Research Article | Open Access
Xin Fan, Mike Tebyetekerwa, Yilan Wu, Rohit Ranganathan Gaddam, Xiu Song Zhao, "Magnesium/Lithium Hybrid Batteries Based on SnS2-MoS2 with Reversible Conversion Reactions", Energy Material Advances, vol. 2022, Article ID 9846797, 14 pages, 2022. https://doi.org/10.34133/2022/9846797
Magnesium/Lithium Hybrid Batteries Based on SnS2-MoS2 with Reversible Conversion Reactions
The magnesium/lithium hybrid batteries (MLHBs) featuring dendrite-less deposition with Mg anode and Li-storage cathode are a promising alternative to Li-ion batteries for large-scale energy storage. However, their limited energy density limits their practical implementation. To improve this, beyond the commonly proposed intercalation compounds, high-capacity conversion-type cathodes based on heterostructures of tin sulphide-molybdenum disulphide (SnS2-MoS2) are proposed in this work. Individual SnS2 is already a promising high-capacity electrode material for multivalent batteries and undergoes conversion reactions during the ion storage process. The introduction of S-deficient MoS2 enhances the reversibility of SnS2 in the conversion reaction via strong polysulfide anchoring and catalytic effect. Our results show that the SnS2-MoS2 electrode achieves a high charge capacity of ~600 mAh g-1 at 50 mA g-1 and an excellent rate capability of 240 mAh g-1 at 1000 mAh g-1 with a negligible capacity fading rate of 0.063% per cycle across 1000 cycles. The results highlight a new direction toward designing 2D heterostructures as high-capacity cathodes beyond intercalation-type cathodes for multivalent-ion batteries.
Lithium-ion batteries (LIBs) with high energy density and portability are now well-positioned to offer one of the most appealing options for future electric transportation and large-scale grid storage [1, 2]. However, lithium scarcity and the potential safety issues of LIBs impose restrictions on their penetration into the large-volume markets [3, 4]. In addition, LIB technology is approaching a fundamental limit in terms of energy density (slightly above 300 W h kg-1). For these reasons, research interests in battery systems are starting to pivot towards multivalent metal batteries involving magnesium, zinc, calcium, and aluminium [5–7]. Multivalent metals that undergo beyond one-electron redox reactions have great potential to provide higher volumetric energy densities than the monovalent LIBs . Rechargeable magnesium (Mg) batteries (MBs) are one of the potential contenders to replace current LIBs to power future electric vehicles with lower cost, higher safety, and extended mileage . Compared with the Li metal anode, the metallic Mg anode shows higher volumetric capacity, safety and easier processing features, more uniform and nondendritic electroplating during cell charging, and cheaper costs due to its greater abundance in the earth crust [5, 7–11]. They are gaining more widespread interest since the pioneering work of Aurbach et al. , which envisioned their possible breakthroughs with proper electrolyte and cathode materials. Unfortunately, the high charge density of divalent Mg2+ (120 C mm-3 of Mg2+vs. 54 C mm-3 of Li+) leads to sluggish solid-state diffusion kinetics in MBs. More still, the commonly used Li-storage electrode materials show low Mg2+ intercalation levels, significant cycling polarisation, and rapid capacity decay for common cathode materials . Hence, a suitable electrode material can possibly further the development of MB technologies.
On the other hand, utilising magnesium metal with less dendrite formation and the fast reaction kinetics of Li+ for the whole electrochemical redox processes, as in the case of rechargeable Mg2+/Li+ hybrid battery (MLHB) [14–16], enables batteries with both the advantages of LIBs and MBs. These mainly include excellent reaction kinetics and battery safety. Specifically, the reaction kinetics at the cathode is highly improved as the reactions are dominated by Li+. In addition, the advantages of using metal Mg as the anode are maintained because Mg2+ are deposited at the anode side with reversibility [14, 17–19]. The dual-salt electrolyte acts as the Li-ion reservoir to ensure the reaction on the cathode side during cycling. The addition of Li salts not only commands good reaction kinetics at the cathodes but also reduces the migration barrier and activation energy of the subsequent Mg2+ insertion [18, 20–23], thus resulting in remarkable electrochemical performances. Despite that, an anodic stability limit of 3.0 V (vs. Mg/Mg2+) has been reported for the Mg electrolyte all-phenyl-complex (APC) . This limit decreases to 2.5 V (vs. Mg/Mg2+) when LiCl or LiBF4 is added [14, 15, 17–19]. Therefore, the redox potential of the chosen cathode must be within the electrochemical stability window of the complex electrolyte and remain chemically inert to electrolyte components to avoid any parasitic reactions.
In light of the success in the commercialisation of LIBs, many efforts have been devoted to the intercalation-type cathode materials, such as Mo6S8, Li4Ti5O12, VS2, and TiO2, shown in Table S1 [14, 17–19, 24–30]. However, the development of these cathode materials is hindered by their low specific and energy density. Conversion-type cathodes can provide another route to realise the practical MLHBs with high energy density. Compared to the intercalation-type electrode materials, they can transfer multiple electrons and yield remarkably high specific capacity [31–35]. As one type of conversion cathode candidate, transition metal sulphides have received significant attention due to their high theoretical capacity, good electrical conductivity, and relatively low cost, although the kinetic reaction is slow [31–35]. One example includes tin sulphides. Tin sulphides are usually selected as cathode materials for MLHBs due to their high theoretical specific capacity and moderate voltage, matching the current dual-salt electrolytes . Tin sulphides can store Li+ through electrochemical conversion and alloying mechanisms, giving rise to a high charge storage capacity. Citing SnS2 as the typical example among tin sulphides, it offers a two-dimensional layered structure where its layers are loosely bound by weak van der Waals forces and hence easily susceptible to the intercalation of Mg2+ and Li+ ions. Further intercalation with Li+ and Mg2+ in a discharge process results in the conversion process of the lithiated SnS2, which forms metallic Sn, Li2S, and MgS [36, 37]. In the subsequent discharge process, the liberated tin alloys are reversible, with Mg2+ forming Mg2Sn . However, the conversion reaction between SnS2 and Li is not fully reversible [37, 39]. If this conversion reaction is made reversible, it can possibly provide significant room for capacity improvement.
It is important to note that SnS2 has a low activation energy barrier for the phase conversion reaction related to Li2S deposition among various metal sulphides. And it has a weak interaction with the polysulfides formed during the conversion reactions . MoS2 provides more robust anchoring sites toward the intermediate polysulfides in Li-S batteries [41, 42]. The interaction energy between the soluble polysulfide (i.e., Li2S6) and MoS2 is calculated to be larger than that of both SnS2 and graphene [40–42]. Moreover, MoS2 is highly desirable in MBs because of its layered structure, which offers good ionic mobility for reversible Mg2+ intercalation/deintercalation [18, 43]. In particular, the metallic 1T phase of MoS2 is more favourable for the electrochemical storage performance than its 2H phase due to its high intrinsic conductivity . The high electrical conductivity (10-100 S cm-1) of 1T phase MoS2 allows the formation of reasonably thick films without any binding agent. The electrical conductivity of the 1T phase electrodes compares favourably with some of the best-performing reduced graphene oxide electrodes .
Based on the above discussions, therefore, incorporating MoS2 and SnS2 in the form of vertically layered heterostructures can be a promising strategy for realising high-performance conversion reaction electrode material for MLHBs. In this work, the hierarchical configuration composed of S-deficient MoS2 and SnS2 nanoflakes was obtained via a one-step hydrothermal reaction. The hierarchical arrangement of the as-prepared SnS2-MoS2 increases the contact area between the active material and the electrolyte and shortens the diffusion path of the guest ions. Note that MoS2 cathodes in MLHBs  undergo the Mg2+/Li+ intercalation/deintercalation process, with no obvious conversion reaction due to the confined voltage window (0.01-2.0 V vs. Mg/Mg2+) as evidenced in our previous work . Therefore, the MoS2 nanosheets are stable and capable of providing an aggregation-free state to SnS2 nanocrystals, thus improving the structural stability of the heterostructures during the repeated cycling process. Unlike in previous MLHB systems [14, 17–19, 24–35] characterised by a single electron transfer at the cathode, in this current work, the multielectron transfer is simultaneously satisfied at both the electrodes. The incorporation of S-deficient MoS2 significantly promotes the active reversible conversion process of SnS2. Moreover, the direct evidence of the catalytic effect of MoS2 is also observed via ex situ transmission electron microscopy (TEM) and X-ray photon spectroscopy (XPS) analysis. The SnS2-MoS2 cathode achieves a high charge capacity of ~600 mAh g-1 at 50 mA g-1 and an excellent rate capability of 240 mAh g-1 at 1000 mAh g-1 with a negligible capacity fading rate of 0.063% per cycle across 1000 cycles. This work leverages on the new direction of using “beyond intercalation-type” electrode materials for MBs towards designing high-capacity electrode materials for dual-ion battery systems. Moreover, it also provides a deeper understanding of the mechanism of their charge storage and its associated advantages over single-ion batteries.
2. Results and Discussion
To obtain SnS2-MoS2 heterostructures, the overall process involved two major steps: (1) the slight exfoliation of MoS2 via Li-intercalation liquid exfoliation  and (2) the in situ growth of SnS2 nanocrystals on the surface of MoS2 in the absence of surfactants using a hydrothermal process. As shown in Scheme 1, MoS2 was first intercalated with Li . Upon intercalation, the Li atoms donate electrons to the d-band of the MoS2 host, and Li+ migrate into the interlamellar spacing. The reduced MoS2 layers undergo a structural distortion to 1T-type MoS2 during this step, which is partially maintained upon exfoliation and flocculation . In the exfoliation step, MoS2 intercalated with lithium was then exfoliated in water under ultrasonication. Previous work suggests that the violent nature of this reaction is able to cause the deformation of the MoS2 crystal structure , with the internal edges gaining tears, pinholes, and defects [47, 48]. In the last step, Sn4+ precursor complexes (S atom in thiourea coordinates with the Sn4+ to yield an Sn-thiourea complex)  were added. The nucleation and particle growth of SnS2 on MoS2 happens at 160°C during the hydrothermal process. The exfoliated MoS2 sheets, which possess defects in both internal edges and perimeter edges, can offer rich sites for SnS2 nucleation, creating densely loaded SnS2 nanocrystals.
To characterise the obtained materials, first, the physical structure and morphology of the obtained materials were studied with the field-emission scanning electron microscope (FESEM) (Figures 1(a) and 1(b)). The results revealed an ultrathin nanosheet morphology of the SnS2-MoS2 composite. The lateral size of the nanosheets was in the range of 100-300 nm and assembled into a flower-like structure. The scanning transmission electron microscope (STEM) image and corresponding elemental mapping images (Figure 1(c)) confirmed the homogeneous distribution of Mo, Sn, and S elements over the detected scope of the constructed hybrid material. A high-resolution transmission electron microscope (HRTEM) also was utilised to further study SnS2-MoS2 (Figures 1(f) and 1(g)) in comparison to MoS2 (Figures 1(d) and 1(e)). Fine SnS2 nanocrystals were observed on the reassembled MoS2 nanosheets. Careful investigation of the HRTEM images revealed short-range ordering (nanodomains) on the MoS2 surface. The as-obtained exfoliated MoS2 nanosheets showed a lamellar structure consisting of 6-8 layers with discontinuous curled edges (Figure 1(e)), and this interlayer spacing (6.5-6.9 Å) is slightly larger than that of the bulk MoS2 (6.1 Å) .
Moreover, as shown in Figure 1(g), the directions of individual planes on the basal surface of SnS2-MoS2 were different, with slight rotation from each other, suggesting a relatively disordered atomic arrangement. This disordered atomic arrangement resulted in additional edge sites (S vacancies on basal plane and additional edge sites of nanodomains). This indicates that many defects are formed during the SnS2 growth on MoS2 nanosheets in hydrothermal reaction than during liquid exfoliation of MoS2 to obtain few-layered MoS2 nanosheets (Figures 1(e) and 1(g)). It is worth noting that the chemical reactivity of various atomic sites on crystal surfaces toward polysulfides is different . Taking 2D layered MoS2 as a model, the binding energies of edge sites, including the Mo-edge and S-edge, are higher than those of the terrace counterpart. Thus, the existence of defect sites can maximise the anchoring effect of MoS2 [50, 51].
Additionally, nine diffraction rings corresponding to 1T phase MoS2 and SnS2 lattice constants were observed in the selected area electron diffraction (SAED) pattern of SnS2-MoS2 (see index in Figure 1(h)). From the same analysis, the interplanar spacing of 3.16 Å, 2.78, and 2.15 Å was ascribed to the (100), (101), and (102) lattices of SnS2 (according to JCPDS card no. 00-023-0677), and the lattice fringes of 3.25, 2.69, and 2.56 Å were ascribed to the (110), (012), and (021) lattices of metallic 1T phase MoS2 (according to JCPDS card no. 04-017-0898). As shown in Figure 1(e), the lattice fringe of 2.67 Å is assigned to the (101) crystalline plane of 2H phase MoS2 (according to JCPDS card no. 04-033-3080), which is a stable phase compared to the 1T phase . Compared to the exfoliated MoS2, where only a semiconducting 2H phase MoS2 lattice is found (Figure 1(e)), the highly increased metallic 1T phase MoS2 in the SnS2-MoS2 sample indicates a more distorted and defective structure, which favours the battery dynamics . This also shows that the main phase of MoS2 changed from the 2H phase to the 1T phase during the formation of SnS2 nanocrystals in the hydrothermal process.
The crystalline phases and structures of the as-obtained SnS2-MoS2 were further studied with X-ray diffraction (XRD) (Figure 2(a)). In the diffraction pattern of SnS2-MoS2 composites, the characteristic diffraction peaks of hexagonal SnS2 (JCPDS card no. 00-023-0677) can be observed. The broadened diffraction peaks of the SnS2-MoS2 composite compared to those of pure SnS2 can be attributed to the increased stack disorder and decreased particle size of SnS2 nanocrystals on the MoS2 nanosheets. Please note that some main diffraction peaks of SnS2 and MoS2 were in close angle positions with some overlaps. Besides, the unapparent diffraction peaks of MoS2 in the SnS2-MoS2 composite XRD pattern are indicative of the defect-rich structure of MoS2 and the presence of nanosized domains of mixed crystalline phases within the basal plane. Furthermore, it is noted that a new peak appeared at 2θ at around 7.5°, which corresponds to a spacing of 11.8 Å, assigned to the MoS2 with an enlarged interlayer distance. This new peak is absent in the XRD pattern of pure exfoliated MoS2 (JCPDS card no. 04-003-3080). For the SnS2-MoS2 composite, this increase in the lattice parameter is caused by the in situ growth of a bilayer of SnS2 nanocrystals with small domains in the interlayer spaces of MoS2 sheets, which led to a further expansion of the MoS2 layered structure. This can be further confirmed by the HRTEM images shown in Figure 1.
Raman spectroscopy was also carried out to collaborate TEM and XRD results of the as-synthesised materials (Figure 2(b)). The peak at 314 cm-1 is attributed to the characteristic out-of-plane vibration A1g mode for SnS2 , while the peaks at 378 and 405 cm-1 arise from the in-plane E2g and out-of-plane A1g modes of MoS2 [54, 55]. The A1g represents out-of-layer symmetric displacements of sulphur atoms along the -axis, and the E2g involves in-layer displacements of Mo and S atoms [55, 56]. The relatively lower intensity and broader E2g peak of MoS2 suggested the presence of substantial defect sites in the MoS2 structure, which could favour Mg2+ intercalation [53, 57].
X-ray photoelectron spectroscopy (XPS) was carried out on SnS2-MoS2 composite materials to determine the chemical valence of Sn and Mo in the SnS2-MoS2 composite. The full-scan spectrum in Figure 2(c) shows that the composite was composed of Sn, Mo, and S elements. The Mo, Sn, and S atomic percentages were 8.82%, 62.79%, and 28.39% (Table S2), respectively. In the high-resolution Sn 3d spectra (Figure 2(e)), the two strong peaks at around 486.2 and 494.6 eV were attributed to Sn 3d3/2 and 3d5/2, respectively, consistent with the reference data of Sn4+ in SnS2. In the high-resolution Mo 3d spectrum (Figure 2(f)), the Mo 3d spectrum had two doublet bands at around 229 and 232 eV, corresponding to Mo4+3d5/2 and Mo4+3d3/2 electrons in MoS2, respectively. The peak 226.7 eV can be ascribed to the 2s binding energy of S in MoS2. Each Mo 3d peak can be deconvoluted into a peak with lower binding energy (228.6 eV for Mo 3d5/2 and 231.8 eV for Mo 3d3/2) due to the 1T phase and another peak with higher binding energy (229.4 eV for Mo 3d5/2 and 232.6 eV for Mo 3d3/2) due to the 2H phase . Deconvolution of the Mo 3d regions of MoS2-SnS2 indicates that the 1T phase concentration of the nanosheets reaches ~71% (Table S3), suggesting that the main phase of MoS2 in the SnS2-MoS2 sample was a metallic 1T phase, reconfirming earlier results.
On the other hand, in the high-resolution Mo 3d spectrum of the pristine exfoliated MoS2, the primary phase is the 2H phase (Figure 2(g)), consistent with the results obtained from the HRTEM and XRD measurements. The high-resolution spectrum of the S 2p showed the main doublet located at binding energies of 161.5 and 162.7 eV in Figure 2(h), which was assigned to the spin-orbit coupling S 2p3/2 and S 2p1/2, respectively. Overall, compared with the pristine MoS2, Mo and S characteristic peaks in the SnS2-MoS2 composite exhibit a negative shift in the binding energy (Figures 2(d), 2(e), 2(g), and 2(h) and Table S3). It suggests an increased density of electron clouds around the Mo sites and a decreased binding energy of the S2- ions. This is attributed to the sulphur deficiency in MoS2 caused by a possible S etching process from MoS2 during the SnS2 formation. Considering that the MoS2 is a p-type semiconductor material, the electron clouds biased to MoS2 from SnS2 can form a strong electronic coupling between them . The XPS results indicate strong electron interaction between the SnS2 nanocrystals and MoS2 nanosheets.
The performance of SnS2-MoS2 in MLHB was evaluated in coin cells using the SnS2-MoS2 composite cathode and the Mg ribbon anode in a hybrid electrolyte. The hybrid electrolyte was composed of 0.25 M all phenyl complex (APC) electrolyte and 0.25 M LiCl in anhydrous tetrahydrofuran (THF). SnS2 and exfoliated MoS2 were also tested as the experimental controls. Coin cells with a pure MoS2 and SnS2 cathode were also assembled for comparison. Extended cycling was carried out at a current density of 50 mA g-1. The corresponding cycle performance and the galvanostatic curves of these three samples are provided in Figure 3. The 1st, 2nd, 20th, 50th, and 100th discharge-charge curves of the SnS2-MoS2 and SnS2 electrodes are shown in Figures 3(a) and 3(b), respectively. It can be observed that the potential plateaus of the SnS2 electrode were reduced during cycling. In contrast, those of the SnS2-MoS2 electrode reflect relatively better stability, especially in the charge curves among the potential range of 1.0-2.0 V. This can be attributed to the more facile electrode kinetics on SnS2-MoS2. We explain the observed curves as follows; the first discharge begins with a brief slope at ~1.0 V, indicating the intercalation of Mg2+ and Li+ into layered MoS2 and SnS2 (Equations (1a) and (1b)) . Then, it is followed by a short plateau at ~0.5 V, attributed to the subsequent conversion of SnS2 (Equation (2)) . After that, the curve proceeds with a gentle slope towards 0.01 V, followed by a small plateau. The sloped region corresponds to the alloying process of Mg2+ with Sn (Equation (3)) . In the charge curves, a short plateau at ~0.25 V refers to the dealloying process of Mg2Sn. Upon charging to a higher voltage (1.0-2.0 V), besides LixMgyMoS2 deintercalation, Sn oxidises to SnS2 while Li2S decomposes to S . The above reactions can be described using the following equations:
From the charge-discharge curves of SnS2-MoS2, the capacity of SnS2 in MLHBs mainly comes from the reversible conversion reactions (in the voltage range of 1.0-2.0 V). Thus, realising the highly reversible conversion reaction (Equation (2)) is very important.
In the cyclic voltammetry (CV) profiles of SnS2-MoS2 (Figure 3(c)), the initial discharge cycle showed a small cathodic broad peak centred at around 0.9 V. This is attributed to the initial Li+ or Mg2+ insertion into layered SnS2 and MoS2, without phase transformation (Equations (1a) and (1b)) . The peak around 0.9 V becomes very weak in the substantial cycles. Towards the end of the first discharge cycle, the observed broad peak in the low-voltage area involves the conversion reaction of SnS2 and MoS2 to metallic Sn and Mo (Equation (2)) and the formation of the solid-electrolyte interface (SEI) film between the active materials and electrolyte . In the subsequent cycles, the cathodic peak ascribed to the conversion reaction became stable at around 0.5 V. The cathodic peak around 0.01 V appears due to the alloying reaction between Mg2+ and Sn (Equation (3)). The anodic peak around 0.25 V corresponds to the reversible Sn-Mg dealloying reaction. Considering that the redox potential of Mg2+/Mg is 0.67 V higher than that of Li+/Li, and the typical alloying reaction between Sn and Li+ is in the range of 0.1-0.5 V vs. Li+/Li , the Sn-Li alloying reaction is ruled out in MLHBs in the voltage window 0.01-2.0 V vs. Mg2+/Mg. Upon further charging, anodic peaks at 1.2 and 1.5 V, representing the reformation of SnS2 with the decomposition of Li-S and MgLi-S, are observed .
In the absence of MoS2, the CV curves of SnS2 are similar to those of SnS2-MoS2 but with a quickly decayed redox pair at ~1.2/1.5 V in the subsequent sweeps (Figure 3(c)). The observed weakened redox pair indicates that the conversion reaction is only partially reversible in SnS2. In contrast, SnS2-MoS2 shows strong overlapping of the conversion reaction redox pairs, which implies a highly enhanced electrochemical reversibility and conversion kinetics of SnS2. This discrepancy can be intuitively ascribed to the positive effect of MoS2 on the reversibility of SnS2 during cycling, which facilitates the rapid, reversible conversion process of Sn and decomposition of Li2S. The exfoliated MoS2 electrode showed highly reversible CV curves with stable cycle performance. Figure 3(e) provides the CV profiles of MoS2, which exhibited pronounced intercalation/deintercalation peaks. The first cathodic peak at around 0.15 V vs. Mg/Mg2+ is associated with the irreversible phase transition from phase 2H-MoS2 to phase 1T-MoS2 . In the subsequent CV scans, MoS2 showed two remarkable cathodic peaks at around 0.92 and 1.02 V and two corresponding anodic peaks at around 1.19 and 1.25 V.
The cycle stability of the electrodes was also studied (Figure 3(f)). SnS2-MoS2 gave an initial discharge capacity of 1009 mAh g-1 which rapidly fell to 600 mAh g-1 in the second cycle and remained reversible to about 450 mAh g-1 with plateaus around 1.25 and 1.5 V for a long-term cycle at a current density of 50 mA g-1. The exfoliated MoS2 electrode showed an initial charge capacity of 195 mAh g-1 with a highly stable cyclability. Similar to our previous work , the capacity of MoS2 in MLHB mainly comes from ion intercalation in the layered structure. The exfoliated MoS2 with enlarged interlayer spacing and a few-layered structure can provide efficient diffusion pathways for reversible Li+ and Mg2+ intercalation. In the case of the pure SnS2 electrode, it delivered an initial charge capacity of 460 mAh g-1, which kept decreasing with the increase in cycle numbers. This capacity loss is caused by the irreversible conversion reaction involving the dissolution of polysulfide intermediate products [25, 40, 60]. Figure 4(a) provides the rate performance results of the electrodes in MLHB systems at various current densities. The reversible capacities of the SnS2-MoS2 electrode in MLHBs are measured to be 600, 482, 405, 350, 290, and 220 mAh g-1 at current densities of 50, 100, 200, 500, 800, and 1000 mA g-1, respectively. With reducing current densities, the capacities recovered gradually. When the current density was decreased back to 50 mA g-1, the reversible capacity recovered to 565 mAh g-1. This was distinctly higher than SnS2, with the highest reversible capacity of ~460 mAh g-1 at low current densities of 50 mA g-1 and a high capacity of 110 mAh g-1 at 1000 mA g-1. The capacity difference deviated more at higher rates. The higher affinity of SnS2-MoS2 for polysulfide adsorption and the catalytic effect of sulphur deficiencies in MoS2 for polysulfide conversion are expected to be the contributive factors [62, 63].
Figure 4(b) compares the long-term cycling stability of these samples at 1000 mAh g-1 for 1000 cycles. The SnS2-MoS2 electrode delivers a reversible capacity of 240 mAh g-1. The capacity remains at 150 mAh g-1 after 1000 cycles, corresponding to a capacity retention of 63% and a slow capacity decay rate of 0.063% per cycle. Compared to the poor cyclability of a pure SnS2 electrode (220 mAh g-1 at the first cycle and 50 mAh g-1 after 1000 cycles), SnS2-MoS2 exhibited much higher capacity and better reversibility. Note that though MoS2 contributed a relatively limited capacity of overall capacity in magnesium-based batteries, the existence of MoS2 virtually benefits the overall electrochemical performance. Indeed, earlier reports using MoS2 as active cathode materials in various Li batteries typically revealed some behaviour of the lithium-sulphur batteries, but without the issues of low sulphur conductivity and polysulfide shuttling in discharge and charge . In particular, MoS2 with sulphur deficiencies participates in the polysulfide reactions and significantly enhances the conversion kinetics [51, 65, 66]. For SnS2, it has a low activation energy barrier for the phase conversion reaction related to Li2S deposition, according to the recent mechanism studies , and it has a weak interaction with the polysulfides formed during the conversion reactions, thus leading to a gradual loss of active materials. It is noted that the specific capacity of MoS2 first gradually increased, then decreased when cycled at high current density. We believe this is related to the restack of the pristine exfoliated MoS2. At high cycling current density, the insertion of Li+ and Mg2+ (particularly the latter) needs time to activate the layered structure. As a result, the capacity change of the MoS2 electrode gradually increased in the first few cycles. The insertion also gradually caused structure damage and/or collapse; thus, the capacity gradually decreased after the material was fully activated.
Here, to confirm the improved polysulfide adsorptivity of the SnS2-MoS2 sample, Li2S2 was selected to experimentally model the high-polar lithium polysulfides as the adsorbate for static adsorption. 0.005 M Li2S2 was prepared by chemically reacting sulphur with Li2S in tetrahydrofuran (THF). SnS2 and SnS2-MoS2 with the equivalent total surface area were added to the above solution for comparison. Brunauer-Emmett-Teller (BET) analysis of MoS2, SnS2, and SnS2-MoS2 samples are shown in Figure 4(c) with their N2 adsorption/desorption isotherms and surface area values. According to Figure 4(d), after prolonged contact with Li2S2, the colour of the dissolved polysulphides with SnS2-MoS2 turned to light. However, there was a less observable effect on the solution with SnS2 since the tawny colour remained darker. This difference suggests a significantly promoted affinity of Li2S2 molecules of SnS2-MoS2 compared with SnS2. Therefore, the remarkable improvement in cycling stability and capacity of SnS2-MoS2 can be ascribed to the realisation of a highly reversible conversion reaction in SnS2. This is through immobilisation of soluble polysulfide species via stronger chemical binding with MoS2 and facile redox reaction propelled by SnS2 and S-deficient MoS2.
The full utilisation of both reversible conversion and alloying reactions in SnS2 via the combination with MoS2 can boost its ion storage capability efficiently and result in a significantly prolonged cycle life. Furthermore, the good high-rate cycle performance can be attributed to the favourable ion diffusion and charge-transport kinetics in the electrochemical reactions of the SnS2 towards Mg2+ offered by the 1T phased MoS2. As the main reaction mechanism in MoS2 in Mg-based batteries is intercalation with limited conversion reaction, MoS2 nanosheets with large interlayer spacing can offer stable and high-speed ion transport across the interfaces for the loaded SnS2 nanocrystals [43, 44, 61]. In this work, the MoS2 in the composite contains a high ratio of metallic 1T phase, which has a higher conductivity than the 2H phase, which thus favoured charge-transport kinetics in the electrochemical reactions of the SnS2 towards Li+/Mg2+.
Further evidence of the reversibility of the conversion reaction was investigated with postcycling ex situ TEM and XPS analyses on SnS2-MoS2 electrode materials. First, some coin cells were discharged to 0.01 V, and the others were fully charged to 2.0 V. The electrodes were extracted from the coin cell in the argon-filled glove box, soaked overnight with dimethylformamide to wash off the residual electrolyte, and dried in the vacuum oven. The ex situ HRTEM and the SAED pattern of the fully discharged and charged electrodes are shown in Figures 5(a) and 5(b), respectively.
Figure 5(a) shows the crystalline lattices corresponding to (401) Mg2Sn and (200) Li2S at a full discharge state. Corresponding SAED diffraction rings corresponding to 1T phase MoS2 (according to JCPDS card no. 04-017-0898), Li2S (according to JCPDS card no. 00-023-0369), MgS (according to JCPDS card no. 00-035-0730), and Mg2Sn (according to JCPDS card no. 00-033-0866) lattice constants are shown in Figure 5(b). After a full charge, crystalline lattices are observed upon full charge and assigned to SnS2 and 1T-MoS2 (Figure 5(c)). The diffraction rings corresponding to SnS2 (according to JCPDS card no. 00-023-0677), 1T phase MoS2 (according to JCPDS card no. 04-017-0898), LiSnS2 (according to JCPDS card no. 00-022-0692), and MgS (according to JCPDS card no. 00-035-0730) were observed in Figure 5(d). The recovery of the initial state of SnS2 implies the reversibility of the conversion reaction. But the residual MgS suggested the conversion reaction between SnS2 and Mg was not fully reversible, which is primarily due to the electrochemical inactivity of MgS . The presence of Li+ enables reactivation of MgS and MgS2 either through an ion exchange reaction and transforms them into rechargeable Li2S or through strong coordination of Li+ with the surface S2- of MgS that increases its solubility and reduces the reoxidation kinetic barrier by forming MgLi-S . Therefore, the reversibility of MgS can be further improved by increasing the amount of Li+ in the electrolyte, necessitating an excess electrolyte volume.
We also observed in the Mo 3d XPS spectra the interaction between S-deficient MoS2 and polysulfides during the conversion reactions (Figure 5(e)). The two doublet bands of Mo 3d peaks were deconvoluted into black and blue peaks. The deficiencies in MoS2 rendered the surface of MoS2 electron rich. When discharged to 0.01 V, the sulphur-deficient MoS2 component (blue curve) in SnS2-MoS2 gave a significantly diminished intensity. After sulphur deposition, the weaker XPS signal from the sulphur deficiencies suggests the electron transfer from the former to the latter . When charged back to 2.0 V, the sulphur deficiencies were restored, as suggested by the increased integrated intensity of blue peaks. This indicates S release and provides indirect proof for sulphur deficiencies as the origin of enhanced reversibility in polysulfide electrochemical reactions. The high-resolution Sn 3d spectra of SnS2-MoS2 electrodes at different states, pristine, discharged to 0.01 V, and charged to 2.0 V, are compared in Figure 5(f). In the discharge state, Sn 3d peaks shift downwards to 493.5 eV and 485.1 eV, which indicates the reduction of Sn. Upon charge, the Sn 3d shifted back to the 494.5 eV and 486.0 eV positions, indicating oxidation of Sn (i.e., the reformation of SnS2). In a nutshell, the results obtained from the ex situ HRTEM and XPS of the cycled electrode materials provided further evidence of the reversibility of the conversion reactions of MoS2-SnS2 in MLHBs.
Considering the possible influence of electrolytes on the electrochemical performance, the performance of the cells with APC-based mixed electrolyte was further tested. Figure S1 shows the electrochemical performance of SnS2-MoS2 cathode in three cells: SnS2-MoS2|(PhMgCl)2-AlCl3/THF|Mg cell (MIB), SnS2-MoS2|LiCl+(PhMgCl)2-AlCl3/THF|Mg cell (MLHB), and SnS2-MoS2|LiCl/THF|Li cell (LIB). From Figures S1a, d, and g, we observe that SnS2-MoS2 delivers the lowest capacity in the MIB cell than in the other two cells of LIB and MHLB. Past reports show that magnesium storage performance in both SnS2 and MoS2 has a strong dependence on size and morphology of electrodes . Besides the Mg2+ intercalation/deintercation reaction in layered SnS2-MoS2, Figure S1a shows broad cathodic and anodic peaks referring to the conversion reaction between SnS2 and Mg2+. The charge-discharge curves of MIBs show sloping curves with a capacity around 200 mAh g-1 in the first few cycles (Figures S1d and g) and which stabilized at around 80 mAh g-1 after 300 cycles. As compared in Figure S1, the CV curves and discharge/charge curves in MLHBs show more obvious cathodic/anodic peaks and potential plateaus with lower polarization and larger enclosed areas than those in MIBs, suggesting enhanced charge storage performance than the former. This can be attributed to the fast Li+ diffusion kinetics, which leads to more efficient reactions with SnS2-MoS2. The lower polarization is likely resulted from the lower electronic conductivity of the dual salt electrolyte than APC. The additive LiCl enhances the conductivity of the hybrid electrolyte, and the Cl- anion diminishes the interface resistance of Mg deposition/dissolution by destroying the blocked solid electrolyte interphase (SEI) on the Mg anode [15, 28] As a comparison, the SnS2-MoS2|Li+|Li cells with lithium electrolytes are also represented. Considering that the redox potential of Mg2+/Mg is 0.67 V higher than that of Li+/Li, the voltage window of LIB was limited between 0.8 and 2.8 V . Notably, the hybrid cells show lower initial capacities but better long cycle stability than lithium cells (41% and 67% capacity retention for LiCl+(PhMgCl)2-AlCl3/THF and LiCl/THF, respectively). This may partly be due to more stable deposition-dissolution efficiencies of Mg than Li , and the cathode structure distortion caused by a more rapid Li+ insertion/deinsertion kinetics is more severe than that of the Mg2+ case.
The initial coulombic efficiency of the electrode was also evaluated. As shown in Figure 3(f), the initial coulombic efficiency of SnS2, MoS2, and SnS2-MoS2 was 49%, 75%, and 60%, respectively. The increased coulombic efficiency of SnS2-MoS2 than that of pure SnS2 can be attributed to the following reasons. As shown in Figure S2, the pure SnS2 shows a stacking structure with large lateral size (1~2 μm) and thickness (~50 nm). On the contrary, the SnS2-MoS2 shows an assembled flower-like structure with an ultrathin nanosheet morphology of SnS2. Due to the strong electrostatic interactions of Mg2+ in cathode materials, SnS2-MoS2 with shorter diffusion length will significantly improve the diffusion kinetics of electrolyte ions, thus decreasing the Mg2+ trapping. Besides, as shown in Figure 1(e), the exfoliated MoS2 nanosheets showed a lamellar structure with large interlayer spacing (6.5-6.9 Å), which also offers the privilege of Mg2+ and Li+ intercalation. Another possible reason would be related to the different reversibility of the conversion reactions. The electrochemical inactivity of MgS suggested by the conversion reaction between SnS2 and Mg is not fully reversible. However, SnS2-MoS2 shows better reversibility than pure SnS2 because of the immobilisation of soluble polysulfide species via stronger chemical binding with MoS2.
We also compared the coulombic efficiency of SnS2-MoS2 in different electrolytes. Figure S1g shows that the coulombic efficiency of SnS2-MoS2 in MIBs is lower than in LIBs and MLHBs during the whole cycle process, and the SnS2-MoS2 shows rapid capacity drop in Mg-only electrolyte. This is due to the poor reversibility attributed to the formation of electrochemically inactive MgSx species, while Li+ presence in MLHB can improve the reversibility of the cell with a higher coulombic efficiency. This phenomenon can also be observed in the work by Zhang et al. , where both the charge/discharge polarization and coulombic efficiency become worse with less Li-salt addition. This further indicates that the reversibility arises from chemical reactivation of MgS and MgS2 by Li+. One possible solution is to increase the Li+ concentration in the electrolyte. But the solubility of Li salts limits the concentration of Li+ within ethereal solvents, which requires a large amount of electrolyte solvent with enough Li+ [8, 31]. If the amount of electrolyte can be reduced to be comparable to that of conventional LIBs, the energy density of the MLHB would be significantly enhanced and they would be more widely applicable for various fields such as electric vehicles. However, note that the initial coulombic efficiency of SnS2-MoS2 in LIBs is lower than that in MIBs. This is because of the common solid electrolyte interphase formation on the Li anode in LIBs. In MLHBs, the high reversibility and efficiency of the electrochemical deposition-dissolution of magnesium at the anode side in the APC electrolyte can contribute to a higher initial coulombic efficiency.
In summary, a rationally designed SnS2-MoS2 heterostructure composite was applied as a high-capacity cathode material in MLHBs. Such rational design of SnS2-MoS2 features several advantages which can boost their electrochemical performance. First, the incorporation of S-deficient MoS2 and SnS2 offers both strong adsorption for polar polysulfide intermediates and effective transformation between sulphur to Li2S during cycling, thus significantly promoting the active reversible conversion process of SnS2. Second, the MoS2 with a distinctive layered structure ensures a fast ion transport across the interfaces, and the high electrical conductivity of 1T-MoS2 is favourable for the charge-transport kinetics in the electrochemical reactions of the SnS2 towards Mg2+. Third, the stable MoS2 nanosheets enable the aggregation-free state of embedded SnS2 nanocrystals, which improves the structural stability of the hybrid materials during the repeated cycling process. This work unlocks the potential of SnS2 that exploits both conversion and alloying reactions and paves the way to explore advanced multielectron systems coupled with safe Mg anodes.
The authors declare that the main data supporting the findings in this study are available within the article and its supplementary information. Additional data are available from the corresponding author upon reasonable request.
Conflicts of Interest
All authors declare that there is no conflict of interest regarding the publication of this article. Xiu Song Zhao is an editorial board member of Energy Materials Advances.
X.F. and X.S.Z. conceived the idea. X.F. designed the experiments. M.T., Y.W., and R.R.G were involved in the experimental setups of the project. M.T and X.F. wrote the first manuscript draft and revised all the secondary drafts. All authors were involved in the discussion for data analysis and commented on the manuscript.
XF thanks the University of Queensland (UQ) for offering IPRS and UQ Centennial scholarships and the PhD research startup foundation of North University of China. The authors gratefully acknowledge the facilities and technical assistance of the Australian Microscopy and Microanalysis Research Facility at the UQ Centre for Microscopy and Microanalysis. This work was supported by the Australian Research Council (Project ARC FL170100101) and the Research Project Supported by Shanxi Scholarship Council of China (2021-128) and supported by the Fundamental Research Program of Shanxi Province (20210302124356).
Figure S1: the electrochemical performance of SnS2-MoS2 cathode in three cells: SnS2-MoS2|(PhMgCl)2-AlCl3/THF|Mg cell (MIB), SnS2-MoS2|LiCl+(PhMgCl)2-AlCl3/THF|Mg cell (MLHB), and SnS2-MoS2|LiCl/THF|Li cell (LIB). Galvanostatic discharge and charge curves for SnS2-MoS2 at first three cycles in (a) MIBs, (b) LIBs, and (c) MLHBs at a current density of 50 mA g-1. Cyclic voltammograms of SnS2-MoS2 in (d) MIBs, (e) LIBs, and (f) MLHBs at a rate of 0.02 mV s-1. (g) Long cycle performance of SnS2-MoS2 in MIBs, LIBs, and MLHBs at 50 mA g-1. Figure S2: FESEM images revealing structure details of pure SnS2 at high (a) and low (b) magnifications. Table S1: comparison of MLHB cell performance of different electrodes. Table S2: the Mo, Sn, and S atomic percentages from XPS survey spectrum of SnS2-MoS2. Table S3: contributions from 1T and 2H phase components in the Mo 3d spectrum of SnS2-MoS2 and exfoliated MoS2, respectively. (Supplementary Materials)
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