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. 2024 Jul 10;16(29):37994–38005. doi: 10.1021/acsami.4c05810

MoS2/Mayenite Electride Hybrid as a Cathode Host for Suppressing Polysulfide Shuttling and Promoting Kinetics in Lithium–Sulfur Batteries

Niphat Thatsami †,‡, Parinya Tangpakonsab †,‡, Pornsawan Sikam §,∥, Tanveer Hussain ⊥, Orapa Tamwattana †,‡, Anucha Watcharapasorn ∥,#,, Pairot Moontragoon †,‡, Biswarup Pathak ∇, Thanayut Kaewmaraya †,‡,*
PMCID: PMC11295124  PMID: 38985897

Abstract

graphic file with name am4c05810_0008.jpg

The commercial viability of emerging lithium–sulfur batteries (LSBs) remains greatly hindered by short lifespans caused by electrically insulating sulfur, lithium polysulfides (Li2Sn; 1 ≤ n ≤ 8) shuttling, and sluggish sulfur reduction reactions (SRRs). This work proposes the utilization of a hybrid composed of sulfiphilic MoS2 and mayenite electride (C12A7:e–) as a cathode host to address these challenges. Specifically, abundant cement-based C12A7:e– is the most stable inorganic electride, possessing the ultimate electrical conductivity and low work function. Through density functional theory simulations, the key aspects of the MoS2/C12A7:e– hybrid including electronic properties, interfacial binding with Li2Sn, Li+ diffusion, and SRR have been unraveled. Our findings reveal the rational rules for MoS2 as an efficient cathode host by enhancing its mutual electrical conductivity and surface polarity via MoS2/C12A7:e–. The improved electrical conductivity of MoS2 is attributed to the electron donation from C12A7:e– to MoS2, yielding a semiconductor-to-metal transition. The resultant band positions of MoS2/C12A7:e– are well matched with those of conventional current-collecting materials (i.e., Cu and Ni), electrochemically enhancing the electronic transport. The accepted charge also intensifies MoS2 surface polarity for attracting polar Li2Sn by forming stronger bonds with Li2Sn via ionic Li–S bonds than electrolytes with Li2Sn, thereby preventing polysulfide shuttling. Importantly, MoS2/C12A7:e– not only promotes rapid reaction kinetics by reducing ionic diffusion barriers but also lowers the Gibbs free energies of the SRR for effective S8-to-Li2S conversion. Beyond the reported applications of C12A7:e–, this work highlights its functionality as an electrode material to boost the efficiency of LSBs.

Keywords: Li–S batteries, polysulfide shuttling, C12A7 electride, density functional theory, sulfur reduction reaction

Introduction

Lithium-ion batteries (LIBs) are currently the mainstream storage devices for portable electronic gadgets and electric vehicles (EVs).1 Nevertheless, the limited theoretical capacity (<387 Wh kg–1) of LIBs elusively suffices overwhelming demands for ultrahigh energy densities toward extended-range EVs, smart grid applications, and ambitious electric aviation.2,3 In this regard, rechargeable lithium–sulfur batteries (LSBs) have emerged as promising alternatives because of their ultimate theoretical energy density of 2600 Wh kg–1, around five times greater than LIBs.4 Moreover, the use of sulfur (i.e., the 10th most earth-abundant element) as an electrode material not only enhances sustainability due to its abundance but also leverages its status as a byproduct of the petroleum industry.1 These technical advantages position LSBs as leading contenders among next-generation battery technologies.5

Despite their numerous benefits, the technological readiness level of LSBs is around 7–8, meaning full commercialization is not available at the moment.6 LSBs typically face manifold challenges associated with the intrinsic properties of the cathode, electrolyte, and anode.5,7 The sulfur cathode is both electronic (σe = 5 × 10–15 S m–1) and ionic insulator, which retards electrochemical kinetics.8 The conversion of elemental sulfur (S8) into end-product lithium sulfide (Li2S) during the sulfur reduction reaction (SRR) involves the subtle 16-electron process (i.e., S8 + 16Li+ + 16e– = 8Li2S), yielding various heteropolar lithium polysulfides (Li2Sn; 1 ≤ n ≤ 8).9 In particular, the transitions among Li2Sn are energetically hindered by the sluggish SRR derived from the solid–liquid–solid phase evolution. Additionally, certain high-order polysulfides (Li2Sn; 4 ≤ n ≤ 8) can be readily dissolved in polar liquid electrolytes, such as common 1,2-dimethoxyethane (DME) and 1,3-dioxolane (DOL).10 This intricate solubility phenomenon, known as polysulfide shuttling, gradually depletes the active polysulfides, which are subsequently migrated backward to deposit as dendrites at the anode. Consequently, an electrically insulating solid–electrolyte interface is formed at the anode.10 Thus, the battery’s capacity retention and Coulombic efficiency are severely suppressed.11 Furthermore, the end-product Li2S of SRR induces substantial volumetric expansion of around 80.0% because of the pronounced difference in volumetric densities between S8 and Li2S.5 This results in cathode mechanical instability and shortened battery lifespan.5 Addressing these challenges is essential for enhancing the performance of LSBs.

One feasible route to overcoming these technical challenges is through the strategic utilization of the cathode hosting material, also known as the cathode anchoring material.12 The host should possess certain desirable features including high electrical conductivity, a sulfiphilic surface, rich surface area, and superior mechanical flexibility.12 High electrical conductivity is essential for efficient electrochemical reactions.5 The sulfiphilic characteristics facilitate the anchoring of mobile redox centers to directly use chemically adsorbed polysulfides, which bind to the host stronger than polysulfides with electrolytes (i.e., inhibiting the polysulfide solubility). This accelerates Li2S chemical kinetics and homogeneous growth.13 The ultrahigh surface area maximizes polysulfide chemisorption possibilities, thereby improving sulfur utilization.13 Furthermore, exceptional mechanical flexibility is crucial for tolerating cathode volumetric expansion/contraction during lithiation/delithiation.5 Given these requirements, generic two-dimensional (2D) materials become the promising candidates due to their high electrical conductivity, planar geometry with enriched surface area and chemistry, and superior mechanical flexibility.14 Accordingly, numerous works have reported the successful implementations of 2D hosting materials for enhancing the efficiencies of LSBs such as graphene,15 S-terminated Ti2C MXenes,16 biphenylene,17 carbon nitrides CxNy,18 germanene,19 antimonene,20 and generic transition metal dichalcogenides (TMDs).21 Their major accomplishments can be attributed to their sulfiphilic nature for effectively inhibiting polysulfide shuttling and their electrocatalytic properties for accelerating electrochemical reactions.

Among various 2D materials, molybdenum disulfide (MoS2), a reputed 2D material in the family of TMDs, has received intensive attention as the cathode host for LSBs owing to its natural abundance, sulfiphilic nature, ultrafast carrier mobility, less surface dangling bonds with minimal charge trapping, superior mechanical flexibility, and doping engineering and functionalization of 2D metal chalcogenides.22 Importantly, MoS2 can be mass-produced via liquid exfoliation and chemical vapor deposition techniques.23 Nevertheless, the obtained battery performance of LSBs based on MoS2 mandates further improvements (i.e., unfavorable rate performance and poor cycling durability) because of inadequate electrical conductivity and sluggish reaction kinetics during cycling.24 The electronic conductivity and redox kinetics of MoS2 can be exemplarily improved by employing sulfur-doped graphene frameworks supporting atomically dispersed 2H–MoS2,25 sodiated MoS226, and MoS2/MoN lateral heterostructures.27 In particular, the enhanced performance in the MoS2/MoN heterostructure is attributed to the electron donation from MoN to MoS2. As a result, MoS2 is n-doped and becomes more electrically conductive to accelerate the redox reaction of polysulfides and smoothen Li+ diffusion pathways.26 Moreover, the concept of heterostructures for magnifying conductivity extends to van der Waals heterostructures (vdWHs)28 in which distinct 2D materials are vertically assembled via vdW interaction. Specifically, vdWHs comprising MoTe2 and 2D electride [Ca2N]+:e– enable ultimate electronic transport of MoTe2 (i.e., carrier concentration of 1.6 × 1014 cm–2 and a prolonged electron diffusion length of around 100 nm).28 Inspired by these previous works, it is postulated that one can enhance battery performance by rationally utilizing a hybrid cathode host made from sulfiphilic MoS2. This host is coupled with a highly electrically conductive, electron-rich, chemically and thermally stable, low work function, and naturally abundant electron-donor material.

Given the requirements mentioned above, novel electride materials are potential candidates. The electride phase defines peculiar ionic compounds consisting of positively and negatively charged ions.29 The negative ions are simply electrons without nuclei which are weakly bound in the host structure, thereby acting as an electron gas to render superlative electron transport and a platform for topological materials.29 In 2003, the earth-abundant mineral mayenite ([Ca12Al14O33]4+:4O2– or simply known as C12A7), a compound in the family of cement-based calcium aluminates, was synthesized as the electride variant by extracting O2– ions from the C12A7 structure via chemical reduction and replacing them by anionic electrons.30 The resultant chemical formula is [Ca12Al14O33]4+:4e– abbreviated to C12A7:e–. Unlike air-susceptible 2D electride [Ca2N]+:e–, C12A7:e– is intriguingly the first inorganic electride thermally and chemically stable at ambient conditions.31 C12A7:e– exhibits exceptional electrical mobility and transport (conductivity reaching 1500 S cm–1 at room temperature), excellent electron concentration (2.0 × 1023 cm–3), ultralow work function (around 2.4 eV),31 and exceptional mechanical flexibility.32 The unique features enable numerous applications including superconductivity (the critical temperature 0.2–0.4 K), electron-field emitters,33 display devices, catalysis for ambient pressure ammonia production,34 and transparent conductive oxides.35 Beyond these, C12A7:e– is a promising material serving as an appropriate electron donor to intensify the electron transport and catalytic surface reactivity of MoS2 when combined as a heterostructure. In other words, the MoS2/C12A7:e– hybrid is hypothesized to be an efficient conductive host for circumventing the insulating nature of sulfur, inhibiting the shuttle effect, and enhancing the electrochemical reactions in LSBs.

Herein, this work aims at using the MoS2/mayenite electride (C12A7:e–) hybrid as the cathode host to resolve limitations of LSBs. Density functional theory (DFT) was employed to unravel the key aspects including electronic properties, the interfacial binding of Li2Sn on MoS2/C12A7:e–, Li+ diffusion, charge transfer mechanism, and SRR. We discovered that the MoS2/mayenite electride hybrid as a cathode host can effectively suppress polysulfide shuttling while enhancing reaction kinetics in LSBs. Apart from the known applications of C12A7:e–, this work sheds light on its unique role as an electrode material in LSBs.

Modeling and Computational Details

MoS2/C12A7:e– was considered as a hosting material for the sulfur cathode. Modeling of the MoS2/C12A7:e– atomic structure was based on the optimized 4 × 4 × 1 supercell of the single-layer MoS2 unit cell (a = b = 3.18 Å, c = 15.00 Å) placed on top of the cubic unit cell of C12A7:e–, resulting in a 166-atom unit cell. The C12A7:e–(100) plane was considered because it was successfully fabricated in the experiment.36 An adequate 15.0 Å vacuum gap was incorporated in the vertical direction of the pristine MoS2 and MoS2/C12A7:e– hybrids to eliminate self-interactions among periodic replicas.

DFT via the Vienna Ab initio Simulation Package (VASP)37,38 was employed to unravel the fundamental aspects of LSBs. The gradient-corrected Perdew–Burke–Ernzerhof (PBE) formalism was adopted to describe the exchange–correlation energy.39 The projector augmented wave was used to circumvent the rapidly oscillating wave functions of electrons near the ion cores.40 Here, the Mo (4p6 5s1 4d5), S (3s2 3p4), Ca (3s2 3p6 4s2), Al (3s2 3p1), O (2s2 2p4), and Li (1s2 2s1) were treated as the valence states in our calculations. Based on the modeled atomic structures, we employed the converged energy cutoff of 600.0 eV with the Brillouin zone integration of 3 × 3 × 1 (9 × 9 × 1) Monkhorst–Pack grids for structural optimization (calculation of electronic density of states).41 The structural optimization relied on the convergence criterion of Hellmann–Feynman forces acting on each atom (i.e., 20.0 meV/Å.). The total energies of the modeled structures were iteratively minimized until the energy difference between the two consecutive cycles of electronic self-consistency was less than 10–6 eV. Spin–orbit coupling was not considered in our work.

Interfacial interactions of Li2Sn on MoS2 and MoS2/C12A7:e– were investigated by evaluating the binding energies (Eb) according to the following equation

graphic file with name am4c05810_m001.jpg 1

where the first, second, and third terms on the right-hand side of eq 1 define the total energies of Li2Sn-adsorbed MoS2/C12A7:e–, bare MoS2/C12A7:e–, and isolated Li2Sn, respectively. Intrinsic van der Waals (vdW) forces existing at the interface of Li2Sn–MoS2/C12A7:e– were considered according to Grimme’s DFT-D3 approach.42 Furthermore, we calculated the ionic diffusion of Li+ ions on pristine MoS2 and MoS2/C12A7:e– using the climbing-image nudged elastic band (CI-NEB) method.43 Insightful analysis of bonds formed between Li2Sn and MoS2/C12A7:e– was carried out using the crystal orbital Hamilton population (COHP) which applies to studying chemical bonding in absorption, surface processes, and extended solids (i.e., covalent, ionic, and metallic crystals).44 COHP allocates the band-structure energy (here, DOS outputs from plane-wave-based VASP) to orbital–pair interactions (i.e., bond-weighted DOS of a pair of nearby atoms). A COHP diagram indicates bonding and antibonding contributions to the band-structure energy. Moreover, the integrated COHP (ICOHP) represents the contribution of a certain bond to the band energy, hinting at the bond strength. Moreover, Bader charge analysis was conducted to quantify the net atomic charges based on the reconstructed (all-electron) valence density.45

Moreover, electrochemical reactions in LSBs were investigated through SRR which is of considerable interest for achieving high-density kinetics energy storage.46 SRR defines a subtle 16-electron process for reversibly converting sulfur S8 to the end-product Li2S. Mathematically, the overall reaction can be written as S8 + 16Li+ + 16e– = 8Li2S in which there are a series of intermediates Li2Sn as follows9

graphic file with name am4c05810_m002.jpg 2
graphic file with name am4c05810_m003.jpg 3
graphic file with name am4c05810_m004.jpg 4
graphic file with name am4c05810_m005.jpg 5
graphic file with name am4c05810_m006.jpg 6

where * represents Li2Sn which is bound on MoS2 and MoS2/C12A7:e–. The changes in Gibbs free energy (ΔG) at constant pressure and temperature were calculated to quantify the feasibility of chemical reactions according to the equation ΔG = ΔE+ ΔEZPE – TΔS. Here, ΔE, ΔEZPE, and TΔS denote the difference in the binding energies of Li2Sn adsorbed on the host’s surface, zero-point energy (ZPE) associated with molecular vibrations of Li2Sn, and entropy variation between the Li2Sn-adsorbed MoS2 or MoS2/C12A7:e– and isolated Li2Sn. All of the ΔG corresponding to the reaction stages (eqs 2–6) were computed as the following expressions

graphic file with name am4c05810_m007.jpg 7
graphic file with name am4c05810_m008.jpg 8
graphic file with name am4c05810_m009.jpg 9
graphic file with name am4c05810_m010.jpg 10
graphic file with name am4c05810_m011.jpg 11

We set the temperature to ambient 298.15 K. ZPE was obtained by performing vibrational frequency calculations of Li2Sn on MoS2 and MoS2/C12A7:e– by calculating the second-order derivatives of the total energy with respect to the ionic positions via a method of finite differences. The atomic positions of MoS2 and MoS2/C12A7:e– were fixed during the calculations, whereas those of Li2Sn were allowed to move. Furthermore, the ΔG energies for the SRR of Li2Sn in a vacuum are also comparatively calculated to justify the role of the cathode hosts in the SRR.

Results and Discussion

Structural Properties of MoS2 and C12A7:e– Electride

As the first step, Figure 1a shows the crystal structure of monolayer MoS2. It is classified as a threefold-hexagonal symmetry with the space group P6̅m2 (no. 187).47 The S–Mo–S sandwich structure possesses the optimized lattice constant of a = 3.18 Å, and the Mo–S and S–S bond lengths are 2.41 and 3.11 Å, respectively. These crystallographic parameters are consistent with the reported works.48 Moreover, the electronic projected density of states (PDOSs) and the corresponding band structure as calculated by the PBE level of theory manifest a finite direct bandgap of 1.77 eV, agreeing with the reported value of 1.80 eV.49 Meanwhile, Figure 1b shows an optimized unit cell of the C12A7:e– electride comprising empty and occupied cages. There are two oppositely charged components including (i) positive [Ca24Al28O64]4+ and (ii) negative (H–)2. The first part represents a positively charged framework consisting of 12 crystallographic cages. The second part denotes a negatively charged extra-framework that occupies the 2 off-center cages. All cages exhibit S4 symmetry, where the symmetry axes pass through the cage poles (Ca–Ca). The calculated lattice constant a is 12.09 Å, respectively. The average Ca–Ca distance along the cage poles in empty cages is 5.73 Å, whereas that in the cages occupied by electrons is 4.18 Å. The electronic density of states and band structure of C12A7:e– are shown in Figure S1.50

Figure 1.

Figure 1

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(a) Top and side view of the crystal structure, electronic band structure, and PDOS of MoS2. (b,c) Optimized heterostructures of C12A7:e– and MoS2/C12A7:e–. (d) PDOS of MoS2 and optimized heterostructures of MoS2/C12A7:e–. (e) Charge density difference (Δρ) between MoS2 and MoS2/C12A7:e– as calculated by Inline graphic. The isosurface is set to be 0.001 e/Å3. (f) Band alignment diagram of MoS2, C12A7:e–, MoS2/C12A7:e– hybrid, and conventional current collectors (e.g., Al, Cu, and Ni).

Subsequently, Figure 1c depicts the relaxed atomic geometry of the MoS2/C12A7:e– heterostructure. MoS2 in the hybrid shows slight deformation as evidenced by the infinitesimal 0.83% deviation in the Mo–S bond lengths (2.39 and 2.48 Å as compared with original 2.41 Å in pristine MoS2). Likewise, the crystal structure of C12A7:e– in the heterostructure remains virtually retained as compared to the original C12A7:e– because there are small expansions in Ca–O and Al–O distances at the interface. Quantitatively, the Ca–O and Al–O distances at the surface before (after) making the heterostructure are 2.28 Å (2.34) Å and 1.82 Å (1.84 Å), amounting to 6.0 and 1.0% enlargements, respectively. Hence, both MoS2 and C12A7:e– nearly conserved their structural features after being merged into the heterostructure. This implies that the fabricated MoS2/C12A7:e– sample exhibits X-ray diffraction patterns of the component materials resembling their bulk counterparts. Intriguingly, there are bridging S–Ca bonds with the average bond distance of 2.48 Å formed at the boundary. This length is close to the sum of the ionic radii of Ca (1.84 Å) and S (1.00 Å) atoms51 and 2.549 Å of an ionic Ca–S bond in rock-salt CaS,52 indicating that the interface is dominated by ionic chemisorption.

Figure 1d comparatively shows the PDOS of MoS2 before and after being combined with C12A7:e–. MoS2 in the MoS2/C12A7:e– heterostructure adopts the metallic state characterized by the finite states of S-3p and Mo-4d orbitals at the Fermi energy. The resultant metallic character of MoS2 as an electrode material is beneficial for enhancing the electrochemical process in LSBs.5 Such semiconductor-to-metal transition is attributed to the formation of ionic S–Ca bonds at the boundary because the S atoms are highly more electronegative than the Ca atoms.53 This results in notable charge transfer from C12A7:e– to MoS2 as evidenced by the charge density difference plot in Figure 1e. Quantitatively, Bader charge analysis further indicates that Mo and interfacial S atoms accept 0.12e– from C12A7:e–, yielding the modification of Mo–S bond distances. Furthermore, the charge transfer mechanism is also contributed by the electron flow from C12A7:e– to MoS2 because of the lower work function of the former as depicted in Figure 1f. The accepted charges in MoS2 help intensify the surface chemical reactivity for interacting with heteropolar Li2Sn. The work function of MoS2/C12A7:e– is calculated to be 4.65 eV which is well matched with those of current-collecting materials of battery devices such as Cu (4.70 eV) and Ni (5.01 eV).54 The metallic feature of MoS2 and the appropriate work function of MoS2/C12A7:e– consequently facilitate electronic transport for accelerating the electrochemical process in LSBs.

Binding characteristics of Li2Sn with MoS2 and MoS2/C12A7:e–

It is expected that charge donation from C12A7:e– to MoS2 enhances MoS2/C12A7:e– surface affinity toward heteropolar Li2Sn.13 Pristine MoS2 and MoS2/C12A7:e– heterostructure are then utilized as the cathode host material to anchor Li2Sn binding.55 Li2Sn-DME/DOL binding is also considered as a reference for justifying polysulfide shuttling suppression. The lowest-energy conformation of each Li2Sn adsorbed on MoS2 and MoS2/C12A7:e– is obtained by attempting various adsorption orientations and sites. Figure 2 exclusively assembles the lowest-energy atomic configurations of Li2Sn adsorbed on MoS2 and MoS2/C12A7:e–. One can see that S8 energetically adopted the parallel orientation on both MoS2 and MoS2/C12A7:e– for intensifying the interfacial interaction, resembling the horizontal adsorption configuration observed in 2D anchoring materials such as graphene,55 black phosphorene,56 Fe3C,57 and MXenes.16 The bond distances from S8 to MoS2 (MoS2/C12A7:e–) are 3.37 (3.27) Å which are twice the atomic radius of S (1.00 Å).51 Hence, S8 is physisorbed on MoS2 (MoS2/C12A7:e–) via predominant vdW forces. Notably, the relatively shorter S8–MoS2/C12A7:e– distance manifests that S8 on MoS2/C12A7:e– experiences enhanced interaction than that in pristine MoS2.

Figure 2.

Figure 2

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Optimized structures of Li2Sn adsorbed on (a) MoS2 and (b) MoS2/C12A7:e–.

Meanwhile, Li2Sn is adsorbed on MoS2 and MoS2/C12A7:e– by forming Li–S bonds. This scenario is likely driven by the drastic difference in the electronegativity values between S (2.58) and Li (0.98).53 The energetically dissociative adsorption of all polysulfides on MoS2 and MoS2/C12A7:e– is not observed, ensuring Li2Sn conversion reversibility in the redox reaction. Li2Sn–MoS2 bond distances range from 2.20 to 2.58 Å. These distances are within the combination of atomic radii of S (1.00 Å) and Li (1.45 Å),51 characterizing the chemical S–Li bonds formed at the boundary. Notably, the MoS2/C12A7:e– composite induces a stronger interaction with Li2Sn. The shortened bond distances of Li2Sn–MoS2/C12A7:e– range from 2.08 to 2.58 Å. In addition, we calculated the binding energies associated with the structural conformations presented by using eq 1. All obtained binding energies are negative, indicating the attractive interaction. However, Figure 3 shows the strengths of the binding energies (Eb) of S8 and Li2Sn on MoS2 and MoS2/C12A7:e–. There are significant differences in the adsorption energies of Li2S4, Li2S2, and Li2S adsorbed on pristine MoS2 compared to their adsorption on MoS2/C12A7:e–. This is due to the markedly different geometrical orientations of Li2S4, Li2S2, and Li2S adsorbed on pristine MoS2 vs those on MoS2/C12A7:e–. In contrast, Li2S6, Li2S8, and S8 exhibit virtually similar orientations on both MoS2 and MoS2/C12A7:e–, resulting in less pronounced differences in the adsorption energies. The Li2Sn is adsorbed on MoS2 with competitively stronger Eb than DME and DOL, in accordance with reported works.58 Besides, the MoS2/C12A7:e– heterostructure further strengthens the interactions with polysulfide species compared with pristine MoS2. The Eb values of Li2Sn on MoS2/C12A7:e– surpass those on DME and DOL. In particular, MoS2/C12A7:e– greatly maximizes the binding of the liquid high and medium S-content polysulfides (Li2Sn, where n = 8, 6, and 4) which are the liquid phases and highly susceptible to dissolution (i.e., the major cause of polysulfide shuttling).9 Hence, it is conclusive that MoS2/C12A7:e– as the cathode host effectively suppresses polysulfide shuttling.

Figure 3.

Figure 3

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Magnitudes of binding energies of Li2Sn adsorbed on MoS2 and MoS2/C12A7:e-. Here, DOL and DME abbreviate 1,3-dioxolane and 1,2-dimethoxyethane, respectively, which are common electrolytes used in LSBs.

Electronic Properties of MoS2/C12A7:e– Interacting with Li2Sn

The strong binding affinities between Li2Sn and MoS2/C12A7:e– are responsible for the suppression of polysulfide shuttling. Moreover, analyzing the variations in the electronic properties of Li2Sn when interacting with MoS2/C12A7:e– is crucial for comprehending the chemical bonds formed between Li2Sn and MoS2/C12A7:e–. Figure 4a–f displays the total density of states (TDOSs) of Li2Sn before and after adsorption. The bulk sulfur (cyclo-S8) is categorized as an insulator because of its large finite energy gap. Notably, the TDOS of S8 undergoes a notable change after adsorption. Likewise, all Li2Sn intermediates are characterized as electrical insulators owing to the finite energy gaps. Their TDOSs undergo remarkable variations after interacting with MoS2/C12A7:e–, attesting to chemisorption. Furthermore, we calculated charge density difference (Δρ) to support TDOS analysis according to the following expression

graphic file with name am4c05810_m012.jpg 12

where the first, second, and third terms on the right-hand side of eq 12 represent the electron density of the Li2Sn-adsorbed MoS2/C12A7:e–, MoS2/C12A7:e–, and the corresponding isolated Li2Sn molecule, respectively. Figure 4c,d illustrates Δρ of the polysulfides adsorbed on MoS2/C12A7:e–. The adsorption of bulk S8 on MoS2/C12A7:e– causes insignificant charge redistribution, consistent with the weak Eb values and indicative of predominant physisorption. By contrast, the charge transfer becomes pronounced in the case of Li2S6, supporting enhanced Eb values and a tendency toward predominant chemisorption. Specifically, Li2S6 donates charges to the MoS2/C12A7:e– host, resulting in an upshift in the TDOS of the host, as described previously. Likewise, charge transfer is particularly immense in Li2S, supporting the notion of prominent chemisorption.

Figure 4.

Figure 4

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(a–f) TDOS of Li2Sn before and after interacting with MoS2/C12A7:e–. The Fermi energy is shifted to zero for the sake of comparative visualization. (g,i) Charge density difference between Li2Sn-adsorbed MoS2/C12A7:e–.

Additionally, the amount of charge transfer is computed using the Bader charge approach, as shown in Figure 5a. Quantitatively, S8, Li2S6, and Li2S donate 0.0004e (0.0005e), 0.003e (0.006e), and 0.02e (0.03e) to MoS2 (MoS2/C12A7:e–), respectively. The donated charges consequently magnify the Li2Sn–MoS2/C12A7:e– interactions as shown by the analysis of PDOS and COHP in Figure 5b–d. The S-3p orbitals from S8 weakly interact with those from MoS2/C12A7:e–, resulting in infinitesimal ICOHP of −0.097 eV/bond. By contrast, there are immense chemical interplays among Li-2s of Li2Sn (here, selectively shown by high S-content Li2S6 and low S-content Li2S), S-3s, and S-3p orbitals of MoS2/C12A7:e–, as indicated by the antibonding peaks in the conduction band. These antibonding peaks are derived from the charge donation from the Li-2s state to S-3s and S-3p states to form antibonding states of ionic Li–S bonds above the Fermi energy. Moreover, these bonds possess pronounced bond energies, as indicated by the maximized ICOHP of −0.539 and −0.767 eV/bond for Li2S6 and Li2S, respectively. Note that the shape of the Mo-4d states remains nearly conserved during the interaction with Li2Sn. This is because the atomic geometry of MoS2 comprises the Mo atoms in the central sublayer sandwiched by the outer S atoms. Thus, they do not directly interact with Li2Sn.59 Based on our analysis of electronic properties, charge-transfer ionic Li–S bonds formed at the Li2Sn–MoS2/C12A7:e– interface are conclusively responsible for mitigating polysulfide dissolution in DOL and DME electrolytes.

Figure 5.

Figure 5

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(a) Bar chart of Bader charge transfer from Li2Sn to MoS2 and MoS2/C12A7:e–. (b–d) PDOS of Li2Sn-adsorbed MoS2/C12A7:e– in which the corresponding COHP and ICOHP of the bonds between Li (S) of Li2Sn and MoS2/C12A7:e– are also shown.

Li+ Ionic Diffusion

Li+ ionic diffusivity on the cathode of LSBs plays a central role in the reaction kinetics of polysulfide conversion in the redox reaction process as probed by an electrochemical current–voltage measurement.60 The Li+ ions energetically favor the low diffusion energy to maximize the diffusion rate basically according to Arrhenius equation Inline graphic, where k, A, Ea, R, and T represent the rate constant, frequency factor depending on the lattice vibrations of diffusion state, diffusion barrier, universal gas constant, and temperature, respectively.61 In particular, low diffusion energy at a finite temperature on the surface of anchoring materials can particularly enhance the chemical reaction between lithium (anode) and sulfur (cathode).62 According to this, we compute the diffusion barriers of Li+ ions on MoS2 and MoS2/C12A7:e– using the CI-NEB approach to justify the catalytic anchoring roles of MoS2 and MoS2/C12A7:e– in lowering the diffusion potentials. Figure 6a–c shows the diffusion profiles and most energetic paths of the Li+ ion on MoS2 and MoS2/C12A7:e–. The preferential route is obtained by considering the path connecting two neighboring points on the surfaces of MoS2 and MoS2/C12A7:e– at which the Eb values of Li+ are the strongest. The diffusion barriers of Li+ on pristine MoS2, interfacial-strained MoS2, and MoS2/C12A7:e– are 0.25,63 0.15, and 0.12 eV, respectively. The interfacial-strained MoS2 here defines the intrinsic strain imposed by the lattice mismatch between MoS2 and C12A7:e–. The reduction in the ionic diffusion barrier in MoS2/C12A7:e– is simultaneously contributed by (i) the interfacial strain imposing on MoS2 from the lattice incompatibility and (ii) the charge transfer between MoS2 and C12A7:e–. Notably, the strain plays a more dominating role, accounting for more than 50% of the barrier reduction. The pathway follows a curved trajectory between the adjacent hollow sites of the Mo–S hexagons.62 This is because the ion follows the path experiencing less potential, which is the hollow site above the inner Mo atoms. The presence of C12A7:e– remarkably reduces the barrier, evidencing its catalytic ability for Li+ ions’ mobility when in contact with MoS2. Moreover, the obtained values are among the lowest as compared to those of 2D materials such as graphene (0.32 eV),64 single-atom-catalyzed graphene,65 a family of metal sulfides (0.12–0.26 eV),60 and Ti2CS2 MXenes (0.12 eV),16 all calculated from the similar DFT-based NEB approach. The variations in the ionic potential profiles of MoS2 and MoS2/C12A7:e– are associated with the metallic nature of MoS2 in the MoS2/C12A7:e– hybrid as compared with the original semiconducting MoS2. According to the d-band model in metal catalysts,66 the d-band center (εd) is quantitatively computed by the following expression

graphic file with name am4c05810_m014.jpg 13

where ε and nd denote the energy and DOS of the Mo-3d states integrated in a particular range of energy. Moreover, the p-band centers are determined using the same concept as in eq 13.9 The εd values of Mo-3d states in MoS2 and MoS2/C12A7:e– are −3.08 and −3.03 eV, respectively. The values of Mo p-band centers in MoS2 and MoS2/C12A7:e– are −3.335 and −3.554 eV, respectively. Meanwhile, those of S p-band centers in MoS2 and MoS2/C12A7:e– are −3.463 and −3.722 eV, respectively. In particular, the upshift in the S-3p states intensifies the interaction with Li-s states, being consistent with the COHP bonding analysis and Bader charge analysis as presented in the former section.9

Figure 6.

Figure 6

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(a) Diffusion barriers of Li+ ion and (b,c) corresponding path on MoS2 and MoS2/C12A7:e–.

Sulfur Reduction Reaction

The sluggish SRR is one of the major intrinsic drawbacks plaguing the efficiencies of the LSBs. This also results in the extended exposure time of Li2Sn to the electrolytes, worsening the shuttle phenomenon and inefficient sulfur utilization.5 In addition to the improved ionic diffusivity, we elucidate how MoS2 and MoS2/C12A7:e– cathode hosts contribute to accelerating the SRR during the operation of the LSBs. Figure 7a shows the ΔG profiles of the SRR of transforming S8 to Li2S in the presence of MoS2 and MoS2/C12A7:e–, as calculated from eq 2. The first steps ΔG1 for converting solid S8 to liquid Li2S8 are the spontaneously exothermic processes (ΔG < 0) on both MoS2 and MoS2/C12A7:e–, releasing the energies of −3.338 and −3.325 eV, respectively. The initial transition from the S8 ring to the long-change Li2S8 is an energetically favorable process as found experimentally.9 Comparatively, the obtained values are even less than −3.016 eV for converting solid S8 to liquid Li2S8 in vacuum. This means that the processes with the cathode hosts enable the release of more energy to allow Li2S8 to be in a more stable state than that in the vacuum. This corroborates with the occurrence of Li2S8 in the discharging voltage profile.5 Likewise, the second reduction steps ΔG2 for the Li2S8-to-Li2S6 transition on MoS2, MoS2/C12A7:e–, and in vacuum remain the spontaneously exothermic processes with the significantly reduced energies of −0.146, −0.229, and −0.216 eV, respectively. This indicates the catalytic contribution of MoS2/C12A7:e– in accelerating the activity toward Li2S6.67 On the other hand, the remaining processes including ΔG3 (Li2S6-to-Li2S4), ΔG4 (Li2S4-to-Li2S2), and ΔG5 (Li2S2-to-Li2S) are all endothermic. Among all discharge states, the conversion from Li2S4 to Li2S2 notably serves as the rate-determining step requiring 0.64068 and 0.494 eV on MoS2 and MoS2/C12A7:e–, respectively. This manifests that the MoS2/C12A7:e– hybrid plays a superior role in SRR than does MoS2 alone. According to eq 10, the obtained rate-limiting step is predominantly ascribed to the increase in the binding energy and entropy from Li2S4 to Li2S2. Additionally, this elusive Li2S4-to-Li2S2 transition is typical in the SRR of LSBs and it is confirmed by various independent experimental studies (see the review5). For the sake of comparison, this value is among the lowest with other 2D cathode hosts in LSBs, such as 0.513 eV of boron nitride69 and nitrogen-doped graphene 0.72 eV.70 Overall, it is conclusive that MoS2/C12A7:e– as the cathode host can enhance the electrical transport of the cathode, suppress the shuttle effect, and offer efficient electrochemical kinetics.

Figure 7.

Figure 7

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Gibbs free-energy profile of the sulfur reduction in a vacuum, on MoS2, and on MoS2/C12A7:e–.

Conclusions

In conclusion, we carried out first-principles DFT simulations to investigate the key aspects including electronic properties, interfacial binding, Li+ diffusion, charge transfer mechanism, and SRR of LSBs by using MoS2/C12A7:e– as a cathode host. The findings indicate that the MoS2/C12A7:e– hybrid boosts battery efficiency over MoS2. MoS2 in the hybrid undergoes a semiconductor-to-metal transition, induced by electron donation from lower work function C12A7:e–. Moreover, the resultant band positions of MoS2/C12A7:e– are well matched with those of the current collecting materials (i.e., Cu and Ni). The hybrid cathode host effectively prevents polysulfide shuttling by forming stronger bonds with Li2Sn than with electrolytes with Li2Sn. This is because of charge-transfer ionic Li–S bonds. Importantly, MoS2/C12A7:e– not only promotes rapid reaction kinetics by reducing energy barriers for the Li+ diffusion but also lowers the Gibbs free energies of the SRR for the effective S8-to-Li2S conversion.

Acknowledgments

This work is financially supported by the Office of the Ministry of Higher Education, Science, Research and Innovation, Thailand (grant no. RGNS63-005). T.K. acknowledges the high-performance computing facility provided by ThaiSC. Furthermore, T.K. would like to acknowledge Prof. Vittaya Amornkitbamrung for his kind mentoring. This work was supported by the NCI Adapter Scheme, with computational resources provided by NCI Australia, an NCRIS-enabled capability supported by the Australian Government. A.W. and P.S. would like to acknowledge funding support from the NSRF via the Program Management Unit for Human Resources and Institutional Development, Research and Innovation (grant nos. B05F640218 and B05F650023), National Higher Education Science Research and Innovation Policy Council, and Chiang Mai University.

Supporting Information Available

The Supporting Information is available free of charge at https://pubs.acs.org/doi/10.1021/acsami.4c05810.

  • Electronic density of states and band structure of C12A7:e–, optimized structures of S8 and Li2Sn, binding energies of Li2Sn on various 2D materials, and diffusion barriers of Li+ on various 2D materials (PDF)

The authors declare no competing financial interest.

Supplementary Material

am4c05810_si_001.pdf (334.8KB, pdf)

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