Defect Induced Performance Enhancement of Monolayer MoS2 for Li

Subscriber access provided by Nottingham Trent University

C: Energy Conversion and Storage; Energy and Charge Transport

Defect Induced Performance Enhancement of Monolayer MoS for Li- and Na-Ion Batteries 2

Gayatree Barik, and Sourav Pal J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.9b04128 • Publication Date (Web): 14 Aug 2019 Downloaded from pubs.acs.org on August 15, 2019

Just Accepted “Just Accepted” manuscripts have been peer-reviewed and accepted for publication. They are posted online prior to technical editing, formatting for publication and author proofing. The American Chemical Society provides “Just Accepted” as a service to the research community to expedite the dissemination of scientific material as soon as possible after acceptance. “Just Accepted” manuscripts appear in full in PDF format accompanied by an HTML abstract. “Just Accepted” manuscripts have been fully peer reviewed, but should not be considered the official version of record. They are citable by the Digital Object Identifier (DOI®). “Just Accepted” is an optional service offered to authors. Therefore, the “Just Accepted” Web site may not include all articles that will be published in the journal. After a manuscript is technically edited and formatted, it will be removed from the “Just Accepted” Web site and published as an ASAP article. Note that technical editing may introduce minor changes to the manuscript text and/or graphics which could affect content, and all legal disclaimers and ethical guidelines that apply to the journal pertain. ACS cannot be held responsible for errors or consequences arising from the use of information contained in these “Just Accepted” manuscripts.

is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

Page 1 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

The Journal of Physical Chemistry

Defect Induced Performance Enhancement of Monolayer MoS2 for Li- and Na-Ion Batteries Gayatree Barik1 and Sourav Pal1,2* 1Department

of Chemistry, Indian Institute of Technology Bombay, Mumbai 400076, India

2Department

of Chemical Sciences, Indian Institute of Science Education and Research Kolkata, 741246, West Bengal, India Email: [email protected]

Abstract

Dexterity in the application of defect engineering implicates modification in physical properties of two-dimensional (2D) TMDs to enhance their effectiveness towards nanoelectronics applications. Subsequently, the existence of a various type of defects in monolayer MoS2 has been employed to inculcate and implement their significance in enhancing the Li/Na-ion storage capability of MoS2 monolayer as anodes of LIBs/SIBs. DFT calculations have guided us to traverse the effect of various point and anti-site defects on Li/Na adsorption energy and diffusion barrier of monolayer MoS2. Before looking into Li/Na adsorption into defective MoS2, the structural stability of various defects is explored with relevant to their formation energy. This also germinates the quest for the most stable defective structures that could be a reliable anode material for LIBs and SIBs. Enhanced adsorption is found for both Li/Na ions in case of defective MoS2 than corresponding pristine MoS2. To study the level of interaction between Li/Na and defective MoS2 electronic structure analysis has been performed. To evaluate the possible migration pathways and rate of migration of Li/Na over defective MoS2, we calculated diffusion barrier energy through CI-NEB method. Our study demonstrates that the formation of vacancy improves the diffusion performance of both lithium and sodium at the defective region which are prerequisites for LIBs and SIBs. Additionally, we demonstrated that the formation of vacancy could improve the specific capacity of monolayer MoS2 due to a decrease of the molecular mass of defective MoS2 in comparison to pristine MoS2. However, the OCV is not affected much due to enhanced adsorption. Hence, designing MoS2 nanostructures with defects is a useful strategy to achieve an effective anode material for obtaining high capacity LIBs and SIBs by precisely tailoring its properties for desired applications, such as enhancing the adsorption energy, modulating the reaction path way and raising the specific capacity.

1 ACS Paragon Plus Environment

The Journal of Physical Chemistry 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

1. Introduction Over the past decades, rechargeable Li-ion batteries (LIBs)1 has been playing a significant role by virtue of their portability in electronic devices and high-power density. In contemporary days, it is a matter of worries that accessible Li resources from the earth crust would not be adequate to meet the perennial requirements for LIBs. Notwithstanding, because of their inherent security issues and relatively high cost, Li confines the use of LIBs in large scale Nanodevice applications.2 The aforementioned apprehensiveness has prompted to search for suitable alternatives. Compared to LIBs, SIBs3 are rapidly emerging4 as a plausible alternative to LIBs for energy storage, because SIBs accommodate Na which is available in abundance and is also associated with economic advantages as well.5 So, to fulfil the growing demand for modern electronic devices, improvements in the capacity of modern LIBs and SIBs are of paramount significance. Additional improvements in the capacity of modern LIBs and SIBs can be made possible by enhancing adsorption capabilities and diffusion rate of anode materials. In the said context, 2D TMDs6,7 and monochalcogenides8 are most widely used anode material and have pulled in significant consideration as efficient electrode materials. As a standout amongst the most intriguing TMDs, the electrochemical performance of MoS2 towards Li/Na intercalation9,10 has been extensively explored. MoS2 has been used as an exceptional substitute anode material for high capacity LIBs/SIBs in modern electronic devices due to its novel layered structure and tunable physical properties with the substantial interlayer distance which permits quick diffusion of adsorbed atoms that impoverishes volume expansion and also contributes towards its extensive usage. The electronic properties MoS2 can be extended further for its capable usage in future Nano-electronic applications by the formation of vacancies. Further, this may initiate a broad investigation of its physical properties towards its potential usage in LIBs/SIB.11,12 Presence of defects are anticipated to play a key job in determining the impact of defects on the behaviour of MoS2.13,14 Extensive studies have been explored in order to uncover defects in the 2H phase of monolayer MoS2 both by theoretically15–17 and experimentally.13,14,18,19 The presence of different kind of defects in MoS2 distinguishes itself from a defect-free MoS2 by altering the original properties like electronic,20 optical,21 magnetic22 and chemical23 properties of materials. A specific number of defects and impurities are always present in crystalline materials naturally. Defective sites are comparatively more reactive than the perfect sites and 2 ACS Paragon Plus Environment

Page 2 of 32

Page 3 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


even with small defects, the electrical and optical properties, such as the electrical conductivity, are influenced to a large extent. The implication of defect engineering is beneficial for structural designing of materials since the existence of defects leads to the cracking of crystal lattice and subsequently increases the surface area.24 Introduction of defects creates more active sites in the basal plane of MoS2 which leads to significant improvements in the catalytic activity of the Hydrogen evolution reaction (HER).25 Additionally, catalytic activity for the hydrodeoxygenation (HDO) reaction of MoS2 improved significantly through the formation of S-vacancies.26 A fundamental understanding of defect induced Li/Na intercalation capabilities of the monolayer MoS2 towards improving the reliability of LIBs/SIBs is of great importance for large scale energy storage. Many studies have revealed that defects play a key role in improving material intercalation capabilities of electrode materials in batteries. Introduction of defects enhances material intercalation capacities which in turn enhances device performance, by creating unprecedented functionalities which can directly modify both chemical and physical characteristics. DFT investigations recommend that Li storage limits of graphene27,28 can be improved by the presentation of impurities and defects. Surface imperfections can specifically affect alkali metal ion intercalation by moving the thermodynamics and increases surface energy.29,30 The idea of cooperation of intercalating particles with the intrinsic structural defects in MoS2 was studied by DFT calculations and progressive Li+ addition into nonstoichiometric, S- excess MoS2 prompts a semi-conductor to metal electronic phase change.31 Experimental findings suggest that the electrode performance of graphene in LIBs can be upgraded by presenting intrinsic or extrinsic defects.32,33 The defects in 2D materials change the electronic properties by creating active sites for substantial alkali-metal adsorption. The rearrangement of atoms in the region of the defect leads to enhanced binding of alkali metal ions with the defective material.34 Liu et al. investigated that defect-induced Li adsorption and storage capacity of Ti2C monolayers has a significant influence on LIBs anode materials.35 Sun et al.36 revealed that occurrence of defects introduces significant strong binding of Li to the monolayer MoS2 whereas Li-ion migration ability remains relatively unaffected. On the other hand, Zhang et al.37 conveyed that the presence of defects stops the diffusion of Li ions over monolayer black phosphorous and at the same time, the Open circuit voltage (OCV) expanded enormously as the defects appeared. The adsorption capability of alkali metal ions on defective ReS2 monolayer is found to be stronger 3 ACS Paragon Plus Environment


than defect-free ReS2 with the influence of atomic defects.38 Mg adsorption on pristine graphene is not suitable, however defective graphene and graphene allotropes with high vacancies essentially improve Mg storage capacities.28 Similarly, Na and Ca adsorption capability of graphene can be effectively tuned by introducing defects in graphene on account of expanded charge transfer.39 Fan et al.27 have demonstrated that the defect-free graphene can’t upgrade the diffusion energetics of Li-ion. Furthermore, with the formation of defects, make easy diffusion of Li on the surface of graphene. Broad examinations have been undertaken for utilizing MoS2 as electrode material in both LIBs and SIBs. However, studies dealing with intercalation of Li/Na ions on defective MoS2 has been confined to limited studies. So far Sun et al.36 have observed improved lithiation because of the presence of vacancies in contrasted with those on the pristine MoS2. The idea of incorporation of intercalating particles with fundamental structural defects and its role in affecting electronic properties towards device performance applications such as LIBs/SIBS is not known. In the present work, we uncover a broad-range of reversible modulation of structural and electronic properties of MoS2 which is possible through the careful introduction of predominant types of defects in the crystal structure MoS2 by using first-principles DFT calculations. We study a complete theoretical examination of the impact of various type of defects on monolayer MoS2 used in LIBs and SIBs. 2. Computational Details All simulations are performed by the use of density functional theory (DFT) which is implemented in the Quantum ESPRESSO40 package. The exchange-correlation was approximated using the generalized gradient approximation (GGA)41 with Perdew-BurkeErnzerhof (PBE) functional. The atomic positions are optimized using the kinetic energy cut off of 400 eV in the plane-wave expansion. Electron-ion interaction was described by the projector augmented wave (PAW)42 method. The Brillouin zones of 3×3 supercell and 2×2 supercell were modelled using a special kpoint sampling of the Monkhorst-Pack scheme43 with 4×4×1 and 6×6×1 grid of points. The 3×3 hexagonal supercell of monolayer MoS2 was utilized to display various defects in MoS2 and the Li/Na adsorption and diffusion whereas a 2×2 hexagonal supercell is employed to study open-circuit voltage (OCV) and specific capacity. A vacuum spacing of 20 Å was provided along a perpendicular direction to the plane of MoS2 between two adjacent periodic layers in 4 ACS Paragon Plus Environment

Page 4 of 32

Page 5 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


order to avoid any spurious interactions. A convergence threshold of 0.02 eV Å−1 in force is required to reach relaxed configurations. The van der Waals interaction is considered using a dispersion correction term with the vdW-DF44 method. To quantitatively quantify the charge transfer between Li/Na and MoS2, Bader45 charge analysis method has been adopted. The diffusion barrier energies for Li/Na diffusion atoms over MoS2 monolayers with defects along with associated minimum energy pathway (MEP) are calculated by the CI-NEB46 method. The unit cell dimensions of monolayer MoS2, a = b = 3.19 Å and c = 12.294 Å, has been taken.47 The stability of various defects is calculated from their formation energy which is defined as defect ― 𝐸Pristine ± ∑𝑛 𝜇 𝐸defect 𝑖 𝑖 form = 𝐸tot tot

(1)

where 𝐸defect is the total energy of MoS2 with defects, 𝐸Pristine is the total energy pristine MoS2. tot tot ni is the total number of Li/Na atoms added or subtracted to form vacancy, and μi is the chemical potential of individual species i. The chemical potentials 𝜇Mo 𝑎𝑛𝑑 𝜇𝑆 satisfies the relation, 𝜇Mo + 2 × 𝜇𝑆 = 𝜇MoS2 where 𝜇Mo, 𝜇𝑆 are the chemical potentials of Mo, S respectively and 𝜇MoS2 is the chemical potential of pristine monolayer MoS2 per formula unit. To analyse the binding interaction of Li/Na adsorbed on MoS2 with various defects, the adsorption energy ( 𝐸ad ) of a Li/Na adatom was calculated at most favourable adsorption site using 𝐸ad = 𝐸MoS2 ― Li Na ― 𝐸MoS2 ― 𝐸Li Na

(2)

where 𝐸MoS2 ― Li Na is the total energy of Li/Na adsorbed MoS2 with defects, 𝐸MoS2 is the total energy of Li/Na free MoS2 with defects, 𝐸Li Na is the energy of an isolated Li/Na atom. According to the definition, a more negative binding energy determines a favourable exothermic reaction between the monolayer defective MoS2 and Li/Na. To study the interaction of Li/Na over defective MoS2, we calculated the charge density (ρad) isosurface plots which are given by



𝜌ad = 𝜌MoS2 ― Li Na ― 𝜌MoS2 ― 𝜌Li Na

Page 6 of 32

(3)

In the above equation 𝜌MoS2 ― Li Na indicates the charge density of Li/Na adsorbed MoS2 monolayer, 𝜌Li Na indicates charge density of an isolated Li/Na atom, and 𝜌MoS2 signifies the charge density of pristine MoS2. The open-circuit voltage (OCV) depends on Gibb’s free energy of formation and is calculated as follows OCV  

G form

(4)

zF

where Gibb’s free energy change can be defined as G form  E form  PV form  TS form

(5)

The volume and entropy terms are generally considered insignificant during the reaction,48 and hence the effect of entropy and the volume terms are ignored. So, the Gibbs free energy G form becomes equal to the formation energy E form which is calculated from our calculations. The formation energy is calculated as the following way E form 

EMoS2  Li / Na  EMoS2  n  ELi / Na n

(6)

where 𝐸MoS2 ― Li Na is the total energy of the Li/Na adsorbed defective MoS2, n × 𝐸Li Na indicates the total energy of Li/ Na atoms, 𝐸MoS2 represents the total energy of Li/Na free defective MoS2 and n is the number of Li/Na adsorbed. The specific capacity (Q) of a material used in battery is determined by using the following equation, given by

Q

1000  F  z  nLi / Na nMoWMo  nSWS

(7)

where F is the Faraday’s constant, 𝑛Li Na is the total number of electrons, 𝑛Mo , 𝑛𝑆 represents the number of Mo, S atoms in the cell, 𝑊Mo and 𝑊𝑆 represents the molar mass of Mo and S respectively and z is the valence number. 3. Results and Discussion 6 ACS Paragon Plus Environment

Page 7 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


3.1 Stability of various defects and defect formation energies: Before investigating adsorption capability of both Li and Na on defect-free and defective MoS2, we initially talk about their relative strength so as to look for the most reliable defective structures which could act as a reticent substitute for anode materials of LIBs and SIBs. We consider various types of vacancy defects, as well as anti-site defects and investigate their conceivable impacts on adsorption and diffusion of Li/Na to make it for the best use as anodes for LIBs and SIBs. We have taken into account the most usually observed point defects in the monolayer MoS213,14,18 including mono-sulphur vacancy (V1S), di-sulphur vacancy (V2S), vacancy of Mo (VMo), Mo vacancy with three sulphur nearby it (VMoS3), vacancy complex of Mo and nearby six sulphur (VMoS6) and two types of anti-site defects. One anti-site defect is Mo atom replacing an S2 column (MoS2) and another one is S2 column replacing a Mo atom (S2Mo). In order to understand the relative strength under thermodynamic equilibrium, the stability of various defects is investigated through defect formation energies, which is calculated with standard chemical potential of bulk Mo and sulphur by equation 1. In this work, bcc crystal of bulk molybdenum and an orthorhombic lattice of bulk sulphur in α phase49 have been chosen. Eight sulphur atoms are connected in a puckered ring having S-S bond distance of 2.13 Å. The corresponding total energies of bulk sulphur and molybdenum are obtained by full structural relaxation. As indicated by our definition, an increasingly positive value of defect formation energy indicates thermodynamically unfavorability and would thus be more unlikely to exist. Molybdenum disulphide (MoS2) is one of most studied two-dimensional material which has a lamellar shape with one Mo layer is stacked by two sulphur atomic layers arranged in an ABA stacking sequence via strong covalent bonding and each layer are arranged by weak van der Waals forces.50 The optimized lattice constant for MoS2 is 3.19 Å having each Mo-Mo interatomic distance 3.2 Å with a Mo-S bond length of 2.44 Å and thickness of 1.25 Å (Figure 1a) well agreed with the previous report.51 Mono-sulphur vacancy (V1S) arises from a missing of one sulphur atom, from any one of the sulphur atomic layers of MoS2 monolayer (Figure 1b) which are usually produced by



Figure 1: Optimized geometry of various defects in monolayer MoS2. (a) Pristine MoS2, (b) V1S, (c) V2S, (d) VMo, (e) VMoS3, (f) VMoS6, (g) MoS2 and (h) S2Mo vacancy. The purple and yellow coloured balls denote Mo and S atoms respectively. The top layer of S atoms is in golden yellow colour and bottom layers are in light yellow colours. In MoS2 vacancy, the maroon colour ball represents Mo present in 8 ACS Paragon Plus Environment

Page 8 of 32

Page 9 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


anti-site substituting a sulphur group. In S2Mo vacancy, the grey colour ball represents the S present at anti-site substituting a Mo group.

electron bombardment52 and di-sulphur vacancy (V2S) is formed by eliminating two sulphur atoms, one from the top and one from bottom layers (Figure 1c) of MoS2 monolayer. The calculated formation energies for V1S and V2S are 1.95 eV and 3.78 eV, respectively which is found to be lowest among all type of defects. These results well agreed with previous results.14 Experimental52,53 and theoretical studies16,17,54,55 of point defects features the significance of sulphur vacancy defects V1S and V2S, which are abundant because of their low defect formation energy. However, in the case of graphene, mono-vacancy is generally less stable than divacancy because of the presence of dangling bonds.34,56 The optimized lattice constant reduces to 3.18 and 3.15 Å for mono and di-sulphur vacancy respectively (Table S1). When a mono-sulphur vacancy is formed by removing S atom, the Mo-Mo bond length between the three Mo atoms affected by the vacancy site reduces from 3.20 Å to 3.04 Å. Similarly, in case of di-sulphur vacancy, we found that the Mo atoms move inward in the vicinity of the vacancy site with each Mo-Mo distance of ∼2.85 Å. Similar observations were produced by previous work 54 when studied systematically various defects in TMDs. Besides, the formation energy of V2S is observed to be generally twice to that of V1S, proposing that V1S ’s do not possess any tendency to combine, which is a different situation in contradiction to graphene, in which divacancies are enthusiastically supported over monovacancies.57,58 If we look through the Mo vacancy defect, they are formed by the removal of one Mo atom from the monolayer of MoS2 (Figure 1d). The vacancy formation energy of single Mo is 5.73 eV, larger than both mono and di sulphur vacancy. As the Mo atom is connected to six S atoms, removal of one Mo from a monolayer of MoS2 affects coordination sphere of six S atoms, for which Mo vacancy defect needs large formation energy. This agrees with experimental reports, where Mo vacancy needs high formation energy than other types of vacancies.14 A complex type of defect is also observed in monolayer MoS2. One is Mo with three sulphur around it (VMoS3) and another is Mo with three di-sulphur pairs (VMoS6) near it. VMoS3 defect is observed by removing one Mo and three neighbouring S atoms from one layer of MoS2 (Figure 1e) from the same z-axis. Similarly, for VMoS6 vacancy, three sulphur atoms are 9 ACS Paragon Plus Environment


removed from one layer and three sulphur from another layer (Figure 1f). The vacancy formation energy for MoS3 and MoS6 vacancy is 7.90 and 13.89 eV respectively. Another important defect observed in monolayer MoS2 is anti-site defects. The first type of anti-site defects is, one Mo atom substituting two S atoms, MoS2 (Figure 1g) and the second one is, two S atoms shifting the Mo position, namely S2Mo (Figure 1h). The calculated defect formation energy for MoS2 and S2Mo are 6.82 and 9.33 eV respectively. The 3-fold axis of symmetry is maintained by S2Mo defect structure whereas the 3-fold axis of symmetry is, unequivocally broken by MoS2 defects with Mo atom at anti-site moves away from Mo-plane and towards the S-atomic layer. The formation energy of MoS2 and S2Mo anti-site defects are found to be higher than other categories of defects. This is because the formation of anti-site defect needs two steps namely vacancy formation followed by addition of an adatom in the defective site. The calculated vacancy formation energy and optimized lattice constant for all types of defects are given in Table S1. 3.2 Adsorption of Li/Na on Pristine and Defective MoS2: MoS2 is one of the most typical metal dichalcogenide and by virtue of its unique layered structure, offers enormous advantages in the field of nanoelectronics and towards electrochemical energy storage as well. This unique structure with a large interlayer distance accommodates the alkali metal ions in the interlayers which demonstrates excellent performance in anode material applications including lithium- and sodium-ion batteries.59,60 It has been accounted that vacancy defects and grain boundaries in graphene increase the adsorption capability of both Li61 and Na62 ions and for which storage capacity of a material used in alkali metal ion battery is enhanced. There are various types of adsorption sites available for Pristine MoS2 favourable for lithiation/sodiation. To know the adsorption capability of both Li/Na on pristine MoS2, we consider two sites; one is just above the top of a Mo atom (known as Mo Top site) and another one at the mid of the centre of a hexagonal ring (Hollow site). Our calculated results show that top site is the most favourable location (Figure S1a and S2a) for both lithiation and sodiation over pristine MoS2 with Li adsorption energy 2.08 eV and Na adsorption energy -1.28 eV. This result agrees well with previous calculations.11,12 Initially, lower energy adsorption for binding sites was examined towards intercalation of an isolated Li/Na atom in defective MoS2. It has also been brought out by researchers with 10 ACS Paragon Plus Environment

Page 10 of 32

Page 11 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


regard to 2D materials, that metal adatom gets adsorbed at the proximity of defective zone, due to very availability of unused electron from the dangling bond.37,62 Therefore we first checked adsorption energy at defective regions. In case of mono-sulphur vacancy Li/Na atoms tend to adsorb at the vacancy position (Figure S1b and S2b) with adoption energy of -2.57 eV and 1.99 eV for Li and Na respectively, which is much higher than Li/Na absorption energy of pristine MoS2. Similarly, in case of di-sulphur vacancy Li and Na adsorption energies are -2.44 eV and -1.89 eV respectively and both Li/Na tries to remain at one side of sulphur vacancy region (Figure S1c and S2c). However, adsorption of a Li over VMo is found to be highest with an adsorption energy of -3.92 eV and Na adsorption energy is -1.81 eV. Similar result was produced when Li/Na adsorption takes place in monolayer ReS2 in which the adsorption energy of a Li with VRe is of -3.17 eV is strongest among other vacancies in compared to pristine ReS2.38 In case Mo vacancy both the metal atoms prefer to stay on the Mo plane of MoS2 (Figure S1d and S2d). In case of MoS3 vacancy, the calculated adsorption energies are -2.83 and 2.09 eV for Li and Na respectively, but both the metal ions prefer to adsorb away from the Mo plane of MoS2 towards sulphide vacancy region (Figure S1e and S2e). However, in case of MoS6 vacancy, both the metal ions stay at the Mo plane (Figure S1f and S2f) with Li adsorption energy of 3.37 eV and Na adsorption energy of -2.28 eV. Anti-site defects are substitutional defects. As discussed above MoS2 breaks the 3-fold symmetry and Mo atoms try to stay away from the Mo plane, Li/Na metal ions prefer to adsorb at the opposite side of Mo (Figure S1g and S2g) with an adsorption energy of -2.58 eV for Li and -1.99 eV for Na. In case of S2Mo, Li/Na adsorption energy at above the sulphur atom was calculated with adsorption energy for Li is -2.28 eV for Na it is -1.60 eV (Figure S1h and S2h). The high negative value of adsorption energies indicates that the formation of vacancies in monolayer MoS2 promotes enhanced binding energy between monolayer MoS2 and Li/Na, and energetically favourable chemical interaction between Li/Na and layered MoS2. In this manner, it is obvious that the presence of vacancies upgrades the adsorption of Li/Na on the surface of MoS2. These outcomes are in steady with the previously produced results of Sun et al.36 in which, they demonstrate that the existence of a vacancy in MoS2 can improve the adsorption ability of Li. We have also calculated the lattice constant after lithiation/sodiation and the lattice constant for defective MoS2 increases after lithiation/sodiation (Table 1)



Figure 2: Adsorption of Li/Na at the defective site and top site (a) Adsorption Energy at Defective site (b) Adsorption Energy at Top site.

Many studies relative to adsorption of alkali metal on defective systems have shown that possible positions for the alkali metal adsorption are at the defective sites. However, it is pertinent to check the adsorption of adatoms at the top site of Mo (most stable adsorption site for pristine MoS2) in the vicinity of defects. In this way, to inspect the impact of different types of vacancy on adsorption properties of the Li/Na in defective MoS2, adsorption energies for Li/Na near the vacancy are determined. As the top site is the most favourable site for lithiation/sodiation in pristine MoS2, we checked the adsorption energy at the top site for all the defective systems. In comparison to pristine MoS2, the regular adsorption positions for Li/Na in the circumference of vacancy are also equally stable for adsorption of both Li/Na ions. The adsorption energy at defective sites and top sites are shown in the histogram plots (Figure 2a and 2b). The values of the adsorption energies of the Li/Na at the defective MoS2 are expanded more in compared to the pristine one. Therefore, vacancy not only promotes the adsorption positions of both Li and Na atoms at the defective site but also at the top sites in the MoS2 monolayer which recommend that the presence of vacancies can improve the storage capacity of monolayer MoS2 which is a fundamental necessity for LIBs/SIBs. The adsorption energy, adsorption distance, charge transfer and optimized lattice constant at the defective site are shown in Table 1. In the below table, * represents the perpendicular distance of Li/Na from Mo-plane. To acquire an understanding of the interaction of Li/Na ions on defective MoS2, bonding charge-density analysis is carried out. Figure S3 shows charge-density difference plots, plotted 12 ACS Paragon Plus Environment

Page 12 of 32

Page 13 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


between intercalant and the adsorbent at defective MoS2 site. The bonding charge density is obtained by subtracting charge density of Li/Na free defective MoS2 and isolated Li/Na atoms Table 1: Calculated Adsorption energies (Ead) for Li/Na, Bader Charge transfer (|e|) from Li/Na, Optimized Lattice constant (a0), after lithiation and sodiation, the minimum distance between Li/Na atom and neighbouring Mo atom (dLi−Mo), S atom (dLi−S), for Li/Na adsorbed on various types of vacancy in monolayer MoS2.

Lithiation Vacanc E ad y (eV)

Charge Lattice

d Li−Mo dLi−S

transfer constant (Å) (|e|)

Pristine -2.08

Sodiation

(Å)

Ead

Charge Lattice

(eV)

transfer constant (Å)

a0 (Å)

(|e|)

dLi−Mo

dLi−S (Å)

a0 (Å)

0.89

3.23

3.17*

2.41

-1.28

0.79

3.24

3.77*

2.81

MoS2 V1S

-2.57

0.92

3.19

2.60*

3.19

-1.99

0.86

3.19

2.85*

3.28

V2S

-2.44

0.91

3.15

2.62*

3.20

-1.89

0.85

3.15

2.92*

3.32

VMo

-3.92

0.85

3.20

0.00*

2.47

-1.81

0.76

3.21

0.00*

2.57

VMoS3

-2.83

0.86

3.29

0.45*

2.49

-2.09

0.78

3.19

1.64*

3.45

VMoS6

-3.37

0.87

3.13

0.02*

3.76

-2.28

0.66

3.15

0.00*

3.77

MoS2

-2.58

0.91

3.24

2.55*

3.21

-1.99

0.85

3.26

2.85*

3.31

S2Mo

-2.28

0.92

3.29

-----

2.27

-1.60

0.68

3.28

----

2.65

from lithiated/sodiated defective MoS2 as shown in equation 3. A positive value in red colour signifies accumulation of charge, whereas a negative value shown in green colour denotes depletion of charge. These charge density isosurface plots after the formation of defects clearly demonstrate that increased charge transfer occurs from Li/Na towards nearby S atoms of defective MoS2 sheet which prompts to adsorption of Li/Na. To quantitatively quantify the total amount of charge transferred between Li/Na on both pristine and defective MoS2, Bader charge analysis method has been carried out. The Bader charge of isolated Li/Na atom and Bader charge of defective MoS2 is subtracted from the Bader charge of lithiated/sodiated 13 ACS Paragon Plus Environment


Page 14 of 32

defective MoS2. The number of electrons transferred from Li/Na to pristine MoS2 is 0.89 |e| for Li and 0.79 |e| for Na. The amount of electron transfer at various types of defects is shown in Table 1. The charge transfer that takes place from Li/Na to defective MoS2 suggests ionic bonding between Li/Na and defective MoS2. 3.3 Effect of Vacancies on the Li/Na Diffusion Barrier Energy As negligible diffusion barrier energy and fast mobility of ions are key factors for determining the rate performance of a material to use as an efficient anode material during charging and discharging of a battery. So, it is essential to study the diffusion of both Li and Na ions above the surface of the material. The rate execution of the electrodes in LIBs/SIBs is depicted by adsorption and migration of intercalated ions over the surface of electrodes. Before going to study Li/Na diffusion on defective MoS2, first, we explored in-plane migration of both Li and Na atoms on pristine MoS2. Migration of alkali atoms always takes place in a path connecting two adjacent binding sites. But to contemplate the impact of defects on diffusion energy and energy barriers for getting away from the defective regions, we have chosen a different path (Figure 3a) i.e. migration of Li/Na in the two-unit cell in the diagonal direction of monolayer MoS2 has been considered. As the most favourable adsorption position for the Li/Na adsorption on pristine MoS2 monolayer is top site, the diffusion properties of both Li and Na is taken in to account from one top position of one-unit cell to next nearest top site of another unit cell by passing one top site and two hollow sites in between the top sites (Figure 3a) and this path has been fixed for all types of defects studied here. The associated activation energy barriers are calculated by CI-NEB method. The minimum-energy pathways (MEPs) along with activation barrier heights at different coordinates are shown in Figure 3. Table 2. Calculated Barriers Heights for the diffusions of Li and Na over defect-free and defective MoS2. Vacancy

Barrier Heights (eV) Li

Na

Pristine MoS2

0.57

0.28

V1S

0.69

0.43

V2S

0.70

0.40

VMo

0.46

0.32

VMoS3

0.91

0.65


Page 15 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


VMoS6

0.95

0.67

MoS2

0.74

0.47

It is notable that, the rate performance of a battery relies upon the mobility of migrated particles at the anode of the battery. For pristine MoS2 models, when Li/Na crosses one-unit cell, the pathway is characterised by one symmetrical absolute maximum, one symmetrical local



Figure 3. Diffusion path for Li/Na on MoS2 with defects. Diffusion Energy profiles for Li/Na diffusion barriers for (a) Pristine MoS2, (b) Li/Na migration path on Pristine MoS2, (c) mono-sulphur vacancy (d) di-sulphur vacancy, (e) Mo vacancy, (f) MoS2 vacancy, (g) MoS3 vacancy (h) MoS6 vacancy.

maximum and one local minimum. Similarly, when Li/Na crosses two-unit cell, the migration path is portrayed by the two symmetrical absolute maxima, two symmetrical local maxima and two local minima for the diffusion barrier (Figure 3b). In pristine MoS2, to cross the barrier, Li needs only 0.57 eV of energy and Na needs 0.28 eV of energy (Table 2). The typical migration path of both Li/Na on pristine MoS2 displays relatively smaller diffusion barriers when moves from one Mo top site to the adjacent one through the hexagonal hollow site.12,63 As we have taken a different path, these values are quite higher but very close to the diffusion energy of Li/ Na ions when migrates on the surface of graphite i.e. about 0.48 eV for Li diffusion and 0.29 eV for Na diffusion.64 Presence of vacancies affects both barrier heights and minimum-energy pathways to a larger extent. In concurrence to that of pristine MoS2, the entire diffusion pathways of defective MoS2 were simulated with regard to restriction of Li/Na atoms in the peripheral of the defective region and taking into consideration their departure from the same. When a vacancy is created, we find only one maximum in the vicinity of the defect, because the barrier height at the defective region vanishes. The results well agreed with the case of graphene, in which doublevacancy, significantly reduces the diffusion barrier energy at the vacancy sites.27 The energy barrier for diffusion of Li increases to 0.69 eV and for Na, it is 0.43 eV for mono sulphur vacancy (Figure 3c). Again, for a di-sulphur vacancy, the energy barrier for diffusion of Li becomes 0.70 eV and 0.40 eV for Na (Figure 3d). Although the barrier heights in the vicinity 16 ACS Paragon Plus Environment

Page 16 of 32

Page 17 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


of the defect are quite larger than pristine MoS2, barrier height at the defective region are lower than pristine MoS2. So, it is clear that existence vacancy reduces barrier height at the defective region and Li/Na ions diffuse effectively to the vacancy region and afterwards go through the vacancy to another top site of next unit cell of MoS2 plane. However, the diffusion barriers for Mo vacancy is found to be lower than both mono- and di-sulphur vacancy and it is 0.46 eV for Li diffusion and 0.32 eV for Na diffusion (Figure 3e). This results well agreed with previous results of Li/Na migration in ReS2,38 in which Re vacancy facilitates the migration of both Li/Na ions faster than sulphide vacancy. Although the introduction of defect enhances barriers energy in the vicinity of the defect, barrier energy of around 0.70 eV will definitely assure substantial rate performance at room temperature for LIBs.65 This result has an optimistic implication i. e. the presence of vacancy can promote migration of Li/Na ions on MoS2 surface and hence enhances the performance of MoS2 in LIBs/SIBs to a larger extent. Although the results of adsorption energy were quite optimistic in case of VMoS3 and VMoS6 (Figure 3g and 3h) which are preferable for alkali atoms in anode materials, but the phenomena of the enhanced adsorption energy inhibit the atomic motion of both Li and Na ions when moving over VMoS3 and VMoS6 vacancy implicating higher barrier heights in comparison to other types of vacancies. Moreover, a huge amount of adsorption energies due to the formation of defects fundamentally perform as traps and subsequently the accompanying diffusion barrier energies increase abruptly. With a resemblance to vacancy defects, the presence of anti-site defects also facilitates the migration of alkali atoms over the surface of defective MoS2. The diffusion barrier for MoS2 vacancy are 0.74 eV for Li diffusion and 0.47 eV for Na diffusion which is very close to disulphur vacancy (Figure 3f). However, in case of S2Mo, the migration of both Li and Na ions has been stopped, because the migrated Li/Na ions suffer a repulsion from the sulphur atom substituted the Mo atom. The above simulations results suggest that in the vicinity of defects, the diffusion energy barrier is quite large in compared to the pristine case whereas it is comparatively small when escaping the defective region. Defective MoS2 strongly binds Li/Na atoms and subsequently enhances the diffusion barrier heights whereas in the defective region, due to lack of repulsions from other atoms, the barrier height decreases. Similar results have been observed for defective monolayer black phosphorus,37 but, is in contrast to monolayer ReS238 where large energy 17 ACS Paragon Plus Environment


barriers are required to escape defective sites. The barriers heights along the minimum-energy pathways for the diffusion of both Li and Na over pristine and various defective MoS2 are presented in Table2. 3.4 Electronic properties of Li/Na interaction of pristine and defective MoS2 As LIBs/SIBs have been widely used in the electronic device of applications, it is a basic requirement to know the electronic properties of electrode materials concerning their battery execution. Before going to study Li/Na interaction on defective MoS2, the electronic structures of various defects in MoS2 are studied first. Pristine MoS2 possess a direct bandgap of around 1.67 eV, (Figure S4a) consistent with previous DFT calculations66 and experimental results.67,68 Our study suggests that the existence of defect states reduce the bandgap significantly. So, defect engineering is a tool by which the electrical conductivity of semiconducting MoS2 can be enhanced. The localized electrons due to defects create mid-gap states within the valence or conduction band and which in turn results p-type or n-type conductor and the mid-gap states create free charge carriers results enhanced conductivity.69 The presence of both mono sulphur and di sulphur vacancy (V1S and V2S) creates a single peak near the centre of the bandgap towards the conduction band around 1eV (Figure S4b and S4c). The results well agreed with previous calculations.17,70 The presence of the V1S and V2S reduces the bandgap from 1.67 eV to 1 eV and as V1S and V2S are formed by missing of S atoms which hamper the coordination sphere of Mo atom. So, as Mo coordination sphere is affected the mid-gap states arises from Mo 4d atomic orbitals and the bottom of the conduction band is mostly dominated by 4d atomic orbital of Mo, presenting a strong degree of delocalization. VMo (Figure S4d) induces more than one impurity states69 whereas VMoS3 and VMoS6 induce a large no of sates (figure S4e and figure S4f). The anti-site defects also create impurity states near the Fermi level (figure S4f and figure S4g). To analyse the binding interaction between Li/Na and defective MoS2, we studied the electronic properties of both pristine and defective MoS2. To recognize the influence of various types of defects, the DOS of defective MoS2 and lithiated/sodiated defective MoS2 are calculated (Figure 4). The PDOS of Li/Na-adsorbed MoS2 shows that the s orbital of Li/Na interacts with the p orbitals of defective MoS2. These interactions between lithiated/sodiated MoS2 was confirmed by Bader’s charge transfer as discussed before. The enhanced adsorption in 18 ACS Paragon Plus Environment

Page 18 of 32

Page 19 of 32


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60



Figure 4. Spin-up and spin-down Density of state plots for Li/Na adsorbed on various defective monolayer MoS2. The spin up DOS is plotted with thicker lines and spin-down with thinner lines.

defective MoS2 is due to the strong interaction between the localized states located at the mid of bandgap of Mo d-orbital with the s orbital of Li/Na to form strong hybridisation. Another important property that is observed during Li/Na adsorption in defective MoS2 is structural rearrangement. Pristine MoS2 exists in two different polytypic structures depending on the rearrangement sulphur atoms around the Mo atom. One polytype is based on trigonal symmetry, where the metal atoms are sandwiched between two atomic layers of sulphur atoms in a trigonal prismatic geometry known as 2H MoS2. Another polytype is based on the metal atoms octahedrally located in the environment of the sulphur atoms, where atomic layers of sulphur atoms are located in an octahedral geometry which known as T phase.71 The 20 ACS Paragon Plus Environment

Page 20 of 32

Page 21 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


electronic properties of these two phases are completely different. The stable 2H phase shows semiconductor character, whereas the metastable 1T phase of MoS2 is metallic in nature.72 It has been reported that intercalating alkali metals in the stable 2H phase of MoS2 lead conversion of 2H to 1T phase.73 The structural transformation of 2H-MoS2 using electrochemical Li+ intercalation is due to the change in the symmetry of the MoS2 which favours in the formation of metastable 1T phase.74 Xia et. al75 reported that the Li+ supported 1T lattice phase can be spontaneously distorted to a more stable phase (dT), with a distinct crystal structure and maintains Weyl state property. Our band gap calculations show that almost all defective MoS2 are semiconducting in nature with a reduced gap due to mid-gap states, (Table S1) but all lithiated/sodiated defective MoS2 are metallic in nature, which indicates that semiconducting to metallic phase transition occurs after lithiation and sodiation. A semiconducting to metallic phase transition indicates that defective MoS2 also suffers structural rearrangement of atoms after alkali intercalation as like pristine MoS2. We have also studied the magnetic properties of all defective MoS2. As the 2H phase of MoS2 has a trigonal prismatic structure, the Mo4+ ions have two 4d electrons. According to Crystal Field Theory (CFT), these two electrons are located in d z 2 orbital with antiparallel spin76 which results in a magnetic moment of zero.77 But when vacancies are formed, magnetism is induced (Table S1). In the case of VMoS6 and MoS2, a magnetic moment of 2 and 2.25 μB respectively is observed. Because in these two vacancies, the oxidation state is close to +3, (from Mo4+ to Mo3+) in which the number of 4d electrons becomes 3 and one unpaired electron goes to the d x 2  y 2 (according to Crystal Field Theory) orbital for which VMoS6 and MoS2 display magnetic moment. However, effect of magnetism does not have any impact on adsorption and diffusion properties of Li/Na. The adsorption properties of Li/Na with and without magnetism calculations are listed in Table S2. 3.5 Theoretical specific capacity and Open-circuit voltage Since energy density is a critical factor that determines the performance of a material in LIBs/SIBs and energy density directly depends on capacity, which in turn is directly proportional to the amount of charge stored in the electrode material, So, the advancement of advanced LIBs/SIBs with high power density to produce high capacity is the only quest to 21 ACS Paragon Plus Environment


accommodate the increased requirements of powerful electronic devices.78,79 For decades, graphite has been utilized as anodes for LIBs/SIBs due to its wide natural abundance and low cost. However, it’s extremely low theoretical specific capacity of 372 mAhg-1 cannot fully meet the requirements of modern electronic devices.80 Therefore, creating substitute anode for graphite with enhanced capacity is highly needed for fulfilling the higher demand for modern electronic devices.80,81 The arrival of transition metal dichalcogenide MoS2 produces high specific capacity due to its large interlayer distance in between layers which can easily accommodate four alkali metal ions per MoS2 unit cell that enables them to produce a large specific capacity of ∼ 670 mAhg−1 without causing a significant volume expansion.59 Although MoS2 exhibits high Li/Na storage capacities (with 4 mol of Li+/Na+ insertion per formula),82,83 the specific capacity of MoS2 is found to increase further through the formation of defects. The defective monolayer MoS2 is advantageous in specific capacity in comparison with the corresponding pristine MoS2. The specific capacity of defective MoS2 is calculated by equation 7. According to equation 7, the theoretical specific capacity depends on the molecular weight of both Mo and S atom and hence decrease of the molecular weight due to vacancy, will enhance the specific capacity. As like previous calculations, the specific capacity of pristine MoS2 is found to be 670 mAhg−1, when a 2/2 supercell of MoS2 is loaded by 16 Li/Na atoms, i.e. 4 atoms per unit cell. The specific capacity is found to increase to 705 mAhg-1 and 745 mAhg-1 when single and double sulphur vacancies are created. Similarly, Mo vacancy in a 2/2 supercell enhances the specific capacity to 788 mAhg-1.


Page 22 of 32

Page 23 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Figure 5: The calculated voltage profile for pristine and defective MoS2 at the different theoretical specific capacity of Li/Na.

The open-circuit voltage (OCV) being a significant characteristic of an electrode material reflects the performance of a battery and is closely associated with the total energy storage of battery. The OCV can be estimated by using equation 4 and OCV is related to formation energy which is calculated by equation 6.48 So, to study OCV, it is essential to study

the formation energy with respect to Li/Na concentration. For different concentrations of Li/Na adsorbed on MoS2, the formation energy for both lithium and sodium are found to be negative. Commonly, the OCV for anodes in LIBs/SIBs should be low in order to get a maximum voltage.80 Moreover, the OCV value within 0.1 V to 1.0 V suppress the growth of dendrites during the lithiation/sodiation process.84 As shown in Figure 5, It is noted that for all of the cases, the OCV retains a positive value within the range of 0.1-1.2 V. These moderate values of OCVs indicates that high energy density for LIBs/SIBs can be achieved through the formation of defects in MoS2. The lack of negative voltage signifies the Li/Na ions prefer to adsorb on defective MoS2 instead of forming dendrite growth. As the theoretical capacity of defective MoS2 monolayer corresponds to maximum Li/Na concentration, the calculated OCV for pristine MoS2 at maximum concentration is 0.28 V for LIBs and 0.12 V for SIBs. At this stage the lattice constant increases to 3.68% for lithiation and 5.39% for sodiation. The OCV of defective systems is found to be slightly higher than that of pristine MoS2. For example, the OCV at maximum concentration for the mono-sulphur vacancy is 0.33 V for LIBs and 0.25 V for SIBs and the lattice constant enlarges to 3.27% and 4.91% for Li and Na respectively. Similarly, for di-sulphur vacancy, the OCV is 0.33 V for LIBs and 0.14 V for SIBs with an increase of lattice constant of 3.02% for Li and 4.70% for Na. The OCV at maximum concentration for Mo vacancy is 0.64 V for LIBs and 0.36 V for SIBs and the lattice constant increases to 4.68% and 5.78% for Li and Na respectively. The OCV at different specific capacity is shown in Figure 5. Finally, by considering all properties like adsorption energies, migration barriers, storage capacity and OCV, we concluded that VMo is the best type of vacancy for the storage of both Li/Na. However, if we compare between Li and Na, the formation of vacancies is more advantageous for Na storage over Li. Formation of vacancies enhances adsorption energies of both Li/Na to a larger extent, but Li adsorption energies are extremely high in comparison to Na adsorption energy for which Li diffusion barriers are higher than Na diffusion Barrier. So, 23 ACS Paragon Plus Environment


Na storage is more preferable on defective MoS2 as the adsorption energies are higher than pristine MoS2 with negligible barrier energy and high storage capacity. The extremely high adsorption energy for Li is due to the fact that, due to vacancies dangling bond are formed and the unsatisfied valency due to dangling bonds prefer to adsorb adatoms to satisfy its valency. The reason for high Li adsorption energy can be explained on the basis of atomic radii. The atomic radius of Li is 150 pm and for Na, it is 227 pm. The dangling bonds are formed due to S vacancy and atomic radius of S is 100 pm close to the atomic radius of Li. Out of Li and Na, it attracts Li more, and Li adsorption energies are higher with a high barrier. So, Na storage is preferable over Li storage. Conclusions In this work, we have explored the presence of different types of defects that exist in MoS2 in improving the performance of MoS2 in LIBs/SIBs by means of periodic DFT calculations. The Li/Na adsorption energy, migration barrier, theoretical storage capacity, open-circuit voltage as well as electronic properties of pristine and defective MoS2 are systematically investigated. The calculated defect formation energies indicate that S vacancies are highly favourable for formation. Electronic analysis indicates that the existence of vacancies in MoS2 plays a crucial role in modifying the electronic energy levels in MoS2, which displays a convenient strategy to tune the performance of MoS2 in LIBs/SIBs. The existence of vacancies generates mid-gap states in between VBM and CBM which interact strongly with the incoming alkali metal atoms and hence enhances the adsorption energy to a larger extent. Energy barriers in the vicinity of the defect are quite larger than those of pristine MoS2, whereas for escaping defective sites these levels are comparatively small. In addition, we established that the introduction of defects improves the specific capacity of the defective MoS2 to a larger extent than the pristine MoS2 both for LIBs and SIBs. On the other hand, the lithiation/sodiation opencircuit voltage is not affected by enhanced adsorption and the OCV values of defective MoS2 is found to be in the range of reported anode materials which indicates potential application as high-performance anodes of LIBs/SIBs. Among all types of vacancy, VMo is the best vacancy for storages of both Li/Na, However, our study also reveals that formation of vacancy in MoS2 favours replacement of LIBs by SIBS which is a lower-cost alternative with higher natural abundance as well as risk less. Our findings may provide an effective opportunity to advance Li/Na storage technique of MoS2 in LIBs/SIBs, with structural defects being contemplated. 24 ACS Paragon Plus Environment

Page 24 of 32

Page 25 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Supporting Information Optimized geometry and charge density difference iso-surfaces of Li/Na adsorbed at defective sites of monolayer MoS2 with various types of defects. Band structure diagrams of various types of defects in monolayer MoS2. The calculated vacancy formation energy (eV), Optimized lattice constant (Å), Energy gap (eV) and magnetic moment of various defective and pristine MoS2. Effect of magnetism on adsorption energy of Li/Na is given. Author Information Corresponding Author * Email: [email protected] Notes The authors declare no competing financial interest. Acknowledgements S.P. acknowledges the J.C. Bose Grant received from Dept. Of Science and Technology, Govt. Of India, New Delhi, India. G.B. acknowledges Council of Scientific and Industrial Research, New Delhi, India for providing Senior Research Fellowship. The authors are also thankful to High-Performance Computing facilities of IIT Bombay, Mumbai, India. References: (1)

Megahed, S.; Scrosati, B. Lithium-Ion Rechargeable Batteries. J. Power Sources 1994, 51, 79–104.

(2)

Goodenough, J. B. Energy Storage Materials: A Perspective. Energy Storage Mater. 2015, 1, 158–161.

(3)

Palomares, V.; Serras, P.; Villaluenga, I.; Hueso, K. B.; Carretero-González, J.; Rojo, T. Na-Ion Batteries, Recent Advances and Present Challenges to Become Low Cost Energy Storage Systems. Energy Environ. Sci. 2012, 5, 5884–5901.

(4)

Ong, S. P.; Chevrier, V. L.; Hautier, G.; Jain, A.; Moore, C.; Kim, S.; Ma, X.; Ceder, G. Voltage, Stability and Diffusion Barrier Differences between Sodium-Ion and Lithium-Ion Intercalation Materials. Energy Environ. Sci. 2011, 4, 3680–3688.

(5)

Kundu, D.; Talaie, E.; Duffort, V.; Nazar, L. F. The Emerging Chemistry of Sodium Ion Batteries for Electrochemical Energy Storage. Angew. Chemie Int. Ed. 2015, 54, 3431–3448.

(6)

Yoffe, A. D. The Transition Metal Dichalcogenides Discussion and Interpretation of 25 ACS Paragon Plus Environment


the Observed Optical, Electrical and Structural Properties AU - Wilson, J.A. Adv. Phys. 1969, 18, 193–335. (7)

Croguennec, L.; Palacin, M. R. Recent Achievements on Inorganic Electrode Materials for Lithium-Ion Batteries. J. Am. Chem. Soc. 2015, 137, 3140–3156.

(8)

Karmakar, S.; Chowdhury, C.; Datta, A. Two-Dimensional Group IV Monochalcogenides: Anode Materials for Li-Ion Batteries. J. Phys. Chem. C 2016, 120, 14522–14530.

(9)

Xiong, F.; Wang, H.; Liu, X.; Sun, J.; Brongersma, M.; Pop, E.; Cui, Y. Li Intercalation in MoS2: In Situ Observation of Its Dynamics and Tuning Optical and Electrical Properties. Nano Lett. 2015, 15, 6777–6784.

(10)

Gao, P.; Wang, L.; Zhang, Y.; Huang, Y.; Liu, K. Atomic-Scale Probing of the Dynamics of Sodium Transport and Intercalation-Induced Phase Transformations in MoS2. ACS Nano 2015, 9, 11296–11301.

(11)

Mortazavi, M.; Wang, C.; Deng, J.; Shenoy, V. B.; Medhekar, N. V. Ab Initio Characterization of Layered MoS2 as Anode for Sodium-Ion Batteries. J. Power Sources 2014, 268, 279–286.

(12)

Li, Y.; Wu, D.; Zhou, Z.; Cabrera, C. R.; Chen, Z. Enhanced Li Adsorption and Diffusion on MoS2 Zigzag Nanoribbons by Edge Effects: A Computational Study. J. Phys. Chem. Lett. 2012, 3, 2221–2227.

(13)

Zhou, W.; Zou, X.; Najmaei, S.; Liu, Z.; Shi, Y.; Kong, J.; Lou, J.; Ajayan, P. M.; Yakobson, B. I.; Idrobo, J.-C. Intrinsic Structural Defects in Monolayer Molybdenum Disulfide. Nano Lett. 2013, 13, 2615–2622.

(14)

Hong, J.; Hu, Z.; Probert, M.; Li, K.; Lv, D.; Yang, X.; Gu, L.; Mao, N.; Feng, Q.; Xie, L.; et al. Exploring Atomic Defects in Molybdenum Disulphide Monolayers. Nat. Commun. 2015, 6, 6293.

(15)

KC, S.; Longo, R. C.; Addou, R.; Wallace, R. M.; Cho, K. Impact of Intrinsic Atomic Defects on the Electronic Structure of MoS2 monolayers. Nanotechnology 2014, 25, 375703.

(16)

Komsa, H.-P.; Krasheninnikov, A. V. Native Defects in Bulk and Monolayer MoS2 from First Principles. Phys. Rev. B 2015, 91, 125304.

(17)

Noh, J.-Y.; Kim, H.; Kim, Y.-S. Stability and Electronic Structures of Native Defects in Single-Layer MoS2. Phys. Rev. B 2014, 89, 205417.

(18)

Najmaei, S.; Yuan, J.; Zhang, J.; Ajayan, P.; Lou, J. Synthesis and Defect Investigation of Two-Dimensional Molybdenum Disulfide Atomic Layers. Acc. Chem. Res. 2015, 26 ACS Paragon Plus Environment

Page 26 of 32

Page 27 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


48, 31–40. (19)

Liu, X.; Balla, I.; Bergeron, H.; Hersam, M. C. Point Defects and Grain Boundaries in Rotationally Commensurate MoS2 on Epitaxial Graphene. J. Phys. Chem. C 2016, 120, 20798–20805.

(20)

Huang, Y. L.; Chen, Y.; Zhang, W.; Quek, S. Y.; Chen, C.-H.; Li, L.-J.; Hsu, W.-T.; Chang, W.-H.; Zheng, Y. J.; Chen, W.; et al. Bandgap Tunability at Single-Layer Molybdenum Disulphide Grain Boundaries. Nat. Commun. 2015, 6, 6298.

(21)

Wang, K.; Wang, J.; Fan, J.; Lotya, M.; O’Neill, A.; Fox, D.; Feng, Y.; Zhang, X.; Jiang, B.; Zhao, Q.; et al. Ultrafast Saturable Absorption of Two-Dimensional MoS2 Nanosheets. ACS Nano 2013, 7, 9260–9267.

(22)

Zhang, Z.; Zou, X.; Crespi, V. H.; Yakobson, B. I. Intrinsic Magnetism of Grain Boundaries in Two-Dimensional Metal Dichalcogenides. ACS Nano 2013, 7 , 10475– 10481.

(23)

González, C.; Dappe, Y. J.; Biel, B. Reactivity Enhancement and Fingerprints of Point Defects on a MoS2 Monolayer Assessed by Ab Initio Atomic Force Microscopy. J. Phys. Chem. C 2016, 120, 17115–17126.

(24)

Seebauer, E. G.; Noh, K. W. Trends in Semiconductor Defect Engineering at the Nanoscale. Mater. Sci. Eng. R Reports 2010, 70, 151–168.

(25)

Ye, G.; Gong, Y.; Lin, J.; Li, B.; He, Y.; Pantelides, S. T.; Zhou, W.; Vajtai, R.; Ajayan, P. M. Defects Engineered Monolayer MoS2 for Improved Hydrogen Evolution Reaction. Nano Lett. 2016, 16, 1097–1103.

(26)

Liu, G.; Robertson, A. W.; Li, M. M.-J.; Kuo, W. C. H.; Darby, M. T.; Muhieddine, M. H.; Lin, Y.-C.; Suenaga, K.; Stamatakis, M.; Warner, J. H.; et al. MoS2 Monolayer Catalyst Doped with Isolated Co Atoms for the Hydrodeoxygenation Reaction. Nat. Chem. 2017, 9, 810.

(27)

Fan, X.; Zheng, W. T.; Kuo, J.-L. Adsorption and Diffusion of Li on Pristine and Defective Graphene. ACS Appl. Mater. Interfaces 2012, 4, 2432–2438.

(28)

Er, D.; Detsi, E.; Kumar, H.; Shenoy, V. B. Defective Graphene and Graphene Allotropes as High-Capacity Anode Materials for Mg Ion Batteries. ACS Energy Lett. 2016, 1, 638–645.

(29)

Addou, R.; Colombo, L.; Wallace, R. M. Surface Defects on Natural MoS2. ACS Appl. Mater. Interfaces 2015, 7, 11921–11929.

(30)

Bazant, M. Z. Theory of Chemical Kinetics and Charge Transfer Based on Nonequilibrium Thermodynamics. Acc. Chem. Res. 2013, 46, 1144–1160. 27 ACS Paragon Plus Environment


(31)

Puntambekar, A.; Chandrasekeran, N.; Wang, Q.; Roy, I.; Premkumar, V.; Chakrapani, V. Interaction of Native Defects with Ions and Its Role in Inducing Phase Transitions in P-Type S-Excess MoS2. ACS Appl. Energy Mater. 2018, 1, 3093–3102.

(32)

Zhao, X.; Hayner, C. M.; Kung, M. C.; Kung, H. H. Flexible Holey Graphene Paper Electrodes with Enhanced Rate Capability for Energy Storage Applications. ACS Nano 2011, 5, 8739–8749.

(33)

Zhao, X.; Hayner, C. M.; Kung, M. C.; Kung, H. H. In-Plane Vacancy-Enabled HighPower Si–Graphene Composite Electrode for Lithium-Ion Batteries. Adv. Energy Mater. 2011, 1, 1079–1084.

(34)

Carlsson, J. M.; Scheffler, M. Structural, Electronic, and Chemical Properties of Nanoporous Carbon. Phys. Rev. Lett. 2006, 96, 46806.

(35)

Wan, Q.; Li, S.; Liu, J.-B. First-Principle Study of Li-Ion Storage of Functionalized Ti2C Monolayer with Vacancies. ACS Appl. Mater. Interfaces 2018, 10, 6369–6377.

(36)

Sun, X.; Wang, Z.; Fu, Y. Q. Defect-Mediated Lithium Adsorption and Diffusion on Monolayer Molybdenum Disulfide. Sci. Rep. 2015, 5, 18712.

(37)

Zhang, R.; Wu, X.; Yang, J. Blockage of Ultrafast and Directional Diffusion of Li Atoms on Phosphorene with Intrinsic Defects. Nanoscale 2016, 8, 4001–4006.

(38)

Mukherjee, S.; Banwait, A.; Grixti, S.; Koratkar, N.; Singh, C. V. Adsorption and Diffusion of Lithium and Sodium on Defective Rhenium Disulfide: A First Principles Study. ACS Appl. Mater. Interfaces 2018, 10, 5373–5384.

(39)

Datta, D.; Li, J.; Shenoy, V. B. Defective Graphene as a High-Capacity Anode Material for Na- and Ca-Ion Batteries. ACS Appl. Mater. Interfaces 2014, 6, 1788– 1795.

(40)

Wentzcovitch, P. G. and S. B. and N. B. and M. C. and R. C. and C. C. and D. C. and G. L. C. and M. C. and I. D. and A. D. C. and S. de G. and S. F. QUANTUM ESPRESSO: A Modular and Open-Source Software Project for Quantum Simulations of Materials. J. Phys. Condens. Matter 2009, 21, 395502.

(41)

Perdew, J. P.; Burke, K.; Ernzerhof, M. Generalized Gradient Approximation Made Simple. Phys. Rev. Lett. 1996, 77, 3865–3868.

(42)

Kresse, G.; Joubert, D. From Ultrasoft Pseudopotentials to the Projector AugmentedWave Method. Phys. Rev. B 1999, 59, 1758–1775.

(43)

Monkhorst, H. J.; Pack, J. D. Special Points for Brillouin-Zone Integrations. Phys. Rev. B 1976, 13, 5188–5192.

(44)

Grimme, S.; Antony, J.; Ehrlich, S.; Krieg, H. A Consistent and Accurate Ab Initio 28 ACS Paragon Plus Environment

Page 28 of 32

Page 29 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Parametrization of Density Functional Dispersion Correction (DFT-D) for the 94 Elements H-Pu. J. Chem. Phys. 2010, 132, 154104. (45)

Henkelman, G.; Arnaldsson, A.; Jónsson, H. A Fast and Robust Algorithm for Bader Decomposition of Charge Density. Comput. Mater. Sci. 2006, 36, 354–360.

(46)

Henkelman, G.; Uberuaga, B. P.; Jónsson, H. A Climbing Image Nudged Elastic Band Method for Finding Saddle Points and Minimum Energy Paths. J. Chem. Phys. 2000, 113, 9901–9904.

(47)

Ataca, C.; Şahin, H.; Aktürk, E.; Ciraci, S. Mechanical and Electronic Properties of MoS2 Nanoribbons and Their Defects. J. Phys. Chem. C 2011, 115, 3934–3941.

(48)

Courtney, I. A.; Tse, J. S.; Mao, O.; Hafner, J.; Dahn, J. R. Ab Initio Calculation of the Lithium-Tin Voltage Profile. Phys. Rev. B 1998, 58, 15583–15588.

(49)

Rettig, S. J.; Trotter, J. Refinement of the Structure of Orthorhombic Sulfur, α-S8. Acta Crystallogr. Sect. C 1987, 43, 2260–2262.

(50)

Radisavljevic, B.; Radenovic, A.; Brivio, J.; Giacometti, V.; Kis, A. Single-Layer MoS2 Transistors. Nat. Nanotechnol. 2011, 6, 147.

(51)

Johari, P.; Shenoy, V. B. Tuning the Electronic Properties of Semiconducting Transition Metal Dichalcogenides by Applying Mechanical Strains. ACS Nano 2012, 6, 5449–5456.

(52)

Komsa, H.-P.; Kurasch, S.; Lehtinen, O.; Kaiser, U.; Krasheninnikov, A. V. From Point to Extended Defects in Two-Dimensional MoS2: Evolution of Atomic Structure under Electron Irradiation. Phys. Rev. B 2013, 88, 35301.

(53)

Qiu, H.; Xu, T.; Wang, Z.; Ren, W.; Nan, H.; Ni, Z.; Chen, Q.; Yuan, S.; Miao, F.; Song, F.; et al. Hopping Transport through Defect-Induced Localized States in Molybdenum Disulphide. Nat. Commun. 2013, 4, 2642.

(54)

Haldar, S.; Vovusha, H.; Yadav, M. K.; Eriksson, O.; Sanyal, B. Systematic Study of Structural, Electronic, and Optical Properties of Atomic-Scale Defects in the TwoDimensional Transition Metal Dichalcogenides MX2 (M = Mo, W; X = S, Se, Te). Phys. Rev. B 2015, 92, 235408.

(55)

Yuan, S.; Roldán, R.; Katsnelson, M. I.; Guinea, F. Effect of Point Defects on the Optical and Transport Properties of MoS2 and WS2. Phys. Rev. B 2014, 90, 41402.

(56)

Banhart, F.; Kotakoski, J.; Krasheninnikov, A. V. Structural Defects in Graphene. ACS Nano 2011, 5, 26–41.

(57)

Lee, G.-D.; Wang, C. Z.; Yoon, E.; Hwang, N.-M.; Kim, D.-Y.; Ho, K. M. Diffusion, Coalescence, and Reconstruction of Vacancy Defects in Graphene Layers. Phys. Rev. 29 ACS Paragon Plus Environment


Lett. 2005, 95, 205501. (58)

Krasheninnikov, A. V; Lehtinen, P. O.; Foster, A. S.; Nieminen, R. M. Bending the Rules: Contrasting Vacancy Energetics and Migration in Graphite and Carbon Nanotubes. Chem. Phys. Lett. 2006, 418, 132–136.

(59)

Hwang, H.; Kim, H.; Cho, J. MoS2 Nanoplates Consisting of Disordered Graphenelike Layers for High Rate Lithium Battery Anode Materials. Nano Lett. 2011, 11, 4826–4830.

(60)

Lacey, S. D.; Wan, J.; Cresce, A. von W.; Russell, S. M.; Dai, J.; Bao, W.; Xu, K.; Hu, L. Atomic Force Microscopy Studies on Molybdenum Disulfide Flakes as Sodium-Ion Anodes. Nano Lett. 2015, 15, 1018–1024.

(61)

Datta, D.; Li, J.; Koratkar, N.; Shenoy, V. B. Enhanced Lithiation in Defective Graphene. Carbon N. Y. 2014, 80, 305–310.

(62)

Sun, X.; Wang, Z.; Fu, Y. Q. Adsorption and Diffusion of Sodium on Graphene with Grain Boundaries. Carbon N. Y. 2017, 116, 415–421.

(63)

Fan, S.; Zou, X.; Du, H.; Gan, L.; Xu, C.; Lv, W.; He, Y.-B.; Yang, Q.-H.; Kang, F.; Li, J. Theoretical Investigation of the Intercalation Chemistry of Lithium/Sodium Ions in Transition Metal Dichalcogenides. J. Phys. Chem. C 2017, 121, 13599–13605.

(64)

Wang, Z.; Ratvik, A. P.; Grande, T.; Selbach, S. M. Diffusion of Alkali Metals in the First Stage Graphite Intercalation Compounds by VdW-DFT Calculations. RSC Adv. 2015, 5, 15985–15992.

(65)

Ceder, G.; Hautier, G.; Jain, A.; Ong, S. P. Recharging Lithium Battery Research with First-Principles Methods. MRS Bull. 2011, 36, 185–191.

(66)

Kadantsev, E. S.; Hawrylak, P. Electronic Structure of a Single MoS2 Monolayer. Solid State Commun. 2012, 152, 909–913.

(67)

Mak, K. F.; Lee, C.; Hone, J.; Shan, J.; Heinz, T. F. Atomically Thin MoS2: A New Direct-Gap Semiconductor. Phys. Rev. Lett. 2010, 105, 136805.

(68)

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.

(69)

Kunstmann, J.; Wendumu, T. B.; Seifert, G. Localized Defect States in MoS2 Monolayers: Electronic and Optical Properties. Phys. status solidi 2017, 254, 1600645.

(70)

Pandey, M.; Rasmussen, F. A.; Kuhar, K.; Olsen, T.; Jacobsen, K. W.; Thygesen, K. S. Defect-Tolerant Monolayer Transition Metal Dichalcogenides. Nano Lett. 2016, 16, 2234–2239.

(71)

Mattheiss, L. F. Band Structures of Transition-Metal-Dichalcogenide Layer 30 ACS Paragon Plus Environment

Page 30 of 32

Page 31 of 32 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Compounds. Phys. Rev. B 1973, 8, 3719–3740. (72)

Cheng, Y.; Nie, A.; Zhang, Q.; Gan, L.-Y.; Shahbazian-Yassar, R.; Schwingenschlogl, U. Origin of the Phase Transition in Lithiated Molybdenum Disulfide. ACS Nano 2014, 8, 11447–11453.

(73)

Somoano, R. B.; Hadek, V.; Rembaum, A. Alkali Metal Intercalates of Molybdenum Disulfide. J. Chem. Phys. 1973, 58, 697–701.

(74)

Sun, X.; Wang, Z.; Li, Z.; Fu, Y. Q. Origin of Structural Transformation in Mono- and Bi-Layered Molybdenum Disulfide. Sci. Rep. 2016, 6, 26666.

(75)

Xia, J.; Wang, J.; Chao, D.; Chen, Z.; Liu, Z.; Kuo, J.-L.; Yan, J.; Shen, Z. X. Phase Evolution of Lithium Intercalation Dynamics in 2H-MoS2. Nanoscale 2017, 9, 7533– 7540.

(76)

Cai, L.; He, J.; Liu, Q.; Yao, T.; Chen, L.; Yan, W.; Hu, F.; Jiang, Y.; Zhao, Y.; Hu, T.; et al. Vacancy-Induced Ferromagnetism of MoS2 Nanosheets. J. Am. Chem. Soc. 2015, 137, 2622–2627.

(77)

Yang, H.; Kim, S. W.; Chhowalla, M.; Lee, Y. H. Structural and Quantum-State Phase Transitions in van Der Waals Layered Materials. Nat. Phys. 2017, 13, 931.

(78)

Scrosati, B.; Hassoun, J.; Sun, Y.-K. Lithium-Ion Batteries. A Look into the Future. Energy Environ. Sci. 2011, 4, 3287–3295.

(79)

Amine, K.; Belharouak, I.; Chen, Z.; Tran, T.; Yumoto, H.; Ota, N.; Myung, S.-T.; Sun, Y.-K. Nanostructured Anode Material for High-Power Battery System in Electric Vehicles. Adv. Mater. 2010, 22, 3052–3057.

(80)

Goodenough, J. B.; Kim, Y. Challenges for Rechargeable Li Batteries. Chem. Mater. 2010, 22, 587–603.

(81)

Manthiram, A. Materials Challenges and Opportunities of Lithium Ion Batteries. J. Phys. Chem. Lett. 2011, 2, 176–184.

(82)

Cao, X.; Shi, Y.; Shi, W.; Rui, X.; Yan, Q.; Kong, J.; Zhang, H. Preparation of MoS2Coated Three-Dimensional Graphene Networks for High-Performance Anode Material in Lithium-Ion Batteries. Small 2013, 9, 3433–3438.

(83)

Ding, S.; Zhang, D.; Chen, J. S.; Lou, X. W. (David). Facile Synthesis of Hierarchical MoS2 Microspheres Composed of Few-Layered Nanosheets and Their Lithium Storage Properties. Nanoscale 2012, 4, 95–98.

(84)

Koudriachova, M. V; Harrison, N. M.; de Leeuw, S. W. Open Circuit Voltage Profile for Li-Intercalation in Rutile and Anatase from First Principles. Solid State Ion. 2002, 152–153, 189–194. 31 ACS Paragon Plus Environment


TOC Graphic


Page 32 of 32

Defect Induced Performance Enhancement of Monolayer MoS2 for Li

Recommend Documents