Theoretical Prediction of Janus MoSSe as a Potential Anode Material

Article pubs.acs.org/JPCC

Cite This: J. Phys. Chem. C 2018, 122, 23899−23909

Theoretical Prediction of Janus MoSSe as a Potential Anode Material for Lithium-Ion Batteries Chanjuan Shang, Xueling Lei,* Binpeng Hou, Musheng Wu, Bo Xu, Gang Liu, and Chuying Ouyang Department of Physics, Laboratory of Computational Materials Physics, Jiangxi Normal University, Nanchang, Jiangxi 330022, China

J. Phys. Chem. C 2018.122:23899-23909. Downloaded from pubs.acs.org by UNIV OF SUNDERLAND on 10/25/18. For personal use only.

S Supporting Information *

ABSTRACT: Inspired by the synthesis of Janus MoSSe and its beneficial properties, we here report for the first time the adsorption and diffusion of Li-ion on the single-layer MoSSe (SLM) and the double-layer MoSSe (DLM) using first-principle computations. The results show that much more Li-ions can be stored by the SLM and DLM due to their intrinsic dipole moment and the charge redistribution. With a suitable open circuit voltage range vs Li+/Li, the ideal theoretical capacities for the SLM and DLM are 776.5 and 452.9 mAh/g, respectively. Furthermore, the calculated density of states of the lithiated SLM and DLM indicates that they have good electrical conduction, and the smaller Li-ion/Li-vacancy migration barrier ensures fast Li-ion diffusion. Our results suggest that the SLM and DLM can be utilized as a potential anode material for highperformance Li-ion batteries.

graphene (the reversible capacity up to 1290 mAh/g).20 Compared to MoS2 monolayer, Janus MoSSe with an intrinsic dipole may also be an alternative candidate as an anode material for LIBs. In particular, the role of dipole moments on the adsorption and diffusion of lithium ions is unclear. Hence, we explore the potential application of single-layer MoSSe (SLM) as an anode material for LIBs. In addition, the idea of forming van der Waals (vdW) heterostructures by integrating various 2D materials breaks the limitation of the restricted properties of single material systems,21,22 which sheds light on exploring new lithium storage materials. Some vdW heterostructures have been systematically investigated as anode materials for LIBs. For instance, it is reported that BlueP/MS2 vdW heterostructures are ideal candidates used as a promising electrode for highrecycling-rate LIBs with good structural stability, enhanced electrical conductivity, and high capacity.23 Similarly, Fan et al. proposed that the rate performance of the graphene/BlueP heterostructure is better than that of monolayer graphene, and the external electric field can significantly decrease the diffusion energy barrier of Li-ion.24 Also, 2D double-layer BlackP and BlueP were predicted to be good electrodes for high-capacity LIBs.25 Hence, motivated by above investigations, the performances of the MoSSe/MoSSe heterostructure, namely, the double-layer MoSSe (DLM), as an anode material of LIBs have also been systematically investigated in the present work.

1. INTRODUCTION Recently, the Janus MoSSe monolayer has been successfully synthesized by breaking the out-of-plane structural symmetry of the MoS21 and CVD growth,2 respectively. Later, great efforts have been devoted to the studies of various properties of single-layer MoSSe3 or multilayer MoSSe,4 including electronic structure,5 photocatalysts,6 magnetism,7,8 phonotransport,9 and so on.10 The Janus MoSSe has an intrinsic dipole perpendicular to the plane and the resulting internal electric field, the electrostatic potential difference between the upper and the lower surface is 0.78 eV.4 These unique properties of MoSSe engender a versatility that has enabled its use in a wide range of scientific fields, such as optoelectronics, etc. In recent years, the first requirement for lithium-ion batteries (LIBs) is to develop novel electrode materials to meet the increasing demands for energy density.11 To obtain higher capacity, anode materials are widely investigated.12 As we know, layered transition-metal disulfides MS2 (M = Mo, W, Nb, and Ta) can be used as anode materials because they can react with guest atoms to yield intercalation compounds.13,14 For example, layered MoS2 nanoplates or nanosheets are subject of great scientific attention for LIBs or sodium-ion batteries (SIBs).15,16 It is reported that layered MoS2 has a high theoretical specific capacity of 669 mAh/g based on a four-electron transfer per formula unit.17 However, their low intrinsic electrical conductivity leads to sluggish dynamics, poor rate behavior, and rapid capacity fading. Two typical strategies have been introduced to solve these issues. One is to synthesize nanostructured MoS2 with few-layer materials18 and the other is the hybridization of MoS2 and C19 or MoS2 and © 2018 American Chemical Society

Received: August 2, 2018 Revised: September 23, 2018 Published: October 4, 2018 23899

DOI: 10.1021/acs.jpcc.8b07478 J. Phys. Chem. C 2018, 122, 23899−23909

Article

The Journal of Physical Chemistry C In this work, for the first time, we investigate Li-ion storage behaviors in the SLM and DLM and their performance as an anode material of LIBs. With first-principles calculations, both the SLM and DLM are confirmed to be suitable as anode materials for LIBs. Moreover, the specific capacity of the SLM is higher than that of the DLM, but the diffusions of Li-ion and Li-vacancy are independent of the SLM and DLM.

OCV ≈

(x 2 − x1)e

(2)

where ELix1host and ELix2host are the total energy of Lix1host and Lix2host, respectively. μLi is the chemical potential of the intercalating species. According to the above formula, a positive voltage indicates energetically favorable intercalation.

3. RESULTS AND DISCUSSION 3.1. Li Adsorption on the SLM and DLM. The host structures of lithium storage, namely, the SLM and DLM with 3 × 3 × 1 supercell are shown in Figure 1. Considering the

2. COMPUTATIONAL METHODS All the calculations are performed by using the Vienna ab initio simulation package.26,27 The core ion and valence electron interaction is described by the projector-augmented-wave method28,29 and the generalized gradient approximation (GGA)30 with Perdew−Burke−Ernzerhof (PBE) functional is used to calculate the electron exchange−correlation interactions. The valence electron configurations for Li, S, Se, and Mo atom are 1s12s12p1, 3s23p4, 4s24p4, and 4p65s14d5, respectively. A plane-wave cutoff energy is set to be 400 eV for the plane-wave expansion of the wave function (convergence tests see Figure S1 of the Supporting Information). The lattice parameters and the atomic positions are fully relaxed until the force on each atom is less than 0.02 eV/Å. No less than 15 Å vacuum is used to avoid the interaction between the periodically repeated structures. For the structure relaxation and the electronic calculations, 9 × 9 × 1 for primitive cell and 3 × 3 × 1 for supercell Monkhorst−Pack k-point grids in the Brillouin zone are employed.31 Spin nonpolarization has been considered for all calculations. The charge distribution and transfer are studied quantitatively by Bader charge analysis.32 The Li-ion migration pathway is optimized with the climbing image nudged elastic band method.33,34 To better describe the interactions between the adsorbed Li-ions and the host materials and the interactions between the layers of MoSSe, van der Waals (vdW) correction is included in the present calculations. Specifically, the DFT-D2 method of Grimme35 has been employed to treat the vdW interactions in our systems (detailed test refer to Table S1 of the Supporting Information). Moreover, we have also performed the test calculations for PBE and LDA exchange−correlation functionals, respectively. The test results are summarized in Table S2 of the Supporting Information. From Table S2, we know that the adsorption energy obtained from LDA is larger than that obtained from PBE, but the relative stability of configurations does not change. The adsorption energy is important to understand the Li adsorption strength on a host material. Thus, the average adsorption energy is defined by Ead = (E host + nLi − E host − nE Li)/n

E Lix1host − E Lix2host + (x 2 − x1)μLi

Figure 1. Top view and side view of the MoSSe. (a) Single-layer MoSSe (SLM) with 3 × 3 × 1 supercell. (b) Double-layer MoSSe (DLM) with 3 × 3 × 1 supercell.

difference in atoms on each side of the Janus MoSSe, the bilayer configuration of MoSSe has three types of combination modes, that is, the SMoSe/SMoSe, SMoSe/SeMoS, and SeMoS/SMoSe. Our calculations demonstrate that the SMoSe/SMoSe bilayer has an intrinsic electric field caused by its dipole moment, which is different from conventional vdW heterostructures. Therefore, we take the SMoSe/SMoSe as an example to illustrate the adsorption behavior of Li on the DLM. Table 1 lists the lattice constants, bond lengths of Mo− Table 1. Lattice Parameters, Bond Distances of Mo−S and Mo−Se, Dipole Moments, and Band Gaps at PBE Theoretical Level for the SLM and DLM, Respectively bond length (Å)

(1) SLM DLM

where Ehost+nLi and Ehost are the total energy of the optimized ground state of the host (SLM or DLM) with and without n Li-ions adsorption, respectively. ELi is the total energy of an isolated Li atom. According to the definition shown in formula 1, the negative adsorption energy indicates that the Li-ion is tightly bound to the host. In addition, to further understand the performance of MoSSe-based LIBs, the open-circuit-voltage (OCV) has been estimated. In theory, the average OCV can be obtained directly by calculating the energy difference before and after Li intercalation, which is evaluated as follows

a = b (Å)

Mo−S

Mo−Se

μ (Debye)

Eg (eV)

3.26 3.26

2.42 2.42

2.54 2.54

0.18 0.33

1.55 0.81

S and Mo−Se, dipole moments, and band gaps of the SLM and DLM, respectively. The lattice constants are 3.26 Å for primitive cells of the SLM and DLM and the optimized Mo−S and Mo−Se bond lengths are 2.42 and 2.54 Å, respectively, which are in good agreement with previous results.6,36,37 The dipole moment of the DLM is almost 2 times that of the SLM, whereas the band gap of the DLM is about half that of the SLM. It should be noted that the relationship between the 23900


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Figure 2. (a, b) Li-adsorption sites on the SLM. (c−j) Selected top view and side view of the optimal structures of LixMoSSe.

dipole moment and the band gap was discussed in our other paper.38 To check Li-ion adsorption ability, we investigated the adsorption behavior of Li-ion on the SLM and DLM, respectively. Figure 2a,b shows different adsorption sites of one Li-ion on the SLM. For the S and Se side of the SLM, four adsorption sites are considered, respectively. That is top sites of the S/Se atom (denote as TS/Se) and the Mo atom (denote as TMo), hollow site of the hexagon ring (denote as H), and bridge site between the Mo atom and the S/Se atom (denote as B). It is found that the Li-ion only is located at the TS/Se, TMo, and H site, respectively. Instead, Li-ion at the bridge site moves to the top site of Mo atom after full optimization. According to the definition of the adsorption energy in formula 1, we evaluate the preference of the Li-ion adsorption on these sites. Results show that the adsorption energy at the TMo site is larger than those at the TS/Se and H site for the S side and Se side, respectively (see Table 2), indicating that Li-ion prefers to stay at the top of Mo atoms. More importantly, it is noted that Li-ion prefers to stay on the side of the S layer rather than the Se layer. In fact, a single Li-ion cannot be adsorbed on the

Table 2. Adsorption Energies (eV) for All Li-Adsorption Sites of the SLM top site on S/Se (TS/Se) top site on Mo (TMo) hollow site (H) S side Se side

−1.49 −1.05

−2.41 −1.80

−2.25 −1.66

side of the Se layer due to its small adsorption energy, which is discussed in detail later. The optimal adsorption site of Li-ion on the SLM is in good agreement with that on the monolayer MoS2.3939 Thus, the other configurations of Li-ion adsorption on the SLM with different Li concentration are constructed by initially putting Li-ions on the top of Mo atoms on the side of the S layer. The optimal configurations of Li-ion adsorption on the SLM for different Li concentrations are shown in Figure 2c−j. Obviously, for x ≤ 1, all Li-ions prefer to be adsorbed on the top of Mo atoms on the side of the S layer (denote as TMo−S). As for 1 < x ≤ 2, the subsequent Li-ions can be absorbed on the top of Mo atoms on the side of the Se layer (denote as TMo−Se). Here, we note that a single Li-ion cannot be adsorbed on the TMo−Se site with the adsorption energy of 23901


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Figure 3. (a, b) Top view and side view of Li-adsorption sites on the DLM, respectively. (c−j) Top view and side view of the optimal structures of Lix(MoSSe)2.

−1.80 eV, which is higher than the cohesive energy of body center cubic (bcc) Li-metal (−1.92 eV/atom calculated with the GGA-PBE potential in this study). However, the 10th Liion can be adsorbed on the TMo−Se site with the lower adsorption energy of −1.96 eV, which may be due to the dipole moments in the system and the charge redistribution after 9 Li-ions adsorption on the TMo−S sites (Li1MoSSe); the detailed analysis is given in the Discussions section. Similarly, for the case of 2 < x ≤ 3, all Li-ions are located at the hollow sites on the side of the S layer. In the same way, for 3 < x ≤ 4, all Li-ions prefer to locate at the hollow sites on the side of the Se layer. In fact, the Li-ions adsorption sites and configurations are dependent on the dipole moment of the system and the redistribution of electronic charge. For example, for 4 < x ≤ 5, the Li-ions would rather sit on the top of Se atoms than on S atoms, mainly dominated by the dipole moment of system, see

Discussions section. Surprisingly, the SLM accommodates up to six Li-ions per formula unit, corresponding to a chemical ratio of Li6MoSSe; the structure is still nondistorted, see Figure 2j. In this case, the innermost layer adatoms are on the top of Mo atoms, the middle layer adatoms are above the hexagon ring center, and the outmost layer adatoms are on the top of the S/Se atoms. The corresponding Li adsorption energy is −1.94 eV, which is slightly lower than the cohesive energy of the bcc Li-metal, −1.92 eV/atom, indicating that Li adsorption on the SLM with the concentration of x = 6 (54 Li-ions stored in the supercell) is favored, corresponding to a positive discharge potential. When x > 6, the adsorption energy is very close to the cohesive energy of bcc Li-metal, indicating that Limetal formation begins. Thus, x = 6 is considered to be the maximum Li adsorption on the SLM. 23902


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amount of Li-ion that can be stored by the host is determined by the Li-ion adsorption energy (lower than the cohesive energy of Li-metal), and thus the theoretical capacity can be evaluated.40 Here, the adsorption energies for different Li concentration and the cohesive energy of bcc Li-metal are calculated and plotted in Figure 4. As seen from Figure 4, the

Now, we turn to the configurations of lithiated DLM. For one Li-ion adsorption on the DLM, five different adsorption sites are considered in Figure 3a,b. They are top sites of Mo atoms. As there are two types of Mo atoms in the DLM, they are distinguished by TS and TSe, corresponding to the outside surface of the S layer and the Se layer, respectively. Also, there are three adsorption sites in the interlayer of the DLM, namely, the center of the hexagon ring (hollow site, denoted as S1), between the Mo atom and the S atom (denoted as S2), and between the Se atom and the Mo atom (denoted as S3), respectively. The corresponding adsorption energies are calculated and listed in Table 3. It is clear that the S1 site Table 3. Adsorption Energies (eV) for All Li-Adsorption Sites of the DLM TS

TSe

S1

S2

S3

−2.62

−1.83

−2.80

−2.21

−2.43

has the largest adsorption energy of about −2.80 eV, which indicates that Li-ion adsorption on the hollow site is thermodynamically favorable. More importantly, for the top site of Mo atoms, the TS site is shown to have the adsorption energy of −2.62 eV, whereas the TSe site only has the adsorption energy of −1.83 eV. Such a small adsorption energy of the TSe site indicates that a single Li-ion cannot be adsorbed on the side of the Se layer. On the contrary, the adsorption of Li-ion on the side of the S layer is thermodynamically favorable. Then, higher Li-ion adsorption concentrations are further studied. The optimal configurations of Li-ions adsorption on the DLM for different Li concentration are shown in Figure 3c−j, the corresponding chemical formula is Lix(MoSSe)2. Obviously, the hollow sites (T1) remain the most favorable sites for x ≤ 1. When the Li coverage is further increased (1 < x ≤ 2), Li-ions occupy the TS sites. The adatoms gradually occupy the two external surfaces of the DLM with increasing Li concentration. When the Li coverage reaches x = 3, one-third of the Li-ions reside in the hollow sites (T1) and the other adatoms occupy the TS and TSe sites, respectively. For x > 3, the subsequent adatoms adsorb on the outside surface of the DLM. As for the case of x = 5, the configuration of lithiated DLM is similar to that of lithiated BlueP/MS2 (M = Nb, Ta) vdW heterostructures,23 where two layers of Li-ions adsorbed on each side of heterostructure and one layer of Li-ions reside at the interlayer space. We note that there is no structural distortion when the Li coverage reaches a quite large value of x = 7, as shown in Figure 3j. For the case of x = 7, the Li-ions distribution is similar to that for x = 6 for the SLM, with nine Li-ions located at the hollow sites in the interlayer space. For this case, the corresponding adsorption energy is −2.05 eV. The adsorption energy is lower compared with the cohesive energy of bcc Li metal (−1.92 eV/atom), indicating that Liions adsorption on the DLM with x = 7 is favorable. Here, for the lithiated DLM, the Li-ions adsorption configurations are also dependent on the dipole moment of the systems, see the following Discussion section. Generally, if the adsorption energy is lower than −1.92 eV/ atom, Li-ion adsorption on the host is favored, corresponding to a positive discharge potential. Otherwise, Li-metal formation is favored and discharge potential is negative. In a real battery system, formation of Li-metal means a growth of Li-dendrites, which is disadvantage for battery operation. In this sense, the

Figure 4. Adsorption energy as the function of Li content in the LixMSSe and Lix(MSSe)2.

adsorption energies decrease with increase in Li concentration. This is because the attractive interactions between the Li-ions and the host are weakened and the repulsive interactions between the adsorbed Li-ions are strengthened. For the case of Li-ion adsorption on the SLM, the adsorption energies are lower than the cohesive energy of bcc Li-metal until x = 6, indicating that 54 Li-ions can be stored in a 3 × 3 × 1 supercell of the SLM before discharging to 0 V. The corresponding theoretical capacity is 776.5 mAh/g, which is 2 times higher than that of the commercial graphite anode (372 mAh/g).41 As for the case of Li-ion adsorption on the DLM, similar to the SLM, the adsorption energies are shown to be lower than the cohesive energy of bcc Li-metal until the Li concentration reaches x = 7. This implies that 63 Li-ions can be stored in a 3 × 3 × 1 supercell of the DLM with the theoretical capacity of 452.9 mAh/g. The capacity of the DLM is much lower than that of the SLM because the interlayer space of the DLM only can accommodate one layer of Li-ions (nine Li-ions) and the DLM has twice the mass compared with the SLM. From this perspective, the SLM is superior to the DLM as an anode material for LIBs. To visualize the interactions between the adsorbed Li-ions and the host materials, the charge density differences for the lithiated SLM and DLM are calculated and shown in Figure 5. Figure 5a shows the charge density difference of one Li-ion adsorption on either the upper or the bottom of the SLM. Obviously, the Li-ion loses net electronic charge and the SLM gains net electronic charge, indicating that there is significant electronic transfer from the Li-ion to the SLM. As for the DLM, we calculate the charge density difference for one Li adsorption on upper, middle, and bottom of the DLM, respectively, as shown in Figure 5b. Hereafter, for the convenience of description, the upper layer and the bottom layer of the DLM are noted as up-DLM and down-DLM, respectively. Clearly, for one Li-ion adsorption on the upper and the bottom of the DLM, respectively, majority of charge of Li-ion is transferred to the up-DLM and down-DLM, respectively. And in the case of intercalation of one Li-ion in the interlayer of the DLM, net electronic charge is transferred 23903


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Figure 5. Top view and side views of the charge density difference of one Li-ion adsorption on the S side and Se side of the SLM, respectively (a), and the S side, interlayer, and Se side of the DLM, respectively (b). The loss of electrons is indicated in blue and gain of electrons is indicated in yellow.

Li0.11SMoSe system, there is +0.865 |e| on the Li-ion and there are −0.213 |e| and −0.652 |e| on the up-DLM and downDLM, respectively. These net electronic charge distributions reveal that the charge of the Li-ion is predominantly transferred to the adjacent atomic layer. Moreover, when one Li-ion is adsorbed on the outside surface of up-DLM or downDLM, the Bader charge analysis also indicates that obvious charge transfer is from the Li-ion to up-DLM or down-DLM, which again strongly supports the strong ionic interaction between the Li-ion and the SLM or the DLM. This phenomenon is also observed in the BlackP/graphene42 and BlueP/NbS2 heterostructure23 after lithiation. 3.2. Performance Evaluation of the SLM and DLM as Anode Materials. To further understand the performance of MoSSe-based LIBs, in addition to the high theoretical capacity demonstrated above, the open circuit voltage (OCV) and rate performance are also important for them to be a practical anode. According to the definition of OCV in formula 2, the average OCVs for the SLM and DLM have been calculated and plotted in Figure 6. The total energy of the most stable configuration for each Li concentration in the LixMoSSe and Lix(MoSSe)2 is used for computing the voltage values. Moreover, the ideal theoretical capacity as a function of Li concentration is also plotted in Figure 6. As can be seen from Figure 6, the OCV is 0.62 V vs Li+/Li0 for the SLM and 1.01 V vs Li+/Li0 for the DLM, which is comparable to 1.0 V for carbon (LixC6) and 1.12 V for graphite.43 This indicates that

from the Li-ion to both up-DLM and down-DLM. Furthermore, the analysis of charge density difference for the lithiated SLM and DLM demonstrates that there are strong ionic interactions between the Li-ion and the host materials. To further clarify the ion bond strength between the adsorbed Li-ion and the host, we examined the amount of charge transfer in lithiated SLM (Li0.11SMoSe and SMoSeLi0.11 systems) and DLM (Li0.11SMoSeSMoSe, SMoSeLi0.11SMoSe, and SMoSeSMoSeLi0.11 systems) using Bader charge analysis. Table 4 indicates that there is equal amount of positive charge Table 4. Net Charge Transfer (in Electrons) of the Li Ion (Denote as Li), SLM, Up Layer of DLM (Denote as UpDLM), and Bottom Layer of DLM (Denote as Down-DLM) system

Li

SLM

Li0.11SMoSe SMoSeLi0.11 Li0.11SMoSeSMoSe SMoSeLi0.11SMoSe SMoSeSMoSeLi0.11

+0.882 +0.878 +0.882 +0.864 +0.879

−0.882 −0.878

up-DLM

down-DLM

−0.577 −0.213 +0.042

−0.305 −0.652 −0.921

and negative charge on the absorbed Li-ion and the SLM, respectively, demonstrating that the charge of the Li-ion is fully transferred to the SLM. For three cases of lithiated DLM, we also find that the charge of adsorbed Li-ion is transferred to the neighboring atomic layer. For example, for the SMoSe-

Figure 6. Average open circuit voltage (OCV) as the function of Li content in the LixMSSe (a) and Lix(MSSe)2 (b). 23904


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Figure 7. Total DOS of the SLM (a) and DLM (b) and their corresponding lowest and highest lithiated states. The Fermi level is aligned to 0 eV.

Figure 8. Migration path and energy barriers: (a1) and (a2) for one Li-ion migration on the SLM and (b1) and (b2) for one Li-vacancy migration on 17 Li adsorbed on the SLM.

is an ideal candidate with high capacity that can be used as an anode material for LIBs. To evaluate the rate performance of the SLM and DLM as anode materials, both electronic conduction and Li-ion diffusion should be considered. For electronic conduction, we calculate the total density of states (DOS) of the SLM, DLM, and their corresponding lowest and highest lithiated states. As shown in Figure 7, the intrinsic electronic structures of the SLM and DLM are semiconductors with the band gap of 1.55 and 0.81 eV, respectively. Upon Li-ions insertion, the band structures are changed by donating electronic into the

the SLM and DLM could provide a lower discharging voltage. In addition, when the SLM and DLM achieve the highest capacity, corresponding to the Li6MoSSe and Li7(MoSSe)2, the calculated OCV is 0.106 and 0.107 V, respectively, which are positive discharge potentials and very close to 0.11 V for commercial anode material graphite.44 Thus, the applicable OCV provides feasibility for the SLM and DLM to be applied as anodes for LIBs. On the other hand, the ideal capacity of the SLM (776.5 mAh/g) is higher than 669 mAh/g for layered MoS2,17 540 mAh/g for graphene nanosheets,45 and 528.257 mAh/g for BlueP/NbS2 vdW heterostructures.23 So, the SLM 23905


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Figure 9. Migration path and energy barriers: (a1) and (a2) for one Li-ion migration on the DLM and (b1) and (b2) for one Li-vacancy migration on 26 Li adsorbed on the DLM.

Li-ion diffusion on the monolayer graphene (0.28−0.31 eV)42,46 and silicene (0.22−0.27 eV),47,48 our calculated energy barriers suggest that the rate performance of the SLM and DLM is comparable to that of the graphene and silicene. From our calculations on the same theoretical level, the energy barrier of Li-ion migration on the monolayer MoS2 is 0.27 eV, which is between the barrier of the S layer and the barrier of the Se layer and slightly higher than previous reports (0.21− 0.24 eV).39,49,50 For Li-vacancy migration on the SLM, different from Li-ion migration, the pathway looks like an arc of circle and the migration barriers are 0.44 and 0.34 eV for the S layer and Se layer, respectively. This indicates that Li-vacancy diffusion is faster on the Se layer than on the S layer and implies that the diffusion of Li-ions slows down with the increase in Li-ion concentration. Next, we investigated the migration paths and energy barriers of Li-ion and Li-vacancy migration on the DLM. As can be seen from Figure 9, three migration paths are considered, namely, Li-ion or Li-vacancy migration on the S layer, middle, and Se layer. For Li-ion migration on the S layer and Se layer, because the most stable site is also on the top of Mo atom, the migration path is similar to that of Li-ion migration on the SLM. As for Li-ion migration on the middle (interlayer space), because the most stable position is the hollow site, the migration path is from one hollow site to the neighboring one passing through a metal-stable site S3. The calculated energy barriers are 0.28, 0.43, and 0.23 eV for Li-ion migration on the S layer, middle, and Se layer, respectively. As for Li-vacancy migration on the DLM, the migration paths are

SLM or DLM, which induces a semiconductor−metal transition. Moreover, the electronic states at the Fermi level increase with increase in Li concentration, implying that more active electrons are available in the system, and, therefore, better electronic conductivity can be expected. Because the electronic conduction in the SLM and DLM is good, the diffusion of Li-ion in the SLM and DLM becomes critical to the anode performance. As we know, the diffusion coefficient and migration energy barrier of Li-ion are dependent on Li concentration. For simplicity, here we study two extreme cases: Li-ion and Li-vacancy migration. For Li-ion migration, one Li-ion is adsorbed on a supercell of the SLM or DLM, whereas for Li-vacancy migration, one Li-ion is removed from a supercell of the SLM or DLM with large Li concentration. Figure 8 shows the migration paths and energy barriers of one Li-ion and Li-vacancy migration on the SLM, respectively. For comparison, the migration energy barrier of one Li-ion on the monolayer MoS2 is also plotted in Figure 8a1. Here, for the convenience of description, we denote Li-ion or Li-vacancy migration on the side of the S atoms as S layer and on the side of the Se atoms as Se layer. For Li-ion migration, two migration paths are considered based on the symmetry of the SLM, namely, Li-ion migration on the S layer and Se layer. Li-ion from one favorable site (TMo) moves to the nearest neighbor passing through a metal-stable site (H). The migration path is similar to the zigzag direction. Our calculated migration energy barriers for the S layer and Se layer are 0.29 and 0.24 eV, respectively. This indicates that Li-ion diffusion is faster on the Se layer than on the S layer. As compared to the 23906


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The Journal of Physical Chemistry C Table 5. Dipole Moments (in Debye) of LixMoSSe and Lix(MoSSe)2 Systems with Different Li Content LixMoSSe Lix(MoSSe)2

x=0

x=1

x=2

x=3

x=4

x=5

x=6

x=7

−1.61 −2.99

0.78 −1.88

−0.03 0.53

1.08 −0.03

0.24 1.03

−0.25 0.27

0.03 −0.08

0.02

similar to those of Li-ions migration on the DLM, and the corresponding energy barriers are 0.42, 0.36, and 0.31 eV for Li-vacancy migration on the S layer, middle, and Se layer, respectively. Obviously, both the Li-ion diffusion and Livacancy migration are easy on the Se layer than on the S layer. In addition, with the increase in the intercalation of lithium ions, the charge of the system will be redistributed and the dipole moment will be changed, which can affect the migration barrier of lithium. Therefore, the largest barrier occurs on the middle for single Li-ion migration and on the S layer for single Li-vacancy migration. Compared to the SLM, the migration energy barriers for both Li-ion and Li-vacancy are comparable. Moreover, for the SLM and DLM, the energy barriers of Li-ion diffusion slightly increase with the increase in Li concentration. 3.3. Discussions. As mentioned above, the ideal Li-ion storage capacity of the SLM and DLM are 776.5 and 452.9 mAh/g, respectively. Such large capacities can be attributed to the role of intrinsic dipole moments and charge redistributions of the systems. To understand the effect of dipole moments on the adsorption of Li-ions, the dipole moments of LixMoSSe and Lix(MoSSe)2 systems with different Li concentrations are calculated and listed in Table 5. The negative sign represents the direction of dipole moment pointing to the S layer from the Se layer. Clearly, when x = 0, the pristine SLM and DLM have the intrinsic dipole moments of −1.61 and −2.99 Debye, respectively, which induced an internal electric field in the direction from the Se atom layer to the S atom layer. So, the Li-ions prefer to sit on the TMo−S sites with x ≤ 1 for the SLM and the S1 and TS sites with x ≤ 2 for the DLM, respectively. When x = 1 for the SLM and x = 2 for the DLM, the electronic charge of the systems is redistributed due to the introduction of extra electrons after Li-ions adsorption, which leads to the change in dipole moments and engenders the opposite dipole moments. The dipole moments of Li1MoSSe and Li2(MoSSe)2 are 0.78 and 0.53 Debye, respectively, which induced an internal electric field pointing to the Se layer from the S layer. This internal electric field enables the Li-ions adsorption on the TMo−Se sites for SLM with 1 < x ≤ 2 and the TSe sites for DLM with 2 < x ≤ 3, respectively. Similarly, for Li2MoSSe and Li3(MoSSe)2, the negative dipole moments allow the subsequent adatoms to locate on the side of the S layer. In the same way, for Li3MoSSe and Li4(MoSSe)2, the positive dipole moments allow the subsequent adatoms to locate on the side of the Se layer. Interestingly, for Li4MoSSe and Li5(MoSSe)2, the dipole moments are still positive. Therefore, the subsequent Li-ions prefer to still reside on the side of the Se layer (see Figures 2i and 3i). In like manner, for Li5MoSSe and Li6(MoSSe)2, the negative dipole moments enable the subsequent Li-ions to locate on the side of the S layer. To further understand the change in dipole moments, we examined the charge density difference between lithiated SLM and DLM, respectively. Typically, we use the Li1MoSSe and Li2(MoSSe)2 system as examples to illustrate the electronic charge redistribution. As shown in Figure 10a, after adsorption of nine Li-ions on the side of S layer, Li atoms are ionized into positive Li-ions due to loss of electrons. It is clear that there is net electronic charge transfer from Li-ions to the SLM.

Figure 10. Top view and side view of the charge density difference of Li1MoSSe (a) and Li2(MoSSe)2 (b). The loss of electrons is indicated in blue and gain of electrons is indicated in yellow. The isosurfaces value is 0.002 e/Å3.

Therefore, there is an intrinsic dipole moment of 0.78 Debye pointing to the Se layer from the S layer. This is why a single Li-ion cannot adsorb on the side of the Se layer, whereas the 10th lithium-ion can adsorb on the side of the Se layer. In the same way, Figure 10b indicates that there is net electronic charge transfer from the Li-ions to the DLM, resulting in an intrinsic dipole moment of 0.53 Debye pointing to the Se layer. In brief, for the polarized 2D materials of pristine SLM and DLM, the intrinsic dipole moment allows the Li-ions adsorption on the side of the S layer, then the charges of system redistribute after Li-ions adsorption, causing the change in dipole moment. The new dipole moment enables the more Li-ions adsorption on the systems. As a result, much more Liions can be stored by the SLM and DLM before intercalation potential is decreased to 0 V.

4. CONCLUSIONS In summary, the single-layer Janus MoSSe and the double-layer MoSSe (SMoSeSMoSe) as anode materials for LIBs are investigated for the first time by the first-principle calculations. Our present calculations suggest that the 2D SLM and DLM could be potential anode materials for LIBs with excellent performance. Due to the intrinsic dipole moment in the pristine SLM and DLM and the lithiated SLM and DLM, much more Li-ions can be stored by the SLM and DLM, respectively. With a suitable open circuit voltage range vs Li+/ Li, the ideal theoretical capacities for the SLM and DLM achieve 776.5 and 452.9 mAh/g, respectively. The analysis of the electronic structures indicates that the lithiated SLM and DLM are metallic, ensuring a good electrical conduction. The diffusion of the Li-ion and Li-vacancy is easy on the Se layer than on the S layer. The small Li-ion/Li-vacancy migration barriers enable good Li-ion conduction. These features are critical attributes for the SLM and DLM to achieve good rate performances as an anode material for LIBs. 23907


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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.jpcc.8b07478. Total energy vs cutoff energy; adsorption energy (eV) of one Li-ion adsorbed on the monolayer and bilayer (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Xueling Lei: 0000-0002-2482-3728 Bo Xu: 0000-0002-6896-0409 Gang Liu: 0000-0003-3213-3820 Chuying Ouyang: 0000-0001-8891-1682 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors thank the National Natural Science Foundation of China (Grant Nos. 11764019, 11664013, 11664012, and 11564016) for financial support of the current work.

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REFERENCES

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Theoretical Prediction of Janus MoSSe as a Potential Anode Material

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