A General Atomic Surface Modification Strategy for Improving

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Article

A General Atomic Surface Modification Strategy for Improving Anchoring and Electrocatalysis Behavior of TiCT MXene in Lithium-Sulfur Batteries 3

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Dashuai Wang, Fei Li, Ruqian Lian, Jing Xu, Dongxiao Kan, Yanhui Liu, Gang Chen, Yury Gogotsi, and Yingjin Wei ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.9b03412 • Publication Date (Web): 30 Aug 2019 Downloaded from pubs.acs.org on August 30, 2019

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A General Atomic Surface Modification Strategy for Improving Anchoring and Electrocatalysis Behavior of Ti3C2T2 MXene in Lithium-Sulfur Batteries Dashuai Wang†, Fei Li‡, Ruqian Lian†, Jing Xu†#, Dongxiao Kan†, Yanhui Liu#, Gang Chen†*, Yury Gogotsi†§*, Yingjin Wei†*

† Key

Laboratory of Physics and Technology for Advanced Batteries (Ministry of Education),

Jilin University, Changchun 130012, China § Department

of Materials Science & Engineering, and A. J. Drexel Nanomaterials Institute,

Drexel University, Philadelphia, Pennsylvania 19104, United States # Department

‡ Key

of Physics, College of Science, Yanbian University, Yanji 133002, China

Laboratory for UV Light-Emitting Materials and Technology (Ministry of Education),

Northeast Normal University, Changchun 130024, China *E-mail: [email protected]; [email protected]; [email protected]

ACS Paragon Plus Environment

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ABSTRACT

Multiple negative factors, including the poor electronic conductivity of sulfur, dissolution and shuttling of lithium polysulfides (Li2Sn), and sluggish decomposition of solid Li2S, seriously hinder practical applications of lithium-sulfur (Li-S) batteries. To solve these problems, a general strategy was proposed for enhancing the electrochemical performance of Li-S batteries using surface-functionalized Ti3C2 MXenes. Functionalized Ti3C2T2 (T = N, O, F, S, and Cl) showed metallic conductivity, as bare Ti3C2. Among all Ti3C2T2 investigated, Ti3C2S2, Ti3C2O2, and Ti3C2N2 offered moderate adsorption strength, which effectively suppressed Li2Sn dissolution and shuttling. This Ti3C2T2 exhibited effective electrocatalytic ability for Li2S decomposition. The Li2S decomposition barrier was significantly decreased from 3.390 to ~0.4 eV using Ti3C2S2 and Ti3C2O2, with fast Li+ diffusivity. Based on these results, O and S terminated Ti3C2 were suggested as promising host materials for S cathodes. In addition, appropriate functional group vacancies could further promote anchoring and catalytic abilities of Ti3C2T2 to boost the electrochemical performance of Li-S batteries. Moreover, the advantages of Ti3C2T2 host material could be well retained even at high S loading, suggesting the potential of surface-modified MXene for confining sulfur in Li-S battery cathodes.

KEYWORDS: Li-S battery; MXene; Ti3C2; surface modification; first-principles calculation


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The increasing demand for portable electronics and electric vehicles has driven the development of high-performance rechargeable batteries. In various rechargeable battery systems, lithium-sulfur (Li-S) has attracted increasing attention due to its five-fold energy density (2600 W∙h∙kg-1 or 2800 W∙h∙L-1) with respect to traditional lithium ion batteries (LIBs).

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sulfur is abundant, inexpensive, and has low toxicity. Although significant advances have been accomplished in recent years, many scientific and technical challenges still hinder the practical application of Li-S batteries. S cathodes experience multiple reduction/oxidation redox during cycling, consequently forming a series of Li2Sn intermediate products. The insulating property of S, its noticeable volume change (80%) upon cycling as well as the large decomposition energy of Li2S causes sluggish reaction kinetics, resulting in large electrode polarization and poor rate capability. In addition, high-ordered lithium polysulfides (LiPSs; Li2Sn, n = 4, 6, or 8) formed during discharge are soluble in aprotic electrolytes. A portion of soluble LiPSs shuttle to the Li metal anode and deposit there, resulting in rapid capacity decay. 2,4–8 Four aspects have been considered to address the aforementioned issues of Li-S batteries, including (1) conductive matrix, such as carbon materials9–13 and metal nanoparticles,14,15 to facilitate electron transport; (2) confinement framework, such as carbon materials16,17 and transition metal nitrides9,10/sulfides18–23/carbide6,11,24–26/oxides/,1,7,27,28 to mitigate the S volume expansion during lithiation and suppress LiPSs dissolution; (3) blocking layer, such as functional interlayers29 and separators,30,31 to suppress LiPSs shuttling; and finally, (4) electrocatalysts, such as transition metal sulfides,8 to improve Li2S decomposition kinetics. Two-dimensional (2D) materials, such as graphene and transition metal dichalcogenides, have been studied as host materials for S cathodes and effectively improve the electrochemical performance of Li-S batteries. Besides these materials, 2D transition metal carbides and nitrides with a general formula


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of MnCn-1/MnNn-1 (n = 2–4), termed MXenes, where “M” represents an early transition metal, such as Ti, V, Nb, and Mo, have drawn particular attention because of their fascinating electronic conductivity, high surface area with abundant active sites, controllable hierarchical structure with high S loading, and physical confinement for electrode volume changes. Liang et al. have shown that a S/Ti2C composite with 70 wt% S loading exhibits a specific capacity of ~1200 mA∙h∙g-1 at C/5 rate and stable long-term cycling (80% capacity retention over 400 cycles at C/2 rate) due to the strong LiPSs interactions with surface Ti atoms.25 In addition, Ti3C2/Ti3CN nanosheet/carbon nanotube composites enable S cathodes with excellent cycling stability up to 1200 cycles with a low capacity fading rate of 0.043% per cycle.11 Almost all synthesized MXenes are terminated with large quantities of functional groups, including O, OH, or F. These functional groups play important roles in electrochemical energy storage and electrocatalytic properties. For example, compared with F and OH, O-functionalized MXenes have exhibited larger capacities in LIBs, while MXenes terminated with O have shown lower activity in hydrogen evolution reactions due to strong interactions between H and surface O.32–34 Recently, non-native functional groups, such as P, Si, S, and Cl, have been introduced to MXenes, which produces positive effects in electrochemical energy storage, such as in LIBs and supercapacitors.35–38 MXenes have also achieved considerably improved performance in Li-S batteries. However, synergetic effects of MXenes in Li-S batteries, including the anchoring ability for LiPSs as well as the catalytic effect for Li2S decomposition, have not been clearly established, which has prompted study at the atomic level. In this study, Ti3C2 was investigated as the host material for S cathodes. The structural stability, electronic conductivity, and anchoring strength for S and Li2Sn, and the catalytic ability for Li2S decomposition were comprehensively studied by first-principles calculations. The anchoring strength and catalytic ability of Ti3C2 was modulated


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by introducing different surface terminating groups (T), including N, P, O, S, F, and Cl, onto Ti3C2 to screen for the best functionalized Ti3C2T2 as an S cathode host in Li-S batteries.

RESULTS/DISCUSSION

Fig. 1 (a) Side and top views of Ti3C2; (b) density of states of Ti3C2; (c) Li2Sn and S8 structures. Why are surface terminating groups needed? As one of the most common MXene materials, Ti3C2 was constructed by sandwiched Ti and C atom layers stacked in a sequence of TiC-Ti-C-Ti, which can be described as three Ti layers being interleaved with two C layers, forming edge-shared Ti6C octahedral structures (Fig. 1a). Calculated electronic structures showed that Ti3C2 had metallic conductivity (Fig. 1b). The anchoring effect of Ti3C2 was evaluated by calculating adsorption energies (𝐸𝑎𝑑) for S8 and Li2Sn according to the equation, 𝐸𝑎𝑑 = 𝐸𝑀𝑋𝑒𝑛𝑒 + 𝐸𝐿𝑖2𝑆𝑛 ― 𝐸𝑀𝑋𝑒𝑛𝑒 + 𝐿𝑖2𝑆𝑛

(1)

where 𝐸𝑀𝑋𝑒𝑛𝑒 + 𝐿𝑖2𝑆𝑛 and 𝐸𝑀𝑋𝑒𝑛𝑒 are the total energies of Ti3C2 with/without Li2Sn, respectively, and 𝐸𝐿𝑖2𝑆𝑛 the total energy of Li2Sn. A positive 𝐸𝑎𝑑 indicated favored adsorption. The 𝐸𝑎𝑑 for S8 was calculated in the same way. The lowest energy configurations of S8 and Li2Sn were obtained


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using the CALYPSO code (Fig. 1c) and the results showed that the adsorption strength of S8 and Li2Sn on bare Ti3C2 was very strong, indicated by the large 𝐸𝑎𝑑 values in the range of 5–27 eV (Fig. 2a). Under such a large adsorption strength, the S8 ring decomposed into single S atoms, which were then tightly captured by the outer Ti layer of Ti3C2 (Fig. 2b). A similar phenomenon also happened for Li2Sn, where almost all S atoms in Li2Sn were offloaded onto Ti3C2 surfaces. According to these results, it is expected that bare Ti3C2 would be fully terminated by S during preparation of S/Ti3C2 or during initial electrochemical cycling. In other words, bare Ti3C2 cannot be directly used in Li-S batteries. Therefore, in order to achieve reversible Li-S battery operation, surface modification of Ti3C2 was required to obtain suitable adsorption strength with S and Li2Sn.

Fig.2 (a) Adsorption energies of S8 and Li2Sn on Ti3C2; adsorption structures of (b) S8, (c) Li2S8 and (d) Li2S6 on Ti3C2 before (left) and after (right) optimization.


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Structure and electronic properties of Ti3C2T2. Based on the above discussion, six surface terminating groups, T = N, P, O, S, F, and Cl, were considered for Ti3C2. The favored sites for terminating groups was determined by performing total energy calculations for T-saturated Ti3C2 monolayers, termed Ti3C2T2. There were three possible positions for T, i.e., on top of the Ti(1) atom (site α), on top of the C atom (site β), and on top of the Ti(2) atom (site γ, Fig. 1a). Three possible configurations were considered for these terminating groups (Fig. S1, Supporting Information), and the total energies of all possible Ti3C2T2 phases listed in Table S1. In Mode I, T groups were at site α on both sides of the Ti3C2 layer; in Mode II, at site β on both sides of the monolayer; and in Mode III, at site γ on both sides of the monolayer. For the Ti3C2T2 (T= P, O, S, F, and Cl) phases, Mode III had the lowest total energy, while Mode II was the lowest energy phase for Ti3C2N2. The binding energies of terminated atoms on Ti3C2 were calculated to verify the thermodynamic stability of function groups, as listed in Table S2. As a result, all of the terminated atoms could be energetically stable to bind with Ti3C2. In addition, thermal stability of Ti3C2T2 was examined by performing AIMD simulations at 298, 800, and 1000 K. All functionalized Ti3C2T2 maintained their structural integrity after 5 ps of stimulation except Ti3C2P2 (Fig. S2). For Ti3C2P2, although the Ti3C2 part retained structural integrity at 298 K, P atoms bonded to each other in some disorder and the distorted structure could not recover the initial configuration after structural relaxation. One possible reason could be P has the maximum atomic radius (110.5 pm) among all the species of atoms considered, which more easily bond with each other on the limited Ti3C2 surface. What’s more, phonon dispersions along the high symmetrical directions of the Brillouin zone for Ti3C2T2 were calculated to determine the lattice dynamic stabilities, as shown in Fig. S3. Except for Ti3C2P2, no imaginary frequency phonon was found for Ti3C2T2 (T=N/O/S/F/Cl), indicating the lattice dynamic stability of Ti3C2T2 with a variety of


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practically relevant terminations. The results were consistent with AIMD simulations. Consequently, P functional groups were not included in the following discussion. Hence, Ti3C2 could be feasibly functionalized by N, O, S, F, or Cl groups, in which O and F are native functional groups born from material preparation and the others can be subsequently introduced with appropriate experimental control.

Fig. 3 Density of states of Ti3C2T2 (T= N, O, S, F, Cl), respectively. Calculated electron density of states (DOS) showed that, except for ferromagnetic Ti3C2N2, all of functionalized Ti3C2T2 were nonmagnetic by considering spin polarization. For all Ti3C2T2 phases, the majority of electron states at the Fermi level were from Ti-3d states (Fig. 3). As a result, the metallic conductivity of Ti3C2 was well retained after surface functionalization. The excellent electronic conductivity of Ti3C2T2 could cure S’s low conductivity and thus improve the


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electrochemical performance of Li-S batteries. The obvious hybridization between Ti-3d and T(N/O/F)-2p/T(S/Cl)-3p indicated a strong (partial) polar covalent Ti-T bonding.

Fig. 4 (a) Adsorption energies, (b) ratio of vdW interaction of S8 and Li2Sn on Ti3C2T2. Charge density difference of (c, e, g, i, k) S8 and (d, f, h, j, l) Li2S6 on Ti3C2T2 (T = N, O, S, F, Cl, from top to bottom). The isosurface level is set at 0.0003 and 0.003 e Å-3 for S8 and Li2S6 adsorption, respectively. Anchoring ability for Li2Sn. DFT calculations were performed to study the anchoring effects of Ti3C2T2 (T = N, O, S, F, or Cl) for Li-S batteries. The optimized adsorption conformations of Li2Sn and S8 on Ti3C2T2 are shown in Fig. S4. The adsorption strength between Ti3C2T2 and Li2Sn (or S8) was evaluated by 𝐸𝑎𝑑 based on equation (1). The calculated 𝐸𝑎𝑑 with vdWs at different


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lithiation degrees reflected a combination of chemical and physical (vdW) interactions between Ti3C2T2 and Li2Sn (or S8, Fig. 4a). For all Ti3C2T2 phases, the 𝐸𝑎𝑑 for S8 was smaller than that for Li2Sn. The calculated 𝐸𝑎𝑑, 0.75−1.20 eV, was much closer to that of S8 on graphene, transitional metal oxides, and sulfides.39 With S lithiation, different functional groups showed distinct anchoring effects, which were divided into three classes according to the 𝐸𝑎𝑑 magnitude. The adsorption strength for Ti3C2O2 and Ti3C2S2 (at 2.0–4.5 eV) were much stronger than the others during the whole lithiation process. The adsorption strengths for Ti3C2Cl2 and Ti3C2F2 were relatively weak, in the range of 0.8–1.4 eV. In comparison, the adsorption strength for Ti3C2N2 was in the middle of those of the five Ti3C2T2 phases. Consequently, the adsorption strength followed the sequence of Cl < F < N < O < S. For almost all Ti3C2T2 phases, the evolution of 𝐸𝑎𝑑 followed a similar trend, growing stronger with increasing lithiation and reaching a maximum for the fully lithiated Li2S product. The role for suppressing LiPSs shuttling was verified by calculating the 𝐸𝑎𝑑 and comparison with Li2Sn (n ≥ 4) binding strength on DOL/DME electrolyte solvent molecules (DOL, 1,3-dioxolane; DME, 1,2-dimethoxyethane, Fig. S5). Except for Ti3C2Cl2, all 𝐸𝑎𝑑 values were stronger than those for Li2Sn (n ≥ 4) bonded with DOL/DME (0.74– 0.79 eV), which indicated that Ti3C2T2 (T = N, O, S, or F) could be expected to effectively suppress LiPSs shuttling. However, based on the crystal structure model, DFT calculations cannot deal with the effect of electrolyte content on the adsorption and anchoring properties of Ti3C2T2 MXenes. The calculations of the binding effects of the two molecular groups (DOL and DME) with polysulfides could only qualitatively study the advantages of S fixation on Ti3C2T2 by a simplified method. Based on the above analysis, the Ti3C2T2 could fix S and Li2Sn on their surfaces and prevent the polysulfides dissolution and shuttling in the electrolyte solvent. These advantages,


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together with the intrinsic thermal and thermodynamic stabilities of the Ti3C2T2 host materials, may effectively improve the cycle stability of Li-S batteries resulting in prolonged cycle life. Based on increasing 𝐸𝑎𝑑 with lithiation, the anchoring effects of Ti3C2T2 was expected to mainly derive from chemical interactions between Ti3C2T2 and Li2Sn. Interactions between Ti3C2T2 and Li2Sn (or S8) were identified by determining the contributions of chemical and physical (vdW) interactions at different lithiation stages and then evaluated by comparison with adsorption energies without vdW functions (Fig. 4b). For all Ti3C2T2 phases, except Ti3C2N2, vdW interactions dominated S8 adsorption at nearly 100%. In other words, no chemical bonds formed between S atoms and Ti3C2T2 (T= O/S/F/Cl). With lithiation, the changing trend of vdW interactions was almost the opposite to that of 𝐸𝑎𝑑, i.e., vdW interactions gradually decreased with lithiation and reached a minimum value for Li2S. For relatively strongly terminated Ti3C2, such as with O/S/N, vdW interactions decreased linearly with lithiation and chemical interactions increased at the same time. For weakly terminated Ti3C2, such as with F and Cl, vdW interactions dominated the whole lithiation process, to >55% percentage. According to the above discussion, chemical interactions between Ti3C2T2 and Li2Sn (or S8) were mostly derived from intercalated Li+, while vdW interactions were mainly contributed by S. In assessing the nature of the interaction between Ti3C2T2 and Li2Sn (or S8), the charge density difference was analyzed by the equation ∆𝜌 = 𝜌𝑇𝑖3𝐶2𝑇2 + 𝐿𝑖2𝑆𝑛/𝑆8 ― 𝜌𝑇𝑖3𝐶2𝑇2 ― 𝜌𝐿𝑖2𝑆𝑛/𝑆8

(2)

Here, the adsorption conformations of S8 and Li2S6 on Ti3C2T2 (T = N, O, S, or Cl) were taken as examples (Figs. 4c-4l). When S8 was adsorbed on Ti3C2T2 surfaces, most charge transfer occurred inside S8 clusters and Ti3C2T2. This indicated that there was no chemical bonding between S8 and Ti3C2T2 and that the interactions between them mainly depended on vdW adsorption. Only for


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Ti3C2N2, a certain amount of charge transferred between S and Ti, forming Ti-S polar covalent bonds. Meanwhile, when Li2S6 was adsorbed on Ti3C2T2 (T = N, O, S), electron-rich (green) regions appeared between Li and T, indicating the formation of Li-T bonds. At the same time, Li-S and S-S bonds in Li2S6 were weakened with the increase of electron-deficient regions (blue) between Li (or S) and S, which might have contributed to Li2S6 decomposition. Table 1 Calculated Li2S, Li2S6 decomposition barriers and Li+ diffusion barriers on Ti3C2T2.

Substrate

Decomposition barrier (eV)

Li+ diffusion barrier (eV)

Li2S

Li2S6

Ti3C2O2

0.411

0.211

0.236

Ti3C2S2

0.351

0.180

0.188

Ti3C2N2

1.101

0.434

0.573

Ti3C2F2

0.903

1.208

0.187

Ti3C2Cl2

1.625

1.532

0.196

Catalytic and diffusion kinetic properties. The final discharge product of Li-S batteries, Li2S, suffered from low electronic conductivity, low Li+ diffusivity, and high decomposition energy, which resulted in high overpotential and low rate capability. Consequently, Li2S decomposition kinetics (Li-S bond breaking) and Li+ diffusion on the substrate were vital for the electrochemical performance of Li-S batteries. Deeper insight into the reaction mechanisms of Li2S decomposition was gained by analyzing Li2S decomposition kinetics on Ti3C2T2 using the CI-NEB method. The decomposition process was drawn from an original Li2S molecular into a LiS cluster and a single Li+, i.e., 𝐿𝑖2𝑆→𝐿𝑖𝑆 + 𝐿𝑖 + + 𝑒 ― . The results showed that the natural Li2S decomposition barrier was as large as 3.390 eV. All considered Ti3C2T2 could reduce the Li2S decomposition barrier effectively (Table 1, Fig. S6). Among them, Ti3C2Cl2 showed the largest decomposition barrier of ~1.5 eV, which could have been due to weak interactions between Li2S


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and Ti3C2Cl2, such that Li-S bond breaking was weakly affected by the substrate. The decomposition barriers on Ti3C2N2 and Ti3C2F2 were reduced to 1.0 eV. Furthermore, Ti3C2S2 and Ti3C2O2 showed much smaller decomposition barriers, at 0.351 and 0.411 eV respectively. These values were comparable with those reported for traditional metal disulfides (0.31–0.56 eV).8 Therefore, these two Ti3C2T2’s could efficiently catalyze Li2S decomposition and reduce the overpotential of Li-S batteries. The decomposition barrier on Ti3C2S2 was observed to be much smaller than that on Ti2CS2 (~1.51 eV).40 This indicated that the M/C ratio of MXenes might have influenced the catalytic ability for Li2S decomposition. The efficient catalysis of Ti3C2T2 for Li2S decomposition helped improve the electrode kinetics of Li-S batteries. Moreover, fast decomposition of Li2Sn (n = 4, 6, or 8) was also required because it could shorten the LiPSs accumulation time in the cathode and thus reduce LiPSs dissolution. Hence, the Li2S6 decomposition kinetics on Ti3C2T2 were further studied and the trend of the Li2S6 decomposition barrier was found to be like that of Li2S, which was reduced from 3.406 eV (for native Li2S6) to 0.180, 0.211, 0.434, 1.208, and 1.532 eV for Ti3C2S2, Ti3C2O2, Ti3C2N2, Ti3C2F2, and Ti3C2Cl2, respectively (Fig. S7). As Li+ diffusion on substrate materials is closely related to Li2Sn nucleation and decomposition, Li+ diffusivity on Ti3C2T2 was studied by calculating diffusion barriers. The energy profiles along the Li+ diffusion coordinate showed that the minimum Li+ diffusion barriers on Ti3C2S2, Ti3C2O2, Ti3C2F2, and Ti3C2Cl2 were 0.188, 0.236, 0.187, and 0.196 eV, respectively, which were much lower than that of Li+ on graphene (0.300 eV, Fig. S8). With one exception, the Li+ diffusion barrier on Ti3C2N2 was as large as 0.573 eV. This was because the valence electrons of N formed stronger Li-N ionic bonds, which limited Li+ diffusivity. To summarize, the kinetic


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properties of Ti3C2T2 in Li-S batteries, including catalytic effects and Li+ diffusion, followed the priority sequence of Ti3C2S2 > Ti3C2O2 > Ti3C2F2 > Ti3C2N2 > Ti3C2Cl2. Vacancy effect. MoS2 with sulfur deficiencies has been shown to possess better catalytic ability than stoichiometric MoS2 in Li-S batteries.41 To solve the poor adsorption strength and sluggish electrode kinetics of Ti3C2Cl2, a small portion of Cl deficiencies were produced to form Ti3C2Cl2-x (x = 1/16). Here, the vacancy formation energy of one Cl in 4×4 ×1 Ti3C2Cl2 supercell was calculated to be -2.866 eV, indicating Cl vacancy could form on Ti3C2Cl2 surface. Compared with intact Ti3C2Cl2, the 𝐸𝑎𝑑 of Li2Sn (S8) increased from 0.7-1.1 to 2.5-3.5 eV (Fig. 5a). The vdW interaction ratio was decreased to 25–33%, while chemical interactions increased (Fig. 5b). This could have been due to the Cl vacancy with an electron-rich region (Fig. 5c), which could be easily bonded with the S of Li2Sn (S8). In addition, the adsorption strength for Li2Sn was much stronger than the bonding energy with DOL/DME, which greatly suppressed the shuttling effect. In addition, decomposition barriers for Li2S and Li2S6 on Ti3C2Cl2-x declined dramatically, to 0.316 and 0.537 eV (Fig. S9), respectively, which were even lower than those on the optimal Ti3C2S2. Consequently, it was speculated that appropriate functional group vacancies would further promote the anchoring and catalytic capabilities of Ti3C2T2 to boost the electrochemical performance of Li-S batteries.


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Fig. 5 (a) Adsorption energies, (b) ratio of vdW interaction of S8 and Li2Sn on Ti3C2Cl2-x (x=1/16); electron localization functions of the (110) slice of (c) Ti3C2Cl2-x and (d) Ti3C2Cl2. Effects of S loading amount. According to the above analysis, Ti3C2T2, particularly with T = O and S, can improve the electrochemical properties of the sulfur cathode. However, it should be noted that the electrochemical performance of Li-S batteries is strongly affected by the S loading amount in the cathode. The structure model used in the above analysis, one S8 (or Li2Sn) molecular adsorbed on a 4×4×1 supercell, only corresponded to a small S loading (0.9~6.5 wt%) in the MXene host material, which does not reflect the real Li-S batteries. Hence, in order to evaluate the practical application of Ti3C2T2 in Li-S batteries, we further studied the influence of S loading on the anchoring effect, electronic and catalytic properties of Ti3C2S2, which has been predicted to possess the best performance among all Ti3C2T2 MXenes under investigated. First, we studied the anchoring ability of Ti3C2S2 monolayer for the S and Li2Sn electrodes with different S contents, i.e. S8 (35.7 wt%), Li2S6 (30.3 wt%), Li2S4 (32.6 wt%) and Li2S (14.7


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wt%). As shown in Table 2, all the adsorption energies for the above species (E1) were negative, indicating that the Ti3C2S2 host material could effectively anchor sulfur and suppress the dissolution of LiPSs. Especially, the adsorption energies (E2) became even more negative with further adsorption of an additional S8 (or Li2Sn) molecule on the Ti3C2S2 substrate. This indicated that the anchoring effect of Ti3C2S2 could be well retained under high S loading. On the above basis, we further calculated the electronic structures of the Ti3C2S2 host material with/without S and Li2Sn adsorption. The S8 molecule had an intrinsic insulator property (Fig. 6a), while the Ti3C2S2 host material was obviously metallic (Fig. 6b). The electronic structure of the Ti3C2S2/S8 composite exhibited sizeable electron states crossing the Fermi level, which indicated that the insulating property of pristine S8 has been successfully changed by the Ti3C2S2 host (Fig. 6c). In addition, all the Ti3C2S2/Li2Sn (n = 6, 4, 1) composites also showed metallic properties (Fig. 6df). This indicates that the electrode can maintain high electronic conductivity during whole lithiation/de-lithiation process, which is much helpful for the rate capability and cycle stabile of the Li-S battery cells. Finally, we studied the Li2S decomposition barriers on Ti3C2S2 surface with different S loading contents. CI-NEB calculations showed that the decomposition barrier of Li2S increased from 0.351 eV to 1.900 eV with the Li2S concentration increased from 1/16 (corresponding to S loading of 0.9 wt%) to 19/16 (corresponding to S loading of 14.7 wt%) per Ti3C2S2. However, all these decomposition barriers were much lower than the original decomposition barrier of 3.390 eV for bare Li2S (Fig. 6g). The results showed that the catalytic capacity of the Ti3C2S2 host material could be well maintained even at high S loading contents. Table 2. The adsorption energies (E1) for the S and Li2Sn electrodes with different S contents on the Ti3C2S2 monolayer. The adsorption energies (E2) with further adsorption of an additional S8 (or Li2Sn) molecular on the Ti3C2S2 monolayer.


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Adsorbates (S content, wt%)

E1 (eV)

E2 (eV)

S8 (35.7 %)

-0.388

-0.647

Li2S6 (30.3 %)

-1.362

-3.138

Li2S4 (32.6 %)

-2.133

-4.543

Li2S(14.7 %)

-5.103

-6.521

Fig. 6 Density of states for S8, Ti3C2S2, Ti3C2S2/S8, Ti3C2S2/Li2S6, Ti3C2S2/Li2S4, Ti3C2S2/Li2S (a, b, c, d, e, f, respectively). The Fermi level is set to zero. (g) Decomposition barriers of bare Li2S (black) and Li2S on the Ti3C2S2 supercell with 1/16 (green) and 19/16 (red) Li2S to Ti3C2S2 ratios.

CONCLUSIONS

Surface modification of 2D materials has attracted much attention in recent years, including surface functional groups, hybrid heterojunctions, defects and doping. As experienced by other 2D analogs, properties of MXene monolayers can also be improved by proper surface modification. However, although there have been many reports on the native O, F and OH termination of


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MXenes, very limit attention has been paid to functional surface modification of MXene monolayers. In this work, Ti3C2 MXenes were systematically studied as the S cathode host for Li-S batteries by first-principles calculations. It was shown that bare Ti3C2 cannot be directly used in Li-S batteries because of excessively strong interactions between Li2Sn/S8 and Ti3C2 induced decomposition of S8 and Li2Sn. Ti3C2 terminated with functional groups can achieve high performance in Li-S batteries. Among various functional groups, S and O were identified as the best choices for Ti3C2 surface modification, in which O was a native moiety derived from material preparation and S was introduced secondarily with appropriate experimental control. The improved Li-S battery performance benefited from using Ti3C2T2 as the host material due to: (1) metallic properties of Ti3C2T2 conferred high electronic conductivity to the S cathode, (2) the moderate surface adsorption strength effectively suppressed Li2Sn dissolution and shuttling, and (3) the electrocatalytic ability of Ti3C2T2 accelerated the electrode reaction kinetics for Li2S decomposition. These advantages can be well retained even at high S loading contents, supporting the practical application of Ti3C2T2 as a host material for Li-S batteries. Also, appropriate vacancies in surface terminations further promoted the anchoring and electrocatalytic ability of Ti3C2T2 to boost the electrochemical performance of Li-S batteries. This study provides a general strategy for enhancing the electrochemical properties of Li-S batteries by atomic surface modification of MXenes.

METHODS

First-principles calculations were performed within the framework of density functional theory (DFT), as implemented in the Vienna ab initio simulation package.42 Projector-augmented wave potential was used with a plane-wave cutoff energy of 600 eV.43 Exchange correlation energy


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was described by the generalized gradient approximation in the scheme proposed by PerdewBurke-Ernzerhof.44 Spin alignments were studied by spin-polarized calculations. For geometry optimization, the Brillouin-zone integration was performed using a Monkhorst-Pack grid of kpoint sampling.45 Considering the strong correlation in titanium (Ti), the electronic structure was calculated using a generalized gradient approximation (GGA) plus Hubbard U (GGA+U) method with a U value of 4.2 eV.46,47 Geometry optimizations were performed using the conjugated gradient method, with the convergence threshold set at 1×10−6 eV∙atom-1 in energy and 0.01 eV∙Å-1 in force. A vacuum distance of >20 Å was employed to avoid interactions of neighboring images. The configurations of S8 and Li2Sn (n= 1, 2, 4, 6, and 8) were predicted by advanced swarmintelligence structure prediction method implemented in the CALYPSO code.48,49 For all cases, structural searching simulations for each calculation were stopped after generating over 500 structures (e.g., ~20–30 generations). The van der Waals (vdW) forces between Ti3C2T2 and Li2Sn were accurately dealt with vdWinclusive DFT-D3 method, which introduced empirical dispersion corrections implemented by Grimme.50 The different sulfur amount configurations were based on a 4×4×1 supercell of Ti3C2T2. Furthermore, the climbing-image nudged elastic band (CI-NEB) method was applied for computing diffusion and decomposition barriers.51 The force convergence criterion for optimization was set at 0.01 eV∙Å-1. Ab initio molecular dynamics (AIMD) simulations were performed at different temperatures of 298, 800, and 1000 K. AIMD simulations in an NVT ensemble using the Nose-Hoover heat method lasted for 5 ps with a time step of 1.0 fs. The phonon spectra were calculated using the PHONOPY package.52

ASSOCIATED CONTENT


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The authors declare no competing financial interest.

AUTHOR INFORMATION

Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript.

ACKNOWLEDGMENTS

This work was supported by the Ministry of Science and Technology of China (No. 2015CB251103), the National Natural Science Foundation of China (No. 21773091), Science and Technology Department of Jilin Province (No. 20180414004GH), Interdisciplinary research fund of Jilin University (No. 10183201813) and the Independent Industrial Innovation Funding of Jilin University (No. 2018C008).

SUPPORTING INFORMATION

Supporting Information Available: Ti3C2T2 configurations and Snapshots of the equilibrium structures after AIMD simulation, Adsorption configurations of Li2Sn and S8 on Ti3C2T2, Configurations and binding energies of Li2Sn (n =4, 6, 8) bound with DME and DOL, Energy profiles for the decomposition of Li2S and Li2S6 in vacuum and on Ti3C2T2, Li+ diffusion on Ti3C2T2. This material is available free of charge via the Internet at http://pubs.acs.org.

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Table of contents


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A General Atomic Surface Modification Strategy for Improving

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