MAX‐phase Derived Tin Diselenide for 2D/2D Heterostructures with Ultralow Surface/Interface Transport Barriers toward Li‐/Na‐ions Storage

2D tin diselenide and its derived 2D heterostructures have delivered promising potentials in various applications ranging from electronics to energy storage devices. The major challenges associated with large‐scale fabrication of SnSe2 crystals, however, have hindered its engineering applications. Herein, a tin‐extraction synthetic method is proposed for producing large‐size SnSe2 bulk crystals. In a typical synthesis, a Sn‐containing MAX phase (V2SnC) and a Se source are heat‐treated under a reducing atmosphere, by which Sn is extracted from the V2SnC phase as a rectified Sn source to form SnSe2 crystals in the cold zone. After the following liquid exfoliation, the obtained 2D SnSe2 nanosheets have a lateral size of a few centimeters and an atomic thickness. Furthermore, by coupling with 2D graphene to form 2D/2D SnSe2/graphene heterostructured electrodes, as validated by theoretical calculation and experimental studies, the superior Li‐/Na‐ion storage performance with ultralow surface/interface ion transport barriers are achieved for rechargeable Li‐/Na‐ion batteries. This innovative synthetic strategy opens a new avenue for the large‐scale synthesis of selenides and offers more options into the practical application of emerging 2D/2D heterostructure for electrochemical energy storage.


Introduction
2D heterostructures composed of different types of 2D units have been explored as an effective approach to offer the materials with complimentary properties and ultrafast transport

Results and Discussion
As illustrated in Figure 1, the layered SnSe 2 bulk were first synthesized via a direct tin-extraction growth process by using a Sn-containing MAX-phase, V 2 SnC, and Se powders as starting materials. The V 2 SnC MAX-phase, which was synthesized at 1000 °C by using the V, Sn, and graphite as raw materials, [21,22] exhibits a hexagonal crystal system and a space group of P6 3 /mmc with alternative stacks of V 2 C and weakly bonded Sn layers. For the growth of the layered SnSe 2 bulk crystals, a selenization procedure was performed at the temperature of 900 °C with the presence of Se vapor, during which the weakly bonded Sn layers were extracted from the MAX-phase and reacted with the Se vapor to grow into large-size SnSe 2 crystals at the cold end. [23] The atomic structure of SnSe 2 consists of covalently bonded SeSnSe units, which can also be stacked into layered bulk materials via weak vdW interactions between the layers. [24,25] After the extraction of Sn atoms, a V 2 C MXene was left at the site of MAX phase. The V 2 C MXene, however, was promptly oxidized into a V 2 O 3 phase, owning to the extremely high chemical activity of the fresh V 2 C and the possible existence of oxygen-containing species within starting materials and the environment. [26] Different to the previously reported direct growth of SnSe 2 by using SnI 2 , tin chloride-iodine (SnCl 2 -I 2 ), or diphenyl diselenide (Ph₂Se₂) as the starting materials, this proposed method uses economic and safe precursors and without high-quality substrates, which is easy for scaling up for the massive and controllable synthesis. [16,[27][28][29][30][31] After post-exfoliation of the resultant layered SnSe 2 bulk into 2D SSN unit, 2D/2D SSG heterostructures are easily constructed by the vacuum filtration, which are expected to manifest low transport barriers toward Li and Na-ions when utilized as free-standing electrodes for lithium-ion batteries (LIBs) and sodium-ion batteries (SIBs). [32,33] The surface morphological and structural characterizations of the V 2 SnC MAX-phase, the synthesized layered SnSe 2 crystals, and the resultant 2D SnSe 2 unit were examined, as presented in Figure 2. Similar to other MAX phases, V 2 SnC shows a lamellar microstructure ( Figure 2a) and a hexagonal crystal structure ( Figure S1, Supporting Information). The obtained large-size SnSe 2 bulk crystals presented a lateral size up to several centimeters, as confirmed by the optical image in Figure 2b. Low-magnification transmission electron microscopy Figure 2. Surface morphological and structural characterizations of the V 2 SnC MAX-phase and the resultant 2D SnSe 2 unit. a) SEM image of the synthesized V 2 SnC MAX powder. b) Photograph of the grown SnSe 2 crystal. c) Low-magnification and d) bright-field TEM image of the layered SnSe 2 and the corresponding element mapping of Se and Sn. e) AFM image of the 2D SSN unit. f-i) High-magnification TEM and scanning TEM (STEM) images on the surfaces or within the interlayers in the 2D SnSe 2 unit. j) STEM images and the corresponding element mapping patterns showing a Sn-rich edges of 2D SSN unit. k) XRD pattern of the SnSe 2 crystals. l,m) High-resolution XPS patterns of Sn 3d and Se 3d in the SnSe 2 crystals. n) Raman spectrum of the SnSe 2 surfaces.
(TEM) confirms the layered structure and the smooth surface of the SnSe 2 sheets (Figure 2c). Bright-field TEM image and the corresponding element mapping patterns (Figure 2d) verify that the layered SnSe 2 structure delivers a uniform distribution of both Se and Sn. Then, the SnSe 2 bulk was further exfoliated into monolayer or bilayers. Figure 2e presents the thickness of a 2D SSN unit of 1.7 nm measured by an atomic force microscopy (AFM), which corresponds to a stack of two monolayers. High-resolution TEM (HRTEM) and scanning TEM (STEM) images indicate a lattice spacing of 0.29 nm, corresponding to the (011) plane of the SnSe 2 crystals (Figure 2f,g) and an interlayer spacing of 0.67 nm for the multilayers (Figure 2h,i). [34,35] Furthermore, an obviously disordered atomic arrangement and a Sn-rich edge were identified from the STEM images (Figure 2g,i,j) and the element mapping patterns (Figure 2j), which are possibly induced by the evaporation and loss of Se at the highly reactive edge regions. Besides, the typical characteristics of a high crystallinity degree without any impurities and the SnSe 2 crystals with a preferred orientation along [001] direction were identified in the X-ray diffraction (XRD) patterns, in which the highest peak located at 14.4° (Figure 2k) corresponds to the (001) planes. [36][37][38] X-ray photoelectron (XPS) and Raman spectra were further applied to collect more details on the 2D SnSe 2 unit. As demonstrated in the XPS spectra of Sn ( Figure 2l) and Se (Figure 2m), a couple of typical peaks located at 494.8 eV and 486.4 eV are attributed to Sn 3d 3/2 and Sn 3d 5/2 , respectively, with a spin-orbit-splitting gap of 8.4 eV, verifying the presence of Sn 4+ . [39] The core level of Se 3d is split into two pairs, 53.6/54.2 and 54.6/55.5 eV, which suggest the nonstoichiometric surface states originated from the loss of Se that agrees well with the STEM results. In the Raman spectrum, the typical peak located at 185 cm −1 is attributed to the in-plane vibrational mode of the SeSnSe bonds, and the other peaks (70.4, 108.4, 130.7, and 150.4 cm −1 ) indicate the presence of SeSn bonding, also suggesting a defective surface (Figure 2n). [40][41][42] To get a clear map on this new synthetic approach on 2D SnSe 2 unit from Sn-containing MAX-phase, the residual powder at the high-temperature region was characterized. As confirmed by XRD, SEM, and TEM, the residual power is dominantly ascribed to hexagonal V 2 O 3 phase ( Figure S2, Supporting Information) with a 3D curled porous structure ( Figures S3 and S4, Supporting Information) consisting of 2D units ( Figure S5, Supporting Information). This is possibly due to the high reactivity of the surface or edge after the removal of Sn layers, resulting in disordered stacking. Quantitative analysis on the energy dispersive spectroscopy (EDS) verifies that the atomic ratio of V: O is over 1, indicating the abundant oxygen vacancies in the V 2 O 3 powder, which could be achieved under a reducing environment for some metal oxides (e.g., TiO 2 and V 2 O 5 ). [43] One of grand challenges in the synthesis of 2D materials is the negative effects originated from impurities, which can be caused by the incomplete competitive reactions, the contaminations from the surrounding environment, such as solvents and gas, and the interactive effects of raw materials. It should be noteworthy that no vanadium species were detected in the SnSe 2 product ( Figure S6, Supporting Information), suggesting a high purity of the crystals. Also, there was no Sn or Se species detected in the residual V 2 O 3 powder ( Figure S7, Supporting Information), which indicates that the selenization reaction was completed within 4 h and all Sn layers in the MAX phases were extracted. Meanwhile, no obvious deposition of Se onto the V2O3 powders. Due to the high reactivity, Se atoms at the edges are easy to be evaporated, resulting in the formation of Se vacancies at the edges, as evidenced by the STEM, Raman and XPS analyses.
Then, the resultant high purity 2D SnSe 2 unit was applied for constructing 2D/2D SSG heterostructures by pairing with 2D graphene. Similar to other 2D nanomaterials, the individual or pristine 2D unit cannot meet all the requirements for the diversity of applications. [44,45] Even though they show excellent performance in FETs and photodetectors, the poor conductivity of most metal selenides, as widely recognized, remains a major challenge for achieving attractive ion transport and storage. [46] To address this issue, some structural modifications, such as heterostructure formation and doping, have been carried out to improve the intrinsic properties and obtain tailored electronic performance. For example, pairing with conductive materials has been verified as one of the most effective approaches. [47][48][49] Herein, to appropriately adapt the ion/electron transport and thus promote the utilization potentials for electrochemical storage devices, 2D/2D SSG heterostructures were designed by coupling 2D graphene with the as-synthesized 2D SSN unit. It is expected that the electrical conductivity can be significantly improved, the active channels can be enlarged within the heterostructured interfaces, and the mechanical robustness can be reinforced, which are desired properties for free-standing membrane electrodes, thus resulting in excellent electrochemical Li-/ Na-ion storage performances for LIBs and SIBs. [50][51][52][53] For constructing 2D/2D SSG heterostructures, the liquid-exfoliated 2D SSN unit and the 2D graphene oxide (GO) nanosheets were first mixed and then vacuum filtrated to produce a piece of 2D/2D SnSe 2 /GO film. After further thermal reduction treatment at 350 °C under Ar/H 2 flow, the free-standing 2D/2D heterostructured SSG films were finally obtained. XRD confirms the coexistence of SnSe 2 and graphene after the removal of oxygencontaining groups from GO (Figure 3a). High-resolution 3D topography scanned using the PeakForce QNM verifies the uniform distribution of 2D SnSe 2 and 2D graphene (inset in Figure 3a). Mechanical performance of the as-obtained 2D/2D heterostructured SSG film was first evaluated via nanoindentation measurements by using a three-sided pyramidal Berkovich indenter (inset in Figure 3b). [54] Based on the loading-depth plots in Figure 3b, the calculated maximum reduced Young's modulus and hardness of the free-standing SSG film is 5.13 GPa and 0.50 GPa ( Figures S8 and S9, Supporting Information), respectively, which is comparable to the previously reported MXene sheets for flexible electronic devices. [55] To reveal structural characterizes of the 2D/2D heterostructured SSG film, near-edge X-ray absorption fine structure (NEXAFS) spectra of Se L-edge, C K-edge, and O K-edge were collected. As shown in Figure 3c, compared with the selenium powder (Se 0 ), which has a sharp onset and two shoulder peaks at above 1444 eV, these resonances are reduced and broadened in the SSN and SSG film, indicating the high degree of local rearrangement on the surfaces, which agrees well with the above-mentioned Raman and XPS results. [56,57] In the C K-edge spectra (Figure 3d), the peak located at 285.3 eV is attributed to the C1s-π* for a typical CC bond or graphitic sp 2 bond, while the peak located at 292.5 eV is ascribed to C1s-σ* for the typical CC bond. [58] Additional primary peaks between these two peaks are related to functional groups or surface defects. Compared to the GO, the relative higher intensity of the carbon-related peak (C1s-π*) and the lower intensity of the oxygen-related peak (CO/COOH) in the SSG film suggest better structural arrangement of the aromatic ring with fewer surface modifications after thermal reduction. [59] The relative intensity ratio of these two resonances, I (1s-π*) /I (1s-σ*) ratio, is a crucial indicator for evaluating the degree of restoration of the electronic structure after the reduction of GO. [60] By calculation, the intensity ratio of GO was 0.34, which increased to 0.88 after reduction in the 2D/2D heterostructures SSG film. As for the O K-edge spectra as displayed in Figure 3e, the peak in the range between 531.5 and 532.5 eV was assigned to the π* state of the typical CO bond in the COOH group, while the peak from 534 to 535 eV was attributed to the π* state of the typical CO bond in the epoxide group. [61] The peaks located at 539.7 eV and above are assigned to the σ* state of the CO, OH, or CO bonds. [62] By comparison with GO, the peaks in the SSG film exhibited lower negative energy shifts, suggesting the partial reduction of GO into graphene in the 2D/2D heterostructured SSG film. [63] Subsequently, the favorable Li-/Na-ion transport pathways and barriers of the 2D/2D SSG heterostructure were examined and identified by both theoretical calculation and experimental validation. As shown in Figure 4a,b, the transport pathways and barriers were analyzed along the interfaces between SnSe 2 and graphene and along the SnSe 2 surfaces. The Li and Naions follow similar transport pathway mechanisms, in which both Sn and Se interstitials are the most favorable positions. For Li-ions, compared with the slow diffusion along the SnSe 2 surfaces (≈0.19 eV), the interfacial Li-ion transport is much easier, in which the diffusion energy barrier is only 0.13 eV (Figure 4b). However, the high Na-ions transport rates are identified along both the interlayer and the surface accompanied by a similar energy barriers as low as 0.1 eV (Figure 4c), which is much smaller than many previously reported advanced heterostructures for Na-ion transport, such as the Na 2 S/Na 2 Se heterostructure (0.39 eV), [64] silicene/graphene heterostructure   [65] SnS 2 /graphene heterostructure (0.25 eV), [66] MoS 2 /MXene heterostructures (0.36-0.37 eV), [67] and graphene/ bismuthene heterostructure (0.10-0.30 eV), [68] suggesting superior Na-ion storage performance of the resultant 2D/2D SSG heterostructure.
To further confirm this, half-cell models were assembled by directly coupling the resultant free-standing 2D/2D SSG electrode with the metallic Li/Na foils as the reference and counter electrodes, and relatively small discharge/charge rates were used for ensuring well ion diffusion states within both the surfaces and the interlayers. Another advantage for these free-standing 2D/2D SSG film electrodes is avoiding the potential influences from the binders and the conductive additives commonly used in traditional powder-based electrodes on the diffusion/transport and transfer of ions and electrons. It is well known that the electrochemical mechanisms of SnSe 2 for LIBs and SIBs are similar, through an intercalation-conversionalloying reaction. Specifically, the intercalation of Li/Na-ions into the SnSe 2 is first proceeded, accompanied by the formation of M x SnSe 2 (M = Li or Na), and further reactions will produce M 2 Se and Sn, in which the resultant Sn will finally react with Li/Na-ions to produce M x Sn. [69,70] For Li-ion storage ( Figure S10, Supporting Information), the initial discharge capacity is over 1100 mAh g electrode −1 at a current density of 5 mA g electrode −1 , corresponding a volumetric capacity of ≈708 mAh cm electrode −3 when a 5 µm thick 2D/2D SSG film electrode was used. When recharging, about 56% of the capacity could be maintained. From the 2 nd discharge/charge cycle, the coulombic efficiency reached more than 91% with a reversible capacity level of ≈500 mAh g electrode −1 . When the 2D/2D SSG film was employed as the electrode for Na-ion storage, an initial discharge capacity of 640 mAh g electrode −1 was obtained, which corresponds to a . After the initial two cycles, the remaining capacity was nearly 300 mAh g electrode −1 , accompanied by a coulombic efficiency of over 87% from the 2 nd cycle. After 20 discharge/charged cycles at 5 mA g electrode −1 , much higher current densities of 10-50 mA g electrode −1 were applied for the assembled LIBs and SIBs. As compared in Figure 4e,f, with the current density increased, the capacity gap between Li-ions and Na-ions gradually decreased. For example, at a current density of 10 mA g electrode −1 , the average Li-ion capacity was ≈236 mAh g electrode −1 , while the Na-ion one was ≈198 mAh g electrode . This suggested that a rapid capacity decay was identified for Li-ions compared to that for Na-ions, and no obvious difference was found for the capacities toward Li-or Na-ions in high rates. In addition, after 150 discharge/charge cycles, the residual capacity for SIBs was 182 mAh g electrode −1 , much larger than that for LIBs (103 mAh g electrode −1 ) (Figure 4g). These results indicate that the resultant 2D/2D SSG film is a promising electrode for stable SIBs. The possible reasons on the improved stability for Na-ions transport should mainly the results of a lower transport barrier for Na-ions and relatively moderate reactivity of Na-ions with the 2D/2D SSG film. Finally, by comparing with these previously reported electrodes for LIBs or SIBs, such as graphene, silicon, oxides, and MXenes (Figure 4h), the maximum volumetric Li-/Na-ion capacities of the resultant 2D/2D SSG film electrode were comparable to these graphene-based freestanding electrodes (Table S1, Supporting Information).
Finally, electronic structure and chemical analysis of the 2D/2D heterostructured SSG film were conducted for in-depth transport mechanisms. Electronic structures were first investigated by density functional theory (DFT) calculation. 2D monolayer SnSe 2 is a semiconductor with a calculated indirect bandgap of 0.87 eV ( Figure S11, Supporting Information), and graphene manifests a metallic characteristic with a Dirac point at the Γ point near the Fermi level. [71,72] Interestingly, by coupling 2D SnSe 2 with 2D graphene to form the 2D/2D SSG heterostructure, as shown in Figure 5a, the Dirac point moves upward, which indicates that the electrons could pass through The static electro potential along z-direction in the 2D/2D SSG heterostructure. c) The charge density differences and the 2D sectional view between the SnSe 2 layer and the graphene layer in the 2D/2D SSG heterostructure. Yellow isosurface represents the electron accumulation and the cyan one stands for the electron depletion. The isosurface value was set to 0.0002 e Å −3 . Comparison on the contribution ratios based on diffusion and capacitance for d) Li-and e) Na-ions. High-resolution XPS spectra of f) C 1s, g) Se 3d, and h) Sn 3d in 2D/2D SSG film electrode for LIBs and SIBs before and after cycling.
Fermi level. Based on the density of states (DOS) analysis, the primary contribution to the conduction bands is from both Sn-d and Se-p orbitals, while the Se-p orbital makes the greatest contribution to the valence band. Due to the lower electro potential of graphene (Figure 5b), the electrons transfer from graphene to the SnSe 2 layer (0.06 e), as confirmed by the differential charge density between layers (Figure 5c). This obvious charge transfer results in the formation of an electrostatic field across the 2D/2D interface and an ohmic contact appears in the 2D/2D SSG heterostructure, which could significantly reduce the interfacial resistance and thus promote electron transfer processes upon discharge and charge. [73] The kinetic was analyzed by using CV technique to distinguish the contribution ratios from either diffusion or capacitance for LIBs and SIBs (Figures S12 and S13, Supporting Information). For Li-ions, when the scan rate increased into 8.0 mV s −1 or above, the initial diffusion-controlled behavior at low scan rates was gradually changed into a capacitancedominated one (Figure 5d). This phenomenon was more obvious for larger Na-ions. With the scan rate increased from 0.2 to 1.0 mV s −1 , the capacitive contribution ratios significantly increased, achieving 85.5% at a rate of 1.0 mV s −1 (Figure 5e), indicating a dominate surface/interface-controlled capacitive storage pathways. As a result, the fast Na + ion transport was primarily achieved based on capacitive-controlled pathways, which could significantly enhance the cycling stability of the 2D/2D SSG film electrode for SIBs.
Besides, the electrolyte-electrode interfaces have crucial roles for the favorable transport of Li/Na-ions and ensure the stable discharge/charge states over cycling. To validate this, the surfaces of 2D/2D SSG film electrodes were examined by SEM and XPS after cycling. The fresh 2D/2D SSG film electrode showed a smooth surface with the presence of some wrinkles caused by the thermal reduction of graphene and a uniform distribution of Sn and Se ( Figure S14, Supporting Information). After cycling, compared with that for LIBs ( Figure S15, Supporting Information), the surface of the 2D/2D SSG film electrode was relatively rough for SIBs ( Figure S16, Supporting Information), implying the possible side reaction with the electrolyte occurred in SIBs. Another characteristic is the low Sn percentage on the electrode surface electrodes, which might be due to the reaction with active Se. Furthermore, XPS was employed to reveal the interfacial composition and the possible origins of the as-formed coating on the 2D/2D SSG film electrode surfaces. The roles of graphene in the 2D/2D SSG film for LIBs and SIBs were first confirmed. As shown in Figure 5f, an obvious Li-C peak was found in the C 1s spectra of the SSG electrode for LIBs, which was induced by the electron transfer process from lithium to carbon. [74,75] This result confirms the reactivity between the graphene and the Li-ions, which is consistent with the previous reports. [76] When the 2D/2D SSG film electrode was used for SIBs, the reactivity became less obvious, as presented by the decreased intensity and the slight peak shift, suggesting the capacitancebased contribution of graphene for Na-ions storage. [77] Then, the chemical states of Se were compared for LIBs and SIBs (Figure 5g). For LIBs, an obvious peak located at ≈52.5 eV was clearly identified, which was attributed to a reduced state of Se, based on the electrochemical reactions between active Se with Li-ions. [78,79] However, the case was distinctly different in SIBs. A sharp peak appeared in a high binding energy over 60 eV, which should be highly associated to the formation of Se-O species aroused by the possible reaction of Se anions with the carbonyl groups of the electrolytes. [80][81][82] On the other hands, by examining the Sn species at different regions of the electrode surfaces after cycling, a trace of Sn specie was found for LIBs, but obvious Sn signals were verified for SIBs (Figure 5h). These results imply that the formation of undesired Li-Se layer deposited onto the 2D/2D SSG film electrode for LIBs hindered the effective ions/electrons transport, leading to a faster performance decay. This could be further verified by the large charge-transfer resistance (R ct ) based on the electrochemical impedance spectroscopy analysis ( Figure S17, Supporting Information). For SIBs, despite the presence of side reactions with the electrolyte, a relatively stable layer consisting of both Sn and Se active species was formed between the film electrode and the electrolyte, giving rise to a stable cycling state. Besides, the rational combination of SnSe 2 with graphene into 2D/2D SSG film can result in the formation of Sn(O)C and/or SeC bonding, which are very favorable for promoting the interfacial charge transfer, improving the structural stability for adapting the undesired volume variation, and benefiting for producing a stable solid electrolyte interface (SEI) layer. [52,79] To summarize, there are several distinct advantages for the synthesis of 2D SSN units for building a 2D/2D SSG heterostructure with ultralow Li-/Na-ion transport barriers in both LIBs and SIBs applications, including: i) the synthesis of 2D SSN unit is simple and easy for scaling up by using a solidstate prepared, low-cost MAX phases; ii) the formation of the 2D/2D SSG heterostructure is easily operated via a vacuum filtration approach; iii) the ohmic contact characteristic of the 2D/2D SSG heterostructure is favorable for electron transfer; iv) the abundant active surface and interlayer spacing in the 2D/2D SSG heterostructure can facilitate electrochemical Li-/ Na-ion transport and storage; (v) the resultant 2D/2D heterostructured films can be directly used as free-standing working electrodes without any additional additives or current collectors; vi) the precursor used for the synthesis of 2D SSN units can be extended to other Sn-containing MAX phases; vii) the proposed design principle for 2D/2D heterostructures with ohmic contact can be employed for the fabrication of other 2D nanomaterials and heterostructures for various applications across electronics and energy storage devices.

Conclusions
To conclude, an innovative synthesis approach for large-scale production of SnSe 2 crystals is proposed through a simple solid-state reaction by using the Sn-containing MAX phase as the precursor. No special substrate is required for this method, and the operation is easily achieved via a one-step thermal reduction approach. By coupling the resultant 2D SSN unit with 2D graphene, the 2D/2D SSG heterostructure manifests intriguing ohmic contact and favorable Li-/Na-ion transport for LIBs and SIBs, as verified by both the theoretical calculations and the experimental validations. This work offers us a general template for the large-scale synthesis 2D metal selenides and provides new insights for 2D/2D advanced heterostructures with superior ion/electron transfer capability for a wide range of electrochemical applications ranging from electronics to rechargeable batteries.

Experimental Section
Fabrication of SnSe 2 Bulk Crystals: Typically, V 2 SnC MAX was synthesized by using metallic vanadium, tin, and graphite as raw materials and the solid-state reaction was conducted at 1000 °C under Ar flow. [21] Then, 300 mg as-prepared V 2 SnC MAX powder and 3 g selenium powder (Se, 99.5%, 200 mesh, Acros Organics) loaded into two quartz boats were placed in the upstream and middle region of the tube furnace (50 mm in diameter for the quartz tube, NBD-O1200), respectively. Subsequently, the tube was heated into 900 °C at a rate of 10 °C min −1 under continuous Ar/H 2 flow (v/v, 95:5) and maintained for 240 min. Finally, the SnSe 2 crystals could be obtained at the downstream part of the tube, and the residual vanadium-containing powder was collected.
Material Characterizations: The structures of the as-synthesized V 2 SnC MAX, SnSe 2 crystal, V 2 O 3 powder and 2D/2D SSG heterostructures were confirmed by X-ray diffraction (XRD) patterns collected on a Rigaku SmartLab diffractometer equipped with a copper radiation (λ = 0.1542 nm) at the voltage and current of 40 kV and 40 mA, respectively. Raman spectra of the as-synthesized SnSe 2 crystals were recorded on a Renishaw inVia Raman microscope using a 532 nm laser excitation. The chemical states of tin, selenium, and vanadium were examined by XPS spectra on a Kratos AXIS Supra photoelectron with a monochromated Al Kα X-ray source, an emission current of 15 mA, and a pass energy of 20 eV. NEXAFS at the C K-edge O K-edge, and Se L-edge for 2D SSN or 2D/2D heterostructured SSG film were collected at the Soft X-ray Spectroscopy (SXR) Beamline of the Australian Synchrotron, part of ANSTO. The thickness of 2D SSN was confirmed on a Bruker Icon Dimension AFM using ScanAsyst technology for automatically scanning in a peak-force tapping mode. The surface morphologies and chemical components of the V 2 SnC MAX, the SnSe 2 crystal, the V 2 O 3 powder and the 2D/2D SSG heterostructures were performed on a Zeiss Sigma VP field-emission scanning electron microscope (FE-SEM) and a JEOL 2100 transmission electron microscope (TEM, 200 kV) coupled with an Oxford XMax Silicon Drift (SDD) X-ray detector for compositional analysis. STEM characterization were performed using an aberration-corrected JEM-ARM200F NeoARM TEM (200 kV) under vacuum of 1.5 × 10 −5 Pa at room temperature, and the elemental maps were acquired by using a JEOL Dual SDD.
Mechanical Characterizations: Nanoindentation measurements on the 2D/2D heterostructured SSG film were conducted on the Hysitron TI950 Triboindentor by using a three-sided pyramidal Berkovich indenter with the tip radius of ≈150 nm, in which the area function was calibrated on fused quartz before tests. The indentation tests were performed at room temperature controlled by a function setting of loading for 15 s, holding for 30 s, and unloading for 15 s. The maximum force was set to 3000 uN, and a 4 × 4 position array was selected for tests. The reduced modulus and hardness were calculated based on the Oliver-Pharr method.
Electrochemical Characterizations: CR 2032-type coin cells were assembled in Ar-filled glove box (Mbraun) by using metallic lithium or sodium discs as the counter and reference electrodes, and freestanding 2D/2D heterostructured SSG films directly as the working electrodes (the diameter is 12 mm, and the massive weight is ≈0.6 mg for each disc film). For the electrolyte systems, 120 µL 1.0 m LiPF 6 in EC/DEC (1:1, v/v) and 1.0 M NaClO 4 in EC/DEC (1:1, v/v) with 5 wt% fluoroethylene carbonate (FEC) additives was applied for LIBs and SIBs, respectively. Electrochemical properties were evaluated by using the cyclic voltammetry (CV) technique on a CHI 760 electrochemical workstation and the galvanostatic charge/discharge technique on a Neware battery tester within the potential range of 0.0-3.0 V (vs Li + /Li) for LIBs and 0.0-3.0 V (vs Na + /Na) for SIBs. Electrochemical impedance spectra (EIS) were recorded under a frequency range between 5 mHz and 100 kHz.
Theoretical Calculation Methods: All the calculations were performed based on the DFT by using the projector augmented plane-wave (PAW) method, as implemented in the Vienna ab initio simulation package (VASP). [83] The generalized gradient (GGA) approximation proposed by Perdew, Burke, and Ernzerhof (PBE) was selected for the exchangecorrelation potential. [84] The long-range vdW interaction was described by the zero damping DFT-D3 method. [85] The cut-off energy for plane wave was set to 500 eV. All the structures were fully relaxed until the residual energy in iterative solution of the Kohn-Sham equation and force declined to less than 10 −5 eV and 0.001 eV Å −1 , respectively. The Brillouin zone integration was performed using an 11 × 11 × 1 k-mesh. The SSG heterostructure model was composed of a 2 × 2 × 1 SnSe 2 layer and a 3 × 3 × 1 graphene layer. A vacuum layer of 15 Å was added in the perpendicular direction to avoid artificial interaction between periodic images. The energies of ion diffusion (Li or Na) were calculated by the climbing-image nudged elastic band (CI-NEB) method. [86]

Supporting Information
Supporting Information is available from the Wiley Online Library or from the author.