A Spin-polarized DFT study of functionalized MXenes as effective anchor materials in lithium-sulfur batteries

Abstract

Lithium-sulfur (Li-S) batteries have attracted great enthusiasm in recent years due to ultra-high theoretical energy densities, abundant sulfur electrode resources and low price. Despite the severe shuttle effect of lithium polysulfides (LiPSs), the poor conductivity of S₈ and its intermediate products, and the relatively slow dynamics, pose significant challenges for the commercial application of Li-S batteries. Here, functionalized MXenes M₂CT₂ (M = V, Cr, Mn, and Mo; T = F and O) used as the sulfur host are studied to build multifunctional cathodes via spin-polarized first-principles calculation. Through analyzing the adsorption energy and configuration of S₈/Li₂S_n adsorbed M₂CT₂, it is found that spin polarization is indispensable to the Li-S battery calculation of MXenes with transition metals. With the spin polarization calculation, the M₂CT₂ exhibit moderate anchoring strengths and stable adsorption structures, which effectively mitigates the polysulfide shuttle phenomenon. The low decomposition barriers of Li₂S (0.27–1.00 eV) and low diffusion barriers of Li⁺ (0.11–0.44 eV) of M₂CT₂ are observed, which effectively improve the rate performance of batteries. Among the studied MXenes, V₂CO₂ and Mo₂CO₂ are the best choices of host materials for LiPSs with metallic characteristics, outstanding electrocatalysis performance, low decomposition barriers of Li₂S, and diffusion barriers of Li⁺. This work provides important insights into spin-polarized electrode materials for enhanced energy storage capabilities by investigating the application of intrinsic magnetic MXene compounds.

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Introduction

With the continuous growth of energy demand, it is an urgent task to develop high-performance and affordable rechargeable batteries. Lithium-sulfur (Li-S) batteries have attracted substantial research interest in recent years, owing to their high theoretical capacity and specific energy.^1–4 Based on the two-electron transfer redox reaction involving multiple steps, the theoretical capacity of Li-S batteries can reach up to 1675 mA h g⁻¹. Moreover, sulfur is abundant on earth and inexpensive.^5,6 As a result, the combination of high theoretical capacity and low cost endows Li-S batteries with great application potential in the field of commercial batteries.⁷ However, there is still a long journey ahead for Li-S batteries to achieve commercial application, which is hindered by several thorny issues. For example, the electrical conductivities of S₈ and Li₂S_n are poor, and the reaction kinetics is sluggish due to the high decomposition energy barrier of Li₂S.^8,9 During the cycling process, intermediate products of lithium polysulfides (LiPSs) are constantly generated. These intermediates can dissolve in the electrolyte, resulting in a reduction of active sulfur species. In particular, long-chain LiPSs (Li₂S_n, n = 4, 6, and 8) will migrate to the lithium metal anode and deposit there by shuttling through the separator, which also results in a rapid decline in the battery's capacity.^10,11 In the future, more efforts should be devoted to suppressing the shuttle effect. An ideal sulfur composite cathode is capable of suppressing the shuttle effect by moderately anchoring Li₂S_n, enhancing the sulfur reduction reaction (SRR), and improving the reaction kinetics.^12–14

MXenes, a novel class of two-dimensional transition metal carbides and nitrides, exhibit remarkable capabilities in electrochemical energy storage systems. These materials combine metallic-grade conductivity with ultrafast charge migration kinetics and exceptional surface charge storage capacity, attributes that have positioned them at the forefront of energy research since their 2011 discovery.¹⁵ Following Gogotsi's pioneering synthesis of Ti₃C₂ in 2011, titanium-based MXenes have dominated Li-S battery research due to their exceptional interfacial compatibility and sulfur confinement capabilities.¹⁶ After the first synthesis of Ti₃C₂, M₃C₂ has been extensively studied and is considered to be a cathode material for Li-S batteries with excellent performance. Recent experimental studies proved that M₂C is thinner than M₃C₂, which has also proven to be a good anchoring material for Li-S batteries with high specific capacity and long cycle life.^17,18 MXenes synthesized via HF etching typically exhibit surface terminations dominated by oxygen (O) and fluorine (F) groups, which could enhance material stability, optimize interface interaction and improve electrochemical performance.^19,20 According to previous reports, Ti₂C(OH)₂ and V₂C(OH)₂ exhibit excessively strong interfacial interactions with LiPSs, inducing premature decomposition of active species.^21,22 Conversely, oxygen-terminated MXenes like Ti₂CO₂ and Ti₂NO₂ demonstrate optimal chemisorption energetics for LiPSs anchoring. Ti₂CF₂ and Ti₂NF₂ achieve balanced binding interactions, with systematic analyses revealing that O- and F-terminated MXenes collectively offer superior polysulfide confinement capabilities compared to OH-terminated MXenes.^23,24 The interaction between the polysulfides and the hydroxide-ternibated or no functionalized MXene is very strong, which could distort the polysulfides and the MXene materials. Therefore, they cannot effectively suppress the shuttle effect.^25–27 Recently, V₂NO₂ and V₂NF₂ MXenes have been widely investigated due to low decomposition barriers of Li₂S and diffusion barriers of Li⁺, which could accelerate reaction kinetics during the discharge and charging process. The structural diversity of MXenes motivates systematic exploration of their electrochemical properties, particularly focusing on surface-engineered variants with functional groups. It is an interesting task to explore other MXene materials that possess outstanding electrochemical performance, particularly those MXene materials with a large number of functional groups and those that have undergone surface modification.^28,29

In our work, a series of M₂CT₂ (M = V, Cr, Mn, and Mo; T = F and O) are selected as sulfur hosts to construct multifunctional cathode material via spin-polarized first-principles calculations. We note the adsorption energy of S₈/Li₂S_n adsorbed M₂CT₂ is greatly reduced from non-spin-polarized system to spin-polarized system. Spin-polarized systems provide a critical framework for comprehensively elucidating interfacial phenomena in transition metal composites at atomic scales. The adsorption energy of S₈/Li₂S_n adsorbed M₂CT₂ range from 0.63 eV to 5.15 eV, which indicates the M₂CT₂ could inhibit the shuttle effect and reduce the loss of active sulfur species in the polarization system. We find that S₈ can be efficiently reduced to Li₂S with a small rare-limiting step (0.51–1.37 eV), indicating the superior sulfur reducing ability of M₂CT₂. Moreover, the decomposition barriers of Li₂S and migration barriers of Li⁺ for M₂CO₂ are lower than those for M₂CF₂ during charging process, which ultimately promotes redox kinetics during cycling. Based on the finding presented above, we identified V₂CO₂ and Mo₂CO₂ as promising candidates for host materials with moderate anchoring strength and excellent electrocatalytic performance. Systematic analysis of MXene functionalities in this work provides actionable strategies for tailoring their interfacial properties toward sulfur cathode optimization.

Computational details

Spin-polarized density functional theory (DFT) calculations are carried out using the Vienna Ab Initio Simulation Package (VASP), employing the Perdew–Burke–Ernzerhof (PBE) generalized gradient approximation (GGA) functional to describe exchange-correlation effects.^30,31 The core-electron interactions are treated through the projector-augmented wave (PAW) methodology, with a plane-wave basis set truncated at 500 eV kinetic energy. For Brillouin zone integration in the 4 × 4 supercell's two-dimensional geometry, a Monkhorst–Pack grid of 5 × 5 × 1 k-points is implemented.³² To better characterize the strongly correlated 3d electron systems, we incorporate the Hubbard U correction within the GGA + U framework.³³ According to the previous works, the U values of V, Cr, Mn, and Mo atoms are set to be 3, 4, 3, and 3 eV, respectively.^22,34–36 Long-range van der Waals interactions are accounted for using the DFT-D3 empirical dispersion correction scheme.³⁷ A minimum vacuum spacing of 20 Å perpendicular to the surface plane is maintained to eliminate artificial periodic interactions. Structural optimizations are performed until achieving convergence thresholds of 1.0 × 10⁻⁵ eV per atom for total energy and 0.01 eV Å⁻¹ for residual forces. The migration mechanisms of Li⁺ on M₂CT₂ surfaces are investigated through the climbing image nudged elastic band (CI-NEB) method, which enables precise determination of minimum energy pathways and activation barriers.³⁸ The thermodynamic stability of Li₂S adsorption on V₂CO₂ and Mo₂CO₂ is simulated by ab initio molecular dynamics (AIMD) at 300 K. The NVT ensemble simulations are conducted for 5 ps with a time step of 1.0 fs, corresponding to 5000 simulation steps. Detailed calculation equation of adsorption energies (E_ads), charge density difference (CDD), and the equations of Gibbs free energy change (ΔG) are comprehensively derived in the ESI† section.

Results and discussions

Configuration, magnetic ground state and electronic conductivity of M₂CT₂

Investigating MXenes with different magnetic transition metal constituents offers a strategic pathway to leverage their compositional versatility for improving sulfur electrochemistry in Li-S systems. Therefore, the M₂C(F/O)₂ with the transition metals (M = V, Cr, Mn, and Mo) are chosen as sulfur hosts in our work. The monolayer M₂C MXene adopts a P3m1-symmetric honeycomb lattice, featuring a stratified architecture where hexagonally arranged carbon atoms are intercalated between two metallic M layers, as structurally resolved in Fig. 1a.³⁹ There are five possible configurations for terminal groups (T = F and O) absorbed on the M atom in Fig. S1†: (a) T atoms locate right above the M atoms (Top sites); (b) T atoms locate at the hollow sites of adjacent C atoms (Hcp sites); (c) T atoms locate at the hollow sites of contralateral M atoms (Fcc sites); (d) T atoms located at the Fcc and Hcp sites (Fcc–Hcp); (e) T atoms located at the Hcp and Top sites (Fcc–Hcp). The stable configurations of M₂CT₂ are Fcc structures, except for Cr₂CO₂ and Mo₂CO₂, which are Hcp structures as shown in Fig. 1c. The basic information of their lattice parameters is listed in Table S1.†

Fig. 1 (a) Top and side views of MXene monolayer with different adsorption sites (hollow site (H), carbon-top site (C) and transition metal-top (T)). (b) Geometrical structures of S₈ and Li₂S_n molecules. (c) Side view of the stable atomic structures of V₂CT₂, Cr₂CT₂, Mn₂CT₂, and Mo₂CT₂ with the terminal groups T = O and F. Yellow, green, dark red, brown, light red, white, blue, dark purple and light purple color balls represent the S, Li, V, C, O, F, Cr, Mn and Mo atoms, respectively.

Considering that magnetic transition metal in the M₂CT₂, five types of magnetic ground state structures are calculated in Fig. 2: non-magnetic (NM), ferromagnetic (FM), and antiferromagnetic (AFM1, 2, and 3). After relaxation, it is found that stable structures of V₂CO₂ and Mo₂CO₂ are NM, while the stable structures of other M₂CT₂ are AFM, except that Mn₂CF₂ and Cr₂CO₂ are FM structures. The calculated magnetic ground states are same as the results of the previous studies.^34,40–42 According to the crystal field theory, the magnetic exchange interactions in the 2D magnetic materials mainly originate from two kinds of mechanisms:^43–45 (i) the super-exchange ferromagnetism between the nearest intra-layered M atoms via the outmost oxygen atom (e.g., Cr–O–Cr and Mn–F–Mn). (ii) The antiferromagnetic interactions between the nearest interlayered M atoms via the carbon atom (e.g. Cr–C–Cr and Mn–C–Mn). In this case, the ferromagnetic super-exchange processes start to compete with the antiferromagnetic super-exchange mechanism.

Fig. 2 Magnetic state configurations of M₂CT₂: (a–e) NM, FM, AFM1, AFM2, AFM3, respectively (taking a 2 × 2 V₂CF₂ supercells as an example). The directions of the arrow represent the spin up and down.

Meanwhile, the limited conductivity of sulfur represents another key challenge impeding the advancement of Li-S batteries. To further study the electronic structural properties of M₂CT₂, we calculate the projected density of states (PDOS) of M₂CT₂. As shown in Fig. 3, V₂CO₂ and Mo₂CO₂ show significant metallic conductivity, while the others are semiconductors, except for Mn₂CF₂, which is a special semimetal. The DOS of the M₂CT₂ mainly comes from the d-orbital of the transition metal. Furthermore, it is found that the d_xy and d_x²−y² orbitals are degenerated, while the d_xz and d_yz orbitals are degenerated at their respective energy levels as presented in Fig. S2.†

Fig. 3 Projected density of states of (a) V₂CF₂, (b) Cr₂CF₂, (c) Mn₂CF₂, (d) Mo₂CF₂, (e) V₂CO₂, (f) Cr₂CO₂, (g) Mn₂CO₂, and (h) Mo₂CO₂. The Fermi level is set at zero.

Anchoring of M₂CT₂ to inhibit shuttle effect

During Li-S battery discharge, LiPSs dissolve into the electrolyte and shuttle toward the anode through the separator, leading to progressive depletion of electroactive sulfur species. Optimal cathode materials require appropriate adsorption energies to immobilize LiPSs and prevent active sulfur depletion.^46,47 Before to study the anchor ability of M₂CT₂ for S₈ and Li₂S_n, the S₈ and Li₂S_n are optimized in 10 × 10 × 10 ų cell, and the relaxed structures are shown in Fig. 1b. The S–S bond length in S₈ is determined to be 2.06 Å. The Li–S bond lengths in Li₂S₈, Li₂S₆, Li₂S₄, Li₂S₂ and Li₂S are determined to be 2.36 Å, 2.39 Å, 2.38 Å, 2.22 Å and 2.09 Å, respectively. These structures of our work are consistent with the previous reports.^40,48,49

To evaluate the anchor performance of the functionalized M₂CT₂, we calculated the adsorption energy (E_ads) of S₈ and Li₂S_n on these M₂CT₂ surface as a key parameter, which is defined in the ESI.† The positive values of the adsorption energies indicate that adsorption could proceed spontaneously, and moderate adsorption energy represents good anchoring ability. The E_ads with spin-polarized and non-spin-polarized calculations are listed in Table 1 and Table S2,† respectively. The calculated E_ads for S₈ and Li₂S_n on V₂CT₂ and Cr₂CT₂ are well consistent with those from the previous report.⁵⁰ For the adsorption of the S₈ and Li₂S_n, a similar binding trend is observed in all the calculated M₂CT₂ as shown in Fig. 4. For all the calculated M₂CT₂, the E_ads values of the S₈ are comparable. In addition, we note that the E_ads values increase with a decrease in the size of LiPSs.

Table 1 The absorption energies E_ads (eV) of the S₈ and LiPSs on M₂CF₂ and M₂CO₂ with the spin-polarized calculations

MXene	S8	Li2S8	Li2S6	Li2S4	Li2S2	Li2S
V2CF2	0.80	1.14	1.05	1.47	2.08	2.27
Cr2CF2	0.76	1.01	0.97	0.91	1.24	1.41
Mn2CF2	0.63	1.17	1.16	1.99	2.35	2.58
Mo2CF2	0.88	1.19	1.12	1.10	1.72	2.50
V2CO2	1.02	2.07	2.02	3.30	3.48	4.37
Cr2CO2	1.04	2.40	2.12	3.64	4.10	3.73
Mn2CO2	0.89	2.37	2.30	3.70	4.27	4.01
Mo2CO2	1.38	3.44	3.13	4.38	4.24	5.15

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Fig. 4 The absorption energies of S₈/Li₂S_n absorbed on (a) M₂CF₂ and (b) M₂CO₂ monolayer. The dotted and solid lines represent the absorption energy with the spin-polarized and the non-spin-polarized calculations, respectively.

It is also found that the E_ads of S₈ and Li₂S_n on M₂CT₂ surface with the spin-polarized calculations are significantly reduced compared with those of the non-spin-polarized calculations. The E_ads of Li₂S_n on V₂CF₂ and Mn₂CF₂ are reduced by 0.73–1.14 eV from the non-spin-polarized system to the spin-polarized system, while the E_ads values of Li₂S_n absorbed on Cr₂CO₂ decreases by 0.62–2.44 eV. Meanwhile, the E_ads of Li₂S_n absorbed on Cr₂CF₂ and Mo₂CF₂ decreases by 2.10–3.55 eV, which is nearly three times that of V₂CF₂ and Mn₂CF₂.^50–52 The effect of the spin-polarization also reflected from the relaxed configuration of M₂CT₂–S₈/Li₂S_n. As presented in Fig. 5 and Fig. S3.† The atomic structure of Mn₂CO₂ cannot even maintain at the NM state. The short-chain polysulfides break their bonds, and the substrate M₂CT₂ undergo severe deformation under the spin-polarized calculation. The spin polarization system has a significant effect on the absorption process for the magnetic M₂CT₂. Therefore, we take the accurate spin-polarized calculation in the subsequent. The adsorption strength of the M₂CT₂ for Li₂S_n follows the sequence: Mn₂CF₂ > V₂CF₂> Mo₂CF₂ > Cr₂CF₂ and Mo₂CO₂ > Mn₂CO₂ > Cr₂CO₂ > V₂CO₂, respectively. It clearly shows that the E_ads of S₈/Li₂S_n on M₂CT₂ are much smaller than that on M₂C.⁵² The enhanced covalent bonding characteristics between Li and S atoms in low-order LiPSs facilitate accelerated electron transfer from Li to S, creating vacant d-orbitals capable of accepting electrons from transition metal (T) atoms. Computational analyses reveal that short-chain LiPSs (Li₂S₂/Li₂S) and Li₂S₄ adopt specific geometric arrangements where dual lithium atoms remain proximal to the substrate surface, demonstrating strong chemisorption behavior. Conversely, long-chain species (Li₂S₈ and Li₂S₆) exhibit monodentate adsorption configurations with a single lithium atom near the interface, accompanied by horizontally oriented S₈ molecular adsorption. Crucially, all adsorbed LiPSs configurations maintain structural integrity during electrochemical cycling, showing no signs of dissociation. It ensures reliable polysulfide retention on M₂CT₂ surfaces throughout charge/discharge processes.

Fig. 5 The optimized structures of S₈/Li₂S_n adsorption on (a) V₂CT₂, (b) Cr₂CT₂, (c) Mn₂CT₂, and (d) Mo₂CT₂ monolayer with the spin-polarized calculations.

To effectively suppress polysulfide shuttling, the adsorption energetics of high-order Li₂S_n (n = 4, 6, and 8) on the sulfur host matrix should surpass those with conventional ether-based electrolytes 1,3-dioxolane/1,2-dimethoxyethane (DOL/DME), thereby preventing electrolyte-mediated active material dissolution and capacity fade.⁵³ Given the preferential solubility of high-sulfur-content LiPSs in organic electrolytes, we systematically evaluated the adsorption energetics of long-chain LiPSs (Li₂S₈, Li₂S₆, and Li₂S₄) on DOL and DME. Structural optimizations of both isolated electrolyte molecules and LiPS-anchored complexes are shown in Fig. S4,† which reveal comparable polysulfide-trapping capabilities between DOL and DME, with adsorption energies spanning 0.75–0.90 eV. These values are significantly lower than those observed for M₂CO₂ substrates and marginally reduced compared to M₂CF₂ interfaces, highlighting the limited chemical confinement capacity of organic electrolytes relative to engineered MXene surfaces. Based on the aforementioned analysis, M₂CT₂ is capable of anchoring S₈/Li₂S_n with a suitable E_ads value, while, it can preserve the integrity of the adsorption structures. In a tug-of-war with solvent molecules, M₂CO₂ shows stronger chemical affinity for the LiPSs, which reveals its potential to serve as a type of host material for sulfur cathodes.

In order to explore the essence of the adsorption behavior between S₈/Li₂S_n and the surface of M₂CT₂, the CDD between M₂CT₂ monolayer and S₈/Li₂S_n are calculated as shown in Fig. 6 with the calculation details located in ESI†. The S₈ and M₂CF₂ interface exhibits negligible charge transfer, signifying the absence of chemical bond formation between the adsorbed sulfur species and the substrate. In contrast, S₈/Li₂S_n adsorption on M₂CO₂ exhibits significantly enhanced interfacial electron redistribution. Lithiation induces pronounced electron accumulation at Li₂S_n/M₂CT₂ interfaces, accompanied by intensified charge transfer indicative of Li–T (T = F/O) covalent bond formation. Concurrently, expanding electron-deficient domains within Li₂S_n reflect destabilization of Li–S and S–S bonding interactions. It can be deduced that chemical bonds are established between Li₂S_n and M₂CT₂ during the lithiation process, which agrees with the trend of the adsorption energy. Moreover, the decomposition pathways of LiPSs are critically governed by two concurrent mechanisms: interfacial stabilization through Li–T (T = F/O) covalent bond formation and structural destabilization via progressive cleavage of Li–S and S–S bonds within Li₂S_n species.

Fig. 6 Side views of CDD maps of S₈ and Li₂S_n on (a) V₂CT₂, (b) Cr₂CT₂, (c) Mn₂CT₂, and (d) Mo₂CT₂ monolayer (isosurface = ± 0.002 au for Li₂S_n, and isosurface = ± 0.0002 au for S₈). The yellow and blue region indicate electron accumulation and depletion, respectively.

The electronic structures of the substrate undergo changes following LiPSs adsorption, which is investigated by analyzing the PDOS of M₂CT₂ after Li₂S adsorption. The PDOS of M₂CT₂ monolayer after the adsorption of Li₂S are shown in Fig. 7. Compared with the PDOS of M₂CT₂ monolayer, the Fermi level of M₂CT₂-Li₂S has shifted upwards with the electronic state of S atoms appearing in the band gap for the Mo₂CF₂−, Cr₂CF₂−, Mn₂CO₂−, Mo₂CO₂−, and Cr₂CO₂-Li₂S systems. Therefore, the band gap of V₂CF₂ and Cr₂CF₂ became narrow, while Mo₂CF₂, Cr₂CO₂, and Mn₂CO₂ exhibit metallic characters after the adsorption of Li₂S. Overall, the conductivity of M₂CT₂ is enhanced with the adsorption of Li₂S, which is benefit for the electron transportation.

Fig. 7 Projected density of states of (a–h) M₂CT₂ monolayer after the adsorption of Li₂S, respectively. The Fermi level is set at zero.

Electrocatalysis of M₂CT₂ on discharge/charging kinetics

The chemisorption intensity of Li₂S_n species critically modulates electrochemical reaction kinetics. Overly robust binding interactions may restrict surface mobility of redox-active intermediates, consequently impairing charge transfer efficiency. Such kinetic constraints during sulfur redox cycling constitute a fundamental performance limitation in Li-S battery systems.^54,55 The SRR pathways on M₂CT₂ surfaces are systematically mapped through density functional theory (DFT) calculations, with the thermodynamic Gibbs free energy change (ΔG) for each reaction intermediate from S₈ to Li₂S quantitatively presented in Fig. 8 and Table S3.† The calculation route of the SRR refers to the previous reports.^51,52 Obviously, the steps from S₈ to Li₂S₈ are all spontaneous exothermal on all the calculated M₂CT₂, which indicates that the S₈ can be readily converted into Li₂S₈. All reduction steps from Li₂S₈ to Li₂S are found to be endothermic, with the sole exception of the Li₂S₆ → Li₂S₄ step. The rate-limiting step is identified as Li₂S₄ → Li₂S₂ for V₂CO₂, Mo₂CF₂, and Mo₂CO₂, whereas for V₂CF₂, Cr₂CF₂, Cr₂CO₂, Mn₂CF₂, and Mn₂CO₂, the Li₂S₂ → Li₂S step served as the kinetic bottleneck. The endergonic Li₂S₂ → Li₂S transition stems from the inherent kinetic limitations of solid-state conversion processes, aligning with established theoretical frameworks in sulfur electrochemistry.^56,57 The Gibbs free energy change barrier (ΔG_barrier) for the rate-limiting step follow the sequence: Mo₂CF₂ (0.51 eV) < Mn₂CF₂ (0.78 eV) < V₂CF₂ (0.84 eV) < Cr₂CF₂ (0.88 eV) < V₂CO₂ (0.89 eV) < Mn₂CO₂ (1.23 eV) < Mo₂CO₂ (1.25 eV) < Cr₂CO₂ (1.37 eV). It is easy to find that the M₂CF₂ have lower rate-limiting step reaction energy barrier than those of M₂CO₂, which proves that the F atom greatly enhances the catalytic activity of the SRR.

Fig. 8 The calculation of Gibbs free energy change (ΔG) of SRR on MXene monolayers.

The elevated dissociation energy barrier of Li₂S, the terminal discharge product, significantly impedes electrochemical cycling reversibility.^58,59 This necessitates efficient delithiation kinetics (Li₂S → LiS + Li⁺ + e⁻) to sustain charging efficiency. Fig. 9 quantitatively delineates the atomic-scale decomposition pathways of Li₂S on M₂CT₂ monolayers, revealing the critical decomposition energy barriers. The minimum decomposition barriers of Li₂S are 0.27, 0.32, 0.51, and 0.53 eV on Mo₂CO₂, V₂CO₂, Cr₂CO₂ and Mn₂CO₂, respectively, while they are 0.74, 0.80, 0.90, and 1.00 eV on Mn₂CF₂, Mo₂CF₂, V₂CF₂ and Cr₂CF₂, respectively. Similar to the decomposition barrier of Li₂S on Ti₃C₂F₂ (0.90 eV) and Ti₃C₂O₂ (0.41 eV),⁶⁰ all the M₂CT₂ greatly reduce the decomposition barriers, which indicate great catalytic effect for the decomposition reaction. Studies attribute this phenomenon to MXenes' metal-to-carbon (M/C) ratio, where elevated transition metal content (M) enhances Li₂S dissociation kinetics through optimized d-band electronic interactions.²⁷ The results indicate that the M₂CT₂ could promote the decomposition of Li₂S, which enhances the electrochemical kinetics performance of Li-S batteries.

Fig. 9 Decomposition energies and decomposition paths of Li₂S on (a) V₂CF₂, (b) Cr₂CF₂, (c) Mn₂CF₂, (d) Mo₂CF₂, (e) V₂CO₂, (f) Cr₂CO₂, (g) Mn₂CO₂, and (h) Mo₂CO₂ monolayer.

The charging mechanism initiates with the delithiation of the terminal discharge product Li₂S, where ion migration dynamics critically govern decomposition and nucleation processes. Of particular significance is the interfacial Li⁺ diffusion kinetics across the cathode substrate, prompting systematic analysis of Li⁺ migration energy barriers. Therefore, the activation energy barriers for Li⁺ diffusion across the cathode are quantified through NEB simulations. Three kinds of the migration paths of Li⁺ are considered for Fcc and Hcp structure as presented in Fig. S5 (a) and (b),† respectively, and the diffusion barriers with different migration paths are shown in Fig. 10. It is found that the path 3 (the path 6) is the most difficult migration path for Fcc (Hcp) structure, while the diffusion barriers of the path 1 (the path 4) and the path 2 (the path 5) are comparable for all the calculated M₂CT₂. As shown in Fig. S5,† on V₂CT₂, Cr₂CF₂, Mn₂CT₂ and Mo₂CF₂, the Li⁺ tend to move along the path 2 (C1 → M → C2), while on Cr₂CO₂ and Mo₂CO₂, the Li⁺ prefers to be along the path 5 (H1 →C→ H2). The smallest diffusion barriers of Li⁺ on the different M₂CT₂ follow the sequence: Mo₂CO₂ (0.11 eV) < Cr₂CO₂ (0.12 eV) < V₂CO₂ (0.13 eV) < Mn₂CF₂ (0.16 eV) < Mn₂CO₂ (0.25 eV) < V₂CF₂ (0.32 eV) < Cr₂CF₂ (0.33 eV) < Mo₂CF₂ (0.44 eV). These relatively low diffusion barriers of Li⁺ promise the rapid diffusion of Li⁺ on the M₂CT₂ surface and conversion of LiPSs during the charge process.

Fig. 10 Diffusion barriers with different migration paths of Li⁺ on (a) V₂CF₂, (b) Cr₂CF₂, (c) Mn₂CF₂, (d) Mo₂CF₂, (e) V₂CO₂, (f) Cr₂CO₂, (g) Mn₂CO₂, and (h) Mo₂CO₂ monolayer.

The decomposition of Li₂S, the diffusion of Li⁺, and the catalytic activity of the SRR of the studied M₂CT₂ are compared with that of the available F/O-functionalized MXenes as listed in Table S4.† (ref. 27 and 61) It is found that the decomposition barriers of Li₂S for M₂CT₂ (0.27–1.00 eV) are apparently lower than those of V₂NF₂ (2.31 eV), V₂NO₂ (1.55 eV), and graphene (1.81 eV). Regarding the migration dynamics of Li⁺, the diffusion barriers of Li⁺ on the M₂CT₂ substrate (0.11 eV–0.44 eV) are close to those of V₂NF₂ (0.17 eV), V₂NO₂ (0.21 eV), and graphene (0.30 eV). Thus, M₂CT₂ show great kinetic performance during the discharge/charging process with relatively small decomposition barriers of Li₂S and diffusion barriers of Li⁺. What's more, the ΔG_barrier of M₂CT₂ (0.51 eV–1.37 eV) are significantly lower than those of V₂NF₂ (2.14 eV) and V₂NO₂ (1.88 eV), while they are comparable or lower than that of graphene (1.07 eV). Considering the above factors, V₂CO₂ and Mo₂CO₂ have the best electrocatalytic performance among the studied M₂CT₂, in which the decomposition barriers of Li₂S of V₂CO₂ and Mo₂CO₂ are only 0.32 eV and 0.27 eV, respectively, and the diffusion barriers of Li⁺ of V₂CO₂ and Mo₂CO₂ are 0.13 eV and 0.11 eV, respectively. Moreover, the V₂CO₂ and Mo₂CO₂ are metallic, which is beneficial for a fast charge/discharge rate. Therefore, V₂CO₂ and Mo₂CO₂ are the best choices for Li-S battery cathode material in the studied M₂CT₂.

The structural stability of electrode materials critically determines their cycling lifespan within operational voltage windows. To further investigate the stability as cathode materials for Li-S batteries, we calculate the energy of Li₂S adsorption on V₂CO₂ and Mo₂CO₂ through AIMD simulations at room temperature (300 K). As illustrated in Fig. S6,† the total energy fluctuations progressively diminish during simulation, while the structural frameworks remain intact throughout the process. Therefore, it is demonstrated that both V₂CO₂ and Mo₂CO₂ maintain excellent thermodynamic stability with Li₂S adsorption, and further validated their promise as highly prospective host materials for Li-S batteries.

Conclusion

In summary, we systematically explore potential of M₂CT₂ (M = V, Cr, Mn, and Mo; T = F and O) as anchoring materials and catalysts for Li-S batteries using DFT spin-polarized calculations. A moderate E_ads for Li₂S_n on M₂CT₂ species are observed, which would restrain the shuttling of LiPSs and prevent capacity fading and enhance the cycling capability for Li-S batteries. The decomposition barriers of Li₂S (0.27 to 1.00 eV) and the Li⁺ diffusion barriers of M₂CT₂ (0.11 to 0.44 eV) are relatively low, which would accelerate the formation and decomposition of solid Li₂S and hence redox reaction kinetics. It is also found that the M₂CO₂ exhibit lower decomposition barriers of Li₂S and diffusion barriers of Li⁺ than M₂CF₂, especially for V₂CO₂ and Mo₂CO₂. Moreover, the structures of Li₂S adsorbed V₂CO₂ and Mo₂CO₂ exhibit good thermodynamic stability at room temperature. All studied M₂CT₂ exhibit relatively low ΔG_barrier during the reaction process (0.51 to 1.37 eV), which would promote the conversion of LiPSs. Generally, the studied oxygen-terminated MXenes exhibit superior electrochemical performance compared to their fluorine-terminated MXenes. V₂CO₂ and Mo₂CO₂ are identified as the most promising candidates for Li-S battery host materials in the explored M₂CT₂. Our study not only provides critical insights for advancing M₂CT₂ implementation in Li-S battery cathodes, but also expands further theoretical investigations into MXene-based materials.

Data availability

The data that supports the findings of this study are available from the corresponding author upon reasonable request.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

This work is supported by the National Natural Science Foundation of China (No. 11904409 and 51572296), Youth Innovation Technology Project of Higher School in Shandong Province (2022KJ139), Postdoctoral Research Foundation of China (No. 2018M642721), and Shandong Postdoctoral Funded Project (201901012).

References

1. B. Ding, J. Wang, Z. Fan, S. Chen, Q. Lin, X. Lu, H. Dou, A. Kumar Nanjundan, G. Yushin, X. Zhang and Y. Yamauchi, Mater. Today, 2020, 40, 114–131.
2. W. Jin, X. Zhang, M. Liu, Y. Zhao and P. Zhang, Energy Storage Mater., 2024, 67, 103223.
3. R. Kumar, J. Liu, J.-Y. Hwang and Y.-K. Sun, J. Mater. Chem. A, 2018, 6, 11582–11605.
4. A. D. Pathak, E. Cha and W. Choi, Energy Storage Mater., 2024, 72, 103711.
5. S. C. Kim, X. Gao, S. L. Liao, H. Su, Y. Chen, W. Zhang, L. C. Greenburg, J. A. Pan, X. Zheng, Y. Ye, M. S. Kim, P. Sayavong, A. Brest, J. Qin, Z. Bao and Y. Cui, Nat. Commun., 2024, 15, 1268.
6. Z. Lv, P. Wang, J. Wang, S. Tian and T. Yi, J. Ind. Eng. Chem., 2023, 124, 68–88.
7. J. Li, Y. Qu, C. Chen, X. Zhang and M. Shao, Nanoscale, 2021, 13, 15–35.
8. J. T. Kim, A. Rao, H. Y. Nie, Y. Hu, W. Li, F. Zhao, S. Deng, X. Hao, J. Fu, J. Luo, H. Duan, C. Wang, C. V. Singh and X. Sun, Nat. Commun., 2023, 14, 6404.
9. Y. Chen, T. Wang, H. Tian, D. Su, Q. Zhang and G. Wang, Adv. Mater., 2021, 33, 2003666.
10. R. B. Nuwayhid, J. Yeom, H. O. Ford, Z. G. Neale, M. W. Swift, N. Bernstein, R. Carter and J. W. Long, RSC Appl. Interfaces, 2025, 2, 472–483.
11. Z. Wu, M. Liu, W. He, T. Guo, W. Tong, E. Kan, X. Ouyang, F. Qiao, J. Wang, X. Sun, X. Wang, J. Zhu, A. Coskun and Y. Fu, Nat. Commun., 2024, 15, 9535.
12. J. Chai, J. Du, Q. Li, N. Han, W. Zhang and B. Tang, Energy Fuel., 2021, 35, 15455–15471.
13. S. Li, D. Leng, W. Li, L. Qie, Z. Dong, Z. Cheng and Z. Fan, Energy Storage Mater., 2020, 27, 279–296.
14. Y. Pan, Y. Zhu, Y. Li, H. Liu, Y. Cong, Q. Li and M. Wu, Appl. Surf. Sci., 2023, 610, 155507.
15. M. Naguib, M. Kurtoglu, V. Presser, J. Lu, J. Niu, M. Heon, L. Hultman, Y. Gogotsi and M. W. Barsoum, Adv. Mater., 2011, 23, 4248–4253.
16. W. Meng, X. Liu, H. Song, Y. Xie, X. Shi, M. Dargusch, Z.-G. Chen, Z. Tang and S. Lu, Nano Today, 2021, 40, 101273.
17. X. Liang, A. Garsuch and L. F. Nazar, Angew. Chem., Int. Ed., 2015, 54, 3907–3911.
18. X. Zhao, M. Liu, Y. Chen, B. Hou, N. Zhang, B. Chen, N. Yang, K. Chen, J. Li and L. An, J. Mater. Chem. A, 2015, 3, 7870–7876.
19. W. Cui, Z.-Y. Hu, R. R. Unocic, G. Van Tendeloo and X. Sang, Chin. Chem. Lett., 2021, 32, 339–344.
20. Q. Yang, S. J. Eder, A. Martini and P. G. Grützmacher, npj Mater. Degrad., 2023, 7, 6.
21. D. Rao, L. Zhang, Y. Wang, Z. Meng, X. Qian, J. Liu, X. Shen, G. Qiao and R. Lu, J. Phys. Chem. C, 2017, 121, 11047–11054.
22. J. Hu, B. Xu, C. Ouyang, S. A. Yang and Y. Yao, J. Phys. Chem. C, 2014, 118, 24274–24281.
23. X. Liu, X. Shao, F. Li and M. Zhao, Appl. Surf. Sci., 2018, 455, 522–526.
24. H. Lin, D.-D. Yang, N. Lou, S.-G. Zhu and H.-Z. Li, Ceram. Int., 2019, 45, 1588–1594.
25. C. Shi, X. Zhang, Z. Li, T. T. Beyene, T. Zheng, Y. Liu and K. Zhu, Energy Fuels., 2024, 38, 14866–14890.
26. Q. Zhang, X. Zhang, Y. Xiao, C. Li, H. H. Tan, J. Liu and Y. Wu, ACS Omega, 2020, 5, 29272–29283.
27. D. Wang, F. Li, R. Lian, J. Xu, D. Kan, Y. Liu, G. Chen, Y. Gogotsi and Y. Wei, ACS Nano, 2019, 13, 11078–11086.
28. C. Lamiel, I. Hussain, J. H. Warner and K. Zhang, Mater. Today, 2023, 63, 313–338.
29. L. Zhang, J. Shi, K. Niu, P. Jia, Y. Gao and G. Gao, ACS Appl. Nano Mater., 2023, 6, 20812–20822.
30. G. F. l. Kresse, Phys. Rev. B:Condens. Matter Mater. Phys., 1996, 54, 11169–11186.
31. J. P. B. Perdew and K. Burke, Phys. Rev. Lett., 1996, 77, 3865–3868.
32. G. F. l. Kresse, Comput. Mater. Sci., 1996, 6, 15–50.
33. V. I. Anisimov and a. O. K. A. Jan Zaanen, Phys. Rev. B:Condens. Matter Mater. Phys., 1991, 44, 943–954.
34. S. Bae, Y. G. Kang, M. Khazaei, K. Ohno, Y. H. Kim, M. J. Han, K. J. Chang and H. Raebiger, Mater. Today Adv., 2021, 9, 100118.
35. Q. Sun, Z. Fu and Z. Yang, J. Magn. Magn. Mater., 2020, 514, 167141.
36. X. Zhang, W. Meng, T. He, L. Jin, X. Dai and G. Liu, Appl. Surf. Sci., 2020, 503, 144091.
37. S. Grimme, J. Antony, S. Ehrlich and H. Krieg, J. Chem. Phys., 2010, 132, 154104.
38. G. U. Henkelman, B. P. Uberuaga and H. Jónsson, J. Chem. Phys., 1991, 113, 9901–9904.
39. M. Khazaei, M. Arai, T. Sasaki, C. Y. Chung, N. S. Venkataramanan, M. Estili, Y. Sakka and Y. Kawazoe, Adv. Funct. Mater., 2012, 23, 2185–2192.
40. P. Gao, M. Song, X. Wang, Q. Liu, S. He, Y. Su and P. Qian, Nanomaterials, 2022, 12, 556.
41. J. He, P. Lyu and P. Nachtigall, J. Mater. Chem. C, 2016, 4, 11143–11149.
42. Q. Sun, Z. Fu, Y. Li and Z. Yang, J. Alloys Compd., 2021, 850, 156769.
43. Z. Tan, Z. Fang, B. Li and Y. Yang, ACS Omega, 2020, 5, 25848–25853.
44. L. Zhang, Y. Liu, M. Wu and G. Gao, Adv. Funct. Mater., 2024, 2417857.
45. L. Zhang, Y. Liu, Z. Xu and G. Gao, 2D Mater., 2023, 10, 045005.
46. C. Zhang, W. Chu, X. Hong, Q. He, R. Lu, X. Liao and Y. Zhao, Chem. Eng. J., 2022, 439, 135679.
47. G. Xu, X. Song, M. Jiang, R. Wang, S. Lian and X. Yang, Appl. Catal., B, 2025, 362, 124707.
48. Z. Chen, S. Huang, X. Yuan, X. Gan and N. Zhou, Appl. Surf. Sci., 2021, 544, 148861.
49. Z. Chen, S. Huang, B. Huang, M. Wan and N. Zhou, Appl. Surf. Sci., 2020, 509, 145319.
50. Z. Chen, Z. Chang, Z. Liu and N. Zhou, Appl. Surf. Sci., 2022, 602, 154375.
51. X. Zhu, M. Ge, T. Sun, X. Yuan and Y. Li, J. Phys. Chem. Lett., 2023, 14, 2215–2221.
52. C. Wei, T. Fang, X. Tang, P. Wang and X. Liu, J. Phys. Chem. C, 2022, 126, 17066–17075.
53. M. Cheviri and S. Lakshmipathi, Comput. Theor. Chem., 2021, 1202, 113323.
54. Z. Du, X. Chen, W. Hu, C. Chuang, S. Xie, A. Hu, W. Yan, X. Kong, X. Wu, H. Ji and L. J. Wan, J. Am. Chem. Soc., 2019, 141, 3977–3985.
55. G. Zhou, S. Zhao, T. Wang, S. Z. Yang, B. Johannessen, H. Chen, C. Liu, Y. Ye, Y. Wu, Y. Peng, C. Liu, S. P. Jiang, Q. Zhang and Y. Cui, Nano Lett., 2020, 20, 1252–1261.
56. Z. Yi, F. Su, L. Dai, Z. Wang, L. Xie, Z. Zuo, X. Chen, Y. Liu and C.-m. Chen, Energy Storage Mater., 2022, 47, 327–335.
57. Z. Jin, T. Lin, H. Jia, B. Liu, Q. Zhang, L. Li, L. Zhang, Z. M. Su and C. Wang, ACS Nano, 2021, 15, 7318–7327.
58. Y. Fu, H. Wang, D. Li, J. Shen, Y. Liu, M. Wei and Q. Hu, J. Alloys Compd., 2023, 965, 171306.
59. X. T. Fang, L. Zhou, C. Chen, D. L. Danilov, F. Qiao, H. Li and P. H. L. Notten, Molecules, 2023, 28, 7304.
60. H. Hong, N. A. R. Che Mohamad, K. Chae, F. Marques Mota and D. H. Kim, J. Mater. Chem. A, 2021, 9, 10012–10038.
61. K. Fan, Y. Ying, X. Luo and H. Huang, J. Mater. Chem. A, 2021, 9, 25391–25398.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d5ra01387a

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