-induced magnetism in the absence of magnetic impurities is remarkable due to possible technological applications regarding both information storage and processing devices. The corresponding energy differences (ΔE) between spin-unpolarized and spin-polarized as a function of strain present the same tendency as shown in Figure 6b. The spinpolarized states are not observed when compressive strain is less than 10%. However, the ferromagnetic (FM) ground state is induced with compression more than 10%. The ΔE is only 0.11 meV at ε = −10%, indicating the presence of FM coupling. With the increase of compressive strain, the stability of the FM coupling is greatly enhanced. The ΔE reaches a considerable value of 23.58 meV at ε = −12%, meaning the FM state is relatively stable. In Figure 6c, the spatial spin distribution under −11% compressive strain is illustrated to understand the origin of the magnetism. It can be seen that the unpaired spin mainly concentrates on the edge Se atoms, contributed by the 4p electrons. Different from zigzag MoS2 nanoribbons10 whose magnetism is contributed by both edge Mo and S atoms, or the VX211 or NbX212 (XS, Se) monolayer whose magnetism is mainly contributed by the metal atoms (V, Nb), the magnetism entirely induced by the nonmetallic anion is extremely intriguing because this observation, to our best knowledge, has not been found thus far. It is worthy to note that during the strain process, the changes of Sn−Se bond lengths in ANR or ZNR are always less than 7% with respect to the one at the equilibrium state (Table 1). The structure breakage does not occur, indicating a large elastic range and a possibility to control well the electronic and magnetic properties. As we know, materials with magnetic properties which can be induced in a well-controlled way often provide advanced applications,11,52,53 such as practical electromechanical nanodevices, a switch for spin-polarized transport, etc. To have a deeper insight into the strain-induce magnetic properties in SnSe2 ANR, the spin-polarized partial density of states (PDOS) of Sn and edge Se atoms under different compressive strains is plotted in Figure 7. Here we only check Sn 5p and Se 4p orbitals since the PDOS analysis shows that these orbitals contribute to the majority of the electronic states under the Fermi level, forming the Sn−Se covalent bonds. As seen from Figure 7, the states near the Fermi level are almost totally contributed by Se 4p orbitals and the contribution of Sn 5p orbitals is insignificant. However, the most important effect for Sn 5p is located at the low-energy regions, ranging from −5.0 to −0.5 eV. Previous studies12,54 showed that the hybridized states in these regions can transform into a perturbation to the polarized atoms, resulting in the decrease of splitting effect. With the increase of compressive strain, the

Figure 5. (a) The top and side view of SnSe2 10-ANR, and the direction of uniaxial strain (blue arrows). (b) The band structures of SnSe2 10-ANR under various strain; (c) energy gaps as a function of strain. The red dashed line represents the Fermi level.

−8%, the semiconductor−metal transition does not occur in the range of ε = −8% to 10%, indicating that the electronic structure of ANRs is less sensitive to the strain. The energy gap always decreases no matter what tensile or compressive strain is applied. The effect of compressive strains is obviously stronger. As seen from Figure 5c, from ε = −2% to −8% the energy gap of 10-ANR reduces linearly and monotonously, which resembles the scenarios in the SnSe2 monolayer and 8-ZNR. Though all the ANRs are calculated to be indirect semiconductors, the energy band of VBM is rather flat, indicating they are easy to transform into direct ones. It is very interesting that the nonmagnetic semiconductor can be turned into ferromagnetic metal with further enhancement of compressive strain in SnSe2 10-ANR. As shown in Figure 6a, a −10% compressive strain can induce about 0.015 μB per unit SnSe2 molecule, and simultaneously the energy gap is closed. Surprisingly, the emerged unit magnetic moment per SnSe2 molecule (Mmol) increases monotonously, exhibiting a nearly linear relationship with the increase of compressive strain (Figure 6a). At ε = −12%, the Mmol is 0.081 μB (0.81 μB per unit cell), which is similar to that of MoS2 ZNR10 with the

Figure 6. Strain dependence of (a) the unit magnetic moment per SnSe2 molecule (Mmol), (b) the energy difference between spin-unpolarized and spin-polarized states (ΔE), and (c) spatial spin distribution (up−down) of 10-ANR under −11% compressive strain. 9255



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Table 1. Sn−Se Bond Lengths (in Å) of SnSe2 8-ZNR and 10-ANR with the Strain 8-ZNR strain (ε, %) −12 −11 −10 −8 −6 −4 −2 0 2 4 6 8 10

LMina

2.557 2.559 2.568 2.581 2.612 2.611 2.628 2.644 2.636 2.627 2.618

10-ANR LMaxb

LMin

LMax

2.801 2.782 2.770 2.763 2.788 2.788 2.815 2.842 2.866 2.915 2.973

2.502 2.501 2.500 2.498 2.500 2.503 2.506 2.511 2.517 2.524 2.532 2.542 2.555

2.884 2.873 2.855 2.840 2.839 2.839 2.838 2.839 2.825 2.820 2.870 2.932 3.010

a

LMin denotes the minimum bond length in SnSe2 nanoribbons. bThe maximum bond length. The bond length is not homogeneous in the ribbon. LMin always appears on the edge while the LMax locates on the domains between the inner part and the edge. Figure 8. Spin-polarized band structures of SnSe2 10-ANR under different compressive strain. The black and blue lines represent the spin-up and spin-down components, respectively. The red dashed line represents the Fermi level.

degeneracy is broken which makes ANR into a magnetic metal when ε is beyond −10%. 3.2.2. SnSe2 Nanotubes. To date, although the synthesis of SnSe2 nanotubes has not been reported, the formation of these closed caged nanoparticles is predicted to be favorable. It is known that the high energy of dangling bonds at the edge of the graphene planes makes graphite preferably form hollow closed structures.55,56 Due to graphene-like structures with anisotropic layer-type, the formation of closed caged nanoparticles is a general property for other inorganic layered structures. For example, as one of the chalcogenides, fullereneand nanotube-like SnS2 nanoparticles were synthesized by Yella et al.57 Previous to this, the electronic properties of SnS2 nanotubes have been revealed.58 Therefore, our theoretical attempts toward the SnSe2 nanotubes can give direct information to the experimental realization and future potential application. Here, we take the cases of the (10, 10) SnSe2 armchair nanotube (ANT) and the (12, 0) zigzag nanotube (ZNT) to investigate the electronic and magnetic properties as well as the effect of strain. By calculating the strain energy, we found that with the comparable diameter, SnSe2 nanotube is less stable than carbon nanotube (CNT), but more stable than the MoS2 one (Figure 9). Since the latter has been successfully synthesized in laboratory, this result further supports the prediction of feasible preparation of SnSe2 nanotubes. Furthermore, the diameter of the chosen SnSe2 nanotubes locates on the relatively flat curve with converged strain energy in Figure 9, demonstrating the representativeness of our model. Figure 10a shows the geometry of (10, 10) SnSe2 ANT. At the equilibrium state, it is an indirect-band gap semiconductor in which the VBM and CBM are located at ΓH and ZL points, respectively (Figure 10c). Both stretching and compression can effectively tune the energy gap. In the course of strain, the diameter of ANT displays a decrease trend from 19.17 to 18.19

Figure 7. Spin-polarized PDOS of Sn and edge Se atoms under different compressive strain. The green and blue regions represent the spin-up and spin-down components, respectively. The red dashed lines and arrows represent the Fermi level and the spin splitting near the Fermi level, respectively.

hybridized states at these energy regions are somewhat decreased, indicating that the perturbations are weakened and the splitting effects are enhanced. Thus, the magnetic moment of SnSe2 10-ANR increases with the compressive strain increasing. The spin polarization with strain is in accordance with the variation of Sn−Se bond length according to the viewpoint of Ma et al.11 It was believed that the elongation of the bond length results in reduction in the covalent bonding interaction and the enhancement in the ionic bonding interaction, thus giving rise to the unpaired electrons accumulating and the interesting variation in the magnetic moment with strain. As a further support, Figure 8 shows the variation of spinresolved band structures with the compressive strain. At ε = −9.5%, the SnSe2 10-ANR transforms into nonmagnetic metal with symmetry spin-up and spin-down bands; however, with the continuous enhancement of compressive strain, the spin 9256



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Figure 9. Strain energy (in eV) as a function of tube diameter (in Å) for SnSe2, MoS2,59 and CNTs.60 Strain energy is defined as the energy difference between the tube and the planar configuration for several sizes of tube.

Figure 11. Top and cross-section view of (12, 0) SnSe2 ZNT (a) without and (b) with tensile strain. Blue arrows denote the strain direction. (c) The diameter (in Å) and energy gap (in eV) variations as a function of strain, and (d) the band structures under different levels of strain. The red dashed line represents the Fermi level.

stems from the structure deformation. As seen from Figure 11b, the diameter reduces sharply from 12.53 to 11.74 Å when the tensile strain changes from 2% to 4% (only 2% rangeability), such that the distances between every two Se atoms reduce greatly, and simultaneously the neighboring Sn−Se bond lengths are enlarged. Thus, some Sn−Se bonds are broken with a concomitant formation of some new Se−Se bonds. With the strain continuously increasing, the diameter further decreases and the corresponding energy gap further opens (Figure 11c,d). The electronic structure is determined by the structure whether it is a perfect or deformed one, which is in agreement with the case of SnSe nanoribbons.50 Above, we demonstrate the semiconductor−metal or metal− semiconductor transition can be respectively realized in ANT by applying compressive strain or in ZNT through the tensile strain. Still, the analysis from PDOS can give direct information for these transitions. Similar to the SnSe2 monolayer and nanoribbons, the different composition of the electronic states near the Ef depends on the direction of strain. As seen from Figure 12a, by applying compressive strain along the z direction in ANT, the energy level of Se pz is drawn near the Ef. On the other hand, because of Se−Se bond formation on ZNT with the enhancement of tensile strain, the PDOS shape of the Se pz orbital turns smooth with more peaks emerging. Obviously, the metal−semiconductor transition derives from the downward shift of the Se px, py orbitals with respect to the Ef (Figure 12b). This is understandable because the z direction tensile strain makes px, py orbitals receive less repulsion, driving these orbitals to the lower energy domains. 3.3. Stabilities of SnSe2 Monolayer, Nanoribbons and Nanotubes under Different Strain. For nanostructures, the stabilities are vital to evaluate the possibility of experimental realizations. Of course, their stabilities also rely heavily on the different levels of strain. Therefore, it is necessary to calculate the formation energies of these nanostructures with and without the presence of strain. To estimate their stabilities,

Figure 10. (a) Top and cross-section view of (10, 10) SnSe2 ANT, blue arrows denote the strain direction. (b) the diameter (in Å) and energy gap (in eV) variations as a function of strain; (c) the band structures of (10, 10) SnSe2 ANT under different levels of strain. The red dashed line represents the Fermi level.

Å as ε changes from −10% to 10% (Figure 10b). The change of diameter does not exceed 6% in the 20% strain range and the structures of ANT remain intact without any bond breaking, once again demonstrating the large elastic range of SnSe2 nanostructures. With the increase of compressive strain, the energy of the ΓH/ZL point increases/decreases, leading to the energy gap monotonous decrease. When compressive strain ε reaches 6%, the (10, 10) SnSe2 ANT becomes metallic with the closure of the energy gap. At variance, the energy gap change is not monotonous with the tensile strain. As shown in Figure 10b, the largest energy gap appears at ε = 2%. When strain exceeds 2%, the energy gap always decreases. Until ε = 10%, the ANT still presents the semiconducting character, indicating stretching cannot induce the semiconductor−metal transition, which is in accordance with the case of NRs. Figure 11a displays the structure of (12, 0) SnSe2 ZNT. Different from other nanostructures presenting the semiconducting character at the equilibrium state, band structure calculation indicates it is a metal. As seen in Figure 11d, the compressive strain cannot alter its metallic property. However, the energy gap can be opened when tensile strain is larger than 4% (Figure 11c,d). The dramatic electronic structure variation 9257



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Figure 12. The PDOS of (10, 10) SnSe2 ANT under compressive strain (a) and (12, 0) ZNT under tensile strain (b). The red dashed line represents the Fermi level.

the binding energy per SnSe2 molecule (Eb) is defined as Eb = (nESn + 2nESe − ESnnSe2n)/n, where ESn, ESe, and ESnnSe2n are the energies of Sn, Se, and SnnSe2n, respectively, and n is the number of SnSe2 molecules. A higher binding energy indicates a more stable structure. As shown in Figure 13, ε is in perfect

semiconductor−metal or metal−semiconductor transition always can be realized on these nanostuctures by applying different levels of strain. The strain direction strongly affects the energy shift of Se p orbitals, which is responsible for the transition. Furthermore, our results clearly demonstrate that, when subjected to a certain compressive strain, SnSe2 10-ANR is able to possess magnetic and possibly ferromagnetic behavior without the introduction of metal impurities. Therefore, the results open up the possibility of having ferromagnetism at room temperature in SnSe2 ANRs. The tunable electronic and magnetic properties of SnSe2 nanostructures are extremely significant for potential applications to nanoelectronics, nanoelectromechanical devices, and spintronics. We believe that our theoretical work will motivate further experimental and theoretical investigations in this field.

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected].

Figure 13. The variations of binding energy of SnSe2 monolayer, nanoribbons, and nanotubes with the strain.

Notes

The authors declare no competing financial interest.

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quadratic function with the Eb for all the nanostructures with the exception of ZNT under the tensile strain, whose structure experiences obvious deformation. These nanostructures are the most stable at the equilibrium state, and the strain always decreases their stabilities. Because of biaxial strain, the decay of the stabilities of the monolayer is more pronounced than for 1D-nanostructures with uniaxial strain. Basically, the stability order of these 1D-nanostructures is as follows: 8-ZNR > 10ANR > (10, 10) ANT > (12, 0) ZNT, indicating the formation of NTs is energetically less favorable than NRs. Generally, all these nanostructures display considerable binding energies with the strain, indicating a feasible experimental realization.

ACKNOWLEDGMENTS This work was supported by National Younger Natural Science Foundation of China No. 21203001, Natural Science Foundation of Anhui Province No. 1208085QB37, and Doctoral Scientific Research Funding of Anhui Normal University.

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REFERENCES

(1) Novoselov, K. S.; Geim, A. K.; Morozov, S.; Jiang, D.; Zhang, Y.; Dubonos, S.; Grigorieva, I.; Firsov, A. Electric Field Effect in Atomically Thin Carbon Films. Science 2004, 306, 666−669. (2) Schwierz, F. Graphene Transistors. Nat. Nanotechnol. 2010, 5, 487−496. (3) Bonaccorso, F.; Sun, Z.; Hasan, T.; Ferrari, A. Graphene Photonics and Optoelectronics. Nat. Photonics 2010, 4, 611−622.

4. CONCLUSION In summary, systematic first-principles calculations were carried out to investigate the effect of strain on the SnSe2 monolayer, nanoribbons, and nanotubes. Our results show that the 9258



Article

(26) Wyckoff, R. W. G. Crystal Structures; Interscience Publishers: New York, NY,1965. (27) Choi, J.; Jin, J.; Jung, I. G.; Kim, J. M.; Kim, H. J.; Son, S. U. SnSe2 Nanoplate-Graphene Composites as Anode Materials for Lithium Ion Batteries. Chem. Commun. 2011, 47, 5241−5243. (28) Wang, R. Y.; Caldwell, M. A.; Jeyasingh, R. G. D.; Aloni, S.; Shelby, R. M.; Wong, H.-S. P.; Milliron, D. J. Electronic and Optical Switching of Solution-Phase Deposited SnSe2 Phase Change Memory Material. J. Appl. Phys. 2011, 109 (113506), 1−6. (29) Huang, Y.; Ling, C.; Fang, Z.; Wang, S. Width- and EdgeDependent Magnetic Properties, Electronic Structures, and Stability of SnSe2 Nanoribbons. Phys. E 2014, 59, 102−106. (30) Ma, Y.; Dai, Y.; Guo, M.; Niu, C.; Yu, L.; Huang, B. StrainInduced Magnetic Transitions in Half-Fluorinated Single Layers of BN, GaN and Graphene. Nanoscale 2011, 3, 2301−2306. (31) Kim, K. S.; Zhao, Y.; Jang, H.; Lee, S. Y.; Kim, J. M.; Kim, K. S.; Ahn, J.-H.; Kim, P.; Choi, J.-Y.; Hong, B. H. Large-Scale Pattern Growth of Graphene Films for Stretchable Transparent Electrodes. Nature 2009, 457, 706−710. (32) Zhang, Y.; Luo, C.; Li, W.; Pan, C. Strain Induced Chemical Potential Difference between Monolayer Graphene Sheets. Nanoscale 2013, 5, 2616−2619. (33) Ni, Z. H.; Yu, T.; Lu, Y. H.; Wang, Y. Y.; Feng, Y. P.; Shen, Z. X. Uniaxial Strain on Graphene: Raman Spectroscopy Study and BandGap Opening. ACS Nano 2008, 2, 2301−2305. (34) Mohiuddin, T.; Lombardo, A.; Nair, R.; Bonetti, A.; Savini, G.; Jalil, R.; Bonini, N.; Basko, D.; Galiotis, C.; Marzari, N. Uniaxial Strain in Graphene by Raman Spectroscopy: G Peak Splitting, Grüneisen Parameters, and Sample Orientation. Phys. Rev. B 2009, 79 (205433), 1−8. (35) Blöchl, P. E. Projector Augmented-Wave Method. Phys. Rev. B 1994, 50, 17953−17979. (36) Kresse, G.; Furthmüller, J. Efficient Iterative Schemes for Ab Initio Total-Energy Calculations Using a Plane-Wave Basis Set. Phys. Rev. B 1996, 54, 11169−11186. (37) Kresse, G.; Joubert, D. From Ultrasoft Pseudopotentials to the Projector Augmented-Wave Method. Phys. Rev. B 1999, 59, 1758− 1775. (38) Perdew, J. P.; Chevary, J.; Vosko, S.; Jackson, K. A.; Pederson, M. R.; Singh, D.; Fiolhais, C. Atoms, Molecules, Solids, and Surfaces: Applications of the Generalized Gradient Approximation for Exchange and Correlation. Phys. Rev. B 1992, 46, 6671−6687. (39) Perdew, J. P.; Wang, Y. Accurate and Simple Analytic Representation of the Electron-Gas Correlation Energy. Phys. Rev. B 1992, 45, 13244−13249. (40) Monkhorst, H. J.; Pack, J. D. Special Points for Brillouin-Zone Integrations. Phys. Rev. B 1976, 13, 5188−5192. (41) He, X.; Shen, H. Ab Initio Calculations of Band Structure and Thermophysical Properties for SnS2 and SnSe2. Phys. B 2012, 407, 1146−1152. (42) Cheiwchanchamnangij, T.; Lambrecht, W. R. Quasiparticle Band Structure Calculation of Monolayer, Bilayer, and Bulk MoS2. Phys. Rev. B 2012, 85 (205302), 1−4. (43) Wu, M.; Wu, X.; Pei, Y.; Zeng, X. C. Inorganic Nanoribbons with Unpassivated Zigzag Edges: Half Metallicity and Edge Reconstruction. Nano Res. 2011, 4, 233−239. (44) Yu, S.-S.; Zheng, W.-T. Effect of N/B Doping on the Electronic and Field Emission Properties for Carbon Nanotubes, Carbon Nanocones, and Graphene Nanoribbons. Nanoscale 2010, 2, 1069− 1082. (45) Chen, W.; Li, Y.; Yu, G.; Li, C.-Z.; Zhang, S. B.; Zhou, Z.; Chen, Z. Hydrogenation: A Simple Approach To Realize SemiconductorHalf-Metal-Metal Transition in Boron Nitride Nanoribbons. J. Am. Chem. Soc. 2010, 132, 1699−1705. (46) Rodríguez-Manzo, J. A.; Tolvanen, A.; Krasheninnikov, A. V.; Nordlund, K.; Demortière, A.; Banhart, F. Defect-Induced Junctions Between Single- or Double-Wall Carbon Nanotubes and Metal Crystals. Nanoscale 2010, 2, 901−905.

(4) Zhu, Y.; Murali, S.; Cai, W.; Li, X.; Suk, J. W.; Potts, J. R.; Ruoff, R. S. Graphene and Graphene Oxide: Synthesis, Properties, and Applications. Adv. Mater. 2010, 22, 3906−3924. (5) Son, Y.-W.; Cohen, M. L.; Louie, S. G. Half-Metallic Graphene Nanoribbons. Nature 2006, 444, 347−349. (6) Yang, L.; Park, C.-H.; Son, Y.-W.; Cohen, M. L.; Louie, S. G. Quasiparticle Energies and Band Gaps in Graphene Nanoribbons. Phys. Rev. Lett. 2007, 99 (186801), 1−4. (7) Yamashiro, A.; Shimoi, Y.; Harigaya, K.; Wakabayashi, K. Spinand Charge-Polarized States in Nanographene Ribbons with Zigzag Edges. Phys. Rev. B 2003, 68 (193410), 1−4. (8) Tans, S. J.; Verschueren, A. R.; Dekker, C. Room-Temperature Transistor Based on A Single Carbon Nanotube. Nature 1998, 393, 49−52. (9) Kong, J.; Franklin, N. R.; Zhou, C.; Chapline, M. G.; Peng, S.; Cho, K.; Dai, H. Nanotube Molecular Wires as Chemical Sensors. Science 2000, 287, 622−625. (10) Li, Y.; Zhou, Z.; Zhang, S.; Chen, Z. MoS2 Nanoribbons: High Stability and Unusual Electronic and Magnetic Properties. J. Am. Chem. Soc. 2008, 130, 16739−16744. (11) Ma, Y.; Dai, Y.; Guo, M.; Niu, C.; Zhu, Y.; Huang, B. Evidence of the Existence of Magnetism in Pristine VX2 Monolayers (X = S, Se) and Their Strain-Induced Tunable Magnetic Properties. ACS Nano 2012, 6, 1695−1701. (12) Zhou, Y.; Wang, Z.; Yang, P.; Zu, X.; Yang, L.; Sun, X.; Gao, F. Tensile Strain Switched Ferromagnetism in Layered NbS2 and NbSe2. ACS Nano 2012, 6, 9727−9736. (13) Osada, M.; Sasaki, T. Two-Dimensional Dielectric Nanosheets: Novel Nanoelectronics from Nanocrystal Building Blocks. Adv. Mater. 2012, 24, 210−228. (14) Radisavljevic, B.; Radenovic, A.; Brivio, J.; Giacometti, V.; Kis, A. Single-Layer MoS2 Transistors. Nat. Nanotechnol. 2011, 6, 147−150. (15) Pu, J.; Yomogida, Y.; Liu, K.-K.; Li, L.-J.; Iwasa, Y.; Takenobu, T. Highly Flexible MoS2 Thin-Film Transistors with Ion Gel Dielectrics. Nano Lett. 2012, 12, 4013−4017. (16) Hu, K. H.; Hu, X. G.; Xu, Y. F.; Pan, X. Z. The Effect of Morphology and Size on the Photocatalytic Properties of MoS2. React. Kinet., Mech. Catal. 2010, 100, 153−163. (17) Splendiani, A.; Sun, L.; Zhang, Y.; Li, T.; Kim, J.; Chim, C.-Y.; Galli, G.; Wang, F. Emerging Photoluminescence in Monolayer MoS2. Nano Lett. 2010, 10, 1271−1275. (18) Lee, H. S.; Min, S.-W.; Chang, Y.-G.; Park, M. K.; Nam, T.; Kim, H.; Kim, J. H.; Ryu, S.; Im, S. MoS2 Nanosheet Phototransistors with Thickness-Modulated Optical Energy Gap. Nano Lett. 2012, 12, 3695−3700. (19) Thomalla, M.; Tributsch, H. Photosensitization of Nanostructured TiO2 with WS2 Quantum Sheets. J. Phys. Chem. B 2006, 110, 12167−12171. (20) Seayad, A. M.; Antonelli, D. M. Recent Advances in Hydrogen Storage in Metal-Containing Inorganic Nanostructures and Related Materials. Adv. Mater. 2004, 16, 765−777. (21) Liu, J.; Liu, X. W. Two-Dimensional Nanoarchitectures for Lithium Storage. Adv. Mater. 2012, 24, 4097−4111. (22) Chang, K.; Chen, W. In Situ Synthesis of MoS2/Graphene Nanosheet Composites with Extraordinarily High Electrochemical Performance for Lithium Ion Batteries. Chem. Commun. 2011, 47, 4252−4254. (23) Scalise, E.; Houssa, M.; Pourtois, G.; Afanas’ev, V.; Stesmans, A. Strain-Induced Semiconductor to Metal Transition in the TwoDimensional Honeycomb Structure of MoS2. Nano Res. 2012, 5, 43− 48. (24) Lu, P.; Wu, X.; Guo, W.; Zeng, X. C. Strain-Dependent Electronic and Magnetic Properties of MoS2 Monolayer, Bilayer, Nanoribbons and Nanotubes. Phys. Chem. Chem. Phys. 2012, 14, 13035−13040. (25) Ma, Y.; Dai, Y.; Guo, M.; Niu, C.; Lu, J.; Huang, B. Electronic and Magnetic Properties of Perfect, Vacancy-Doped, and Nonmetal Adsorbed MoSe2, MoTe2 and WS2 Monolayers. Phys. Chem. Chem. Phys. 2011, 13, 15546−15553. 9259



Article

(47) Woods, L.; Le, N. B. Zigzag Graphene Nanoribbons with Curved Edges. RSC Adv. 2013, 3, 10014−10018. (48) Son, Y.-W.; Cohen, M. L.; Louie, S. G. Energy Gaps in Graphene Nanoribbons. Phys. Rev. Lett. 2006, 97 (216803), 1−4. (49) Nalwa, H. Handbook of Nanostructured Materials and Nanotechnology; Academic Press: New York, NY, 2000. (50) Huang, Y.; Ling, C.; Liu, H.; Wang, S. Edge-, Width- and StrainDependent Semiconductor-Metal Transition in SnSe Nanoribbons. RSC Adv. 2014, 4, 6933−6938. (51) Liu, L.; Gao, J.; Guo, X.; Zhao, J. Electromechanical Properties of Zigzag-Shaped Carbon Nanotubes. Phys. Chem. Chem. Phys. 2013, 15, 17134−17141. (52) Chiba, D.; Yamanouchi, M.; Matsukura, F.; Ohno, H. Electrical Manipulation of Magnetization Reversal in a Ferromagnetic Semiconductor. Science 2003, 301, 943−945. (53) Tanaka, H.; Zhang, J.; Kawai, T. Giant Electric Field Modulation of Double Exchange Ferromagnetism at Room Temperature in the Perovskite Manganite/Titanate p-n Junction. Phys. Rev. Lett. 2001, 88 (027204), 1−4. (54) Santos, E. J.; Ayuela, A.; Fagan, S.; Mendes Filho, J.; Azevedo, D.; Souza Filho, A.; Sánchez-Portal, D. Switching on Magnetism in NiDoped Graphene: Density Functional Calculations. Phys. Rev. B 2008, 78 (195420), 1−5. (55) Kroto, H. W.; Heath, J. R.; O’Brien, S. C.; Curl, R. F.; Smalley, R. E. C60: Buckminsterfullerene. Nature 1985, 318, 162−163. (56) Ichihashi, T.; Ando, Y. Pentagons, Heptagons and Negative Curvature in Graphite Microtubule Growth. Nature 1992, 356, 776− 778. (57) Yella, A.; Mugnaioli, E.; Therese, H. A.; Panthöfer, M.; Kolb, U.; Tremel, W. Synthesis of Fullerene-and Nanotube-Like SnS2 Nanoparticles and Sn/S/Carbon Nanocomposites. Chem. Mater. 2009, 21, 2474−2481. (58) Chang, H.; In, E.; Kong, K.-j.; Lee, J.-O.; Choi, Y.; Ryu, B.-H. First-Principles Studies of SnS2 Nanotubes: A Potential Semiconductor Nanowire. J. Phys. Chem. B 2005, 109, 30−32. (59) Seifert, G.; Terrones, H.; Terrones, M.; Jungnickel, G.; Frauenheim, T. Structure and Electronic Properties of MoS2 Nanotubes. Phys. Rev. Lett. 2000, 85, 146−149. (60) Hernandez, E.; Goze, C.; Bernier, P.; Rubio, A. Elastic Properties of C and BxCyNz Composite Nanotubes. Phys. Rev. Lett. 1998, 80, 4502−4505.

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Versatile Electronic and Magnetic Properties of SnSe2 Nanostructures

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