Controlled-Size Hollow Magnesium Sulfide Nanocrystals Anchored on

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Controlled-Size Hollow Magnesium Sulfide Nanocrystals Anchored on Graphene for Advanced Lithium Storage Baoping Zhang, Guanglin Xia, Wei Chen, Qinfen Gu, Dalin Sun, and Xuebin Yu ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.8b07770 • Publication Date (Web): 28 Nov 2018 Downloaded from http://pubs.acs.org on November 29, 2018

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Controlled-Size Hollow Magnesium Sulfide Nanocrystals Anchored on Graphene for Advanced Lithium Storage Baoping Zhang, † Guanglin Xia, † Wei Chen, † Qinfen Gu,*§ Dalin Sun † and Xuebin Yu *†

†Department

§Australian

of Materials Science, Fudan University, Shanghai 200433, China

Synchrotron (ANSTO), 800 Blackburn Road, Clayton, 3168, Australia

Corresponding Author

* E-mail: [email protected]; [email protected]

ACS Paragon Plus Environment

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ABSTRACT: Magnesium sulfide (MgS), representative of alkaline-earth metal chalcogenides (AEMCs), is a potential conversion/alloy-type electrode material for lithium ion batteries (LIBs), by virtue of its low potential, high theoretical capacity and abundant magnesium resource. However, the limited capacity utilization and inferior rate performance still hinder its practical application, and the progress is rather slow due to the challenging fabrication technique for MgS. Herein, we report a series of controlledsize hollow MgS nanocrystals (NCs) homogenously distributed on graphene (MgS@G), fabricated through a metal hydride framework (MHF) strategy, and its application as advanced electrode material for LIBs. The hollow structure of MgS NCs are mainly attributed to the Kirkendall effect and the escape of H atoms from metal hydride during sulfuration. The as-synthesized MgS@G demonstrates robust nanoarchitecture and admirable interactions, which assure spatially confined lithiation/delithiation process, optimize the dynamics of two-steps conversion/alloying reactions, and induce a synergetic pseudocapacitive storage contribution. As the result, a representative MgS@G composite delivers a largely enhanced capacity of >1208 mAh g−1 at a current density of 100 mA g−1, and a long-term cycle stability at a high current density of 5 A g−1 with a


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capacity of 838 mAh g−1 over 3000 cycles, indicating well-sustained structural integrity. This work presents an effective route towards the development of high-performance magnesium-based material for energy storage.

KEYWORDS: magnesium hydride; magnesium sulfide; hollow nanocrystals; controlledsize; lithium storage


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In the past decades we have witnessed the increasing demand for clean and renewable energies, as well as the significant development of energy storage conversion/storage systems represented by LIBs.1-3 In this context, constant scientific concern focused on the rational design and innovation of electrode materials for LIBs to satisfy the everincreasing demands of high performance, low cost and environmental friendly. Among the myriad of candidates, considerable attention has been paid on metal sulfides (MSx, M=Mn, Co, Ni, etc.) with diverse structures owing to their high specific capacities exceeding that of the commercial graphite (372 mAh g−1).4-6 Beyond that, as representative of alkaline earth metal chalcogenides, magnesium sulfide (MgS) is an attractive conversion/alloy-type electrode material for LIBs by virtue of its fascinating electrochemical activity, low and safe potential (0.47 V vs. Li/Li+) and light weight (1.45 g cm−3) and high theoretical specific capacity (951 mAh g−1) solely calculated based on the conversion process (MgS + 2Li+ + 2e-↔ Mg + Li2S, Equation 1). This value is much competitive comparing with that of most MSx, such as ZnS (825 mAh g−1), MoS2 (669 mAh g−1), and SnS2 (645 mAh g−1). Moreover, the further alloying process of magnesium with Li-ions (Mg + xLi+ + xe- ↔ LixMg, Equation 2) will contribute additional lithium storage


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capacity. More importantly, the raw materials of MgS especially the magnesium resource are low cost, natural abundant and eco-friendly, which provide the potential for its largescale industrial application. In spite of these attractive and competitive features, however, the progress of MgS electrode is rather slow comparing with other MSx. Basically, the main obstacles in MgS electrode is its limited capacity utilization and inferior rate performance as in many other MSx, resulting from the poor electronic/ionic conductivity and the pulverization caused by large volume changes during cycling. Furthermore, the self-growth and aggregation of active materials upon cycling will hinder the complete lithiation/delithiation, leading to poor reversibility.

7, 8

Generally, these

common issues can be effectively solved in most of the metal sulfides by rational design of materials with favourable morphology and structure.9-11 For example, diverse nanostructured metal sulfides (hollow microspheres, nanoplates, nanowires, etc.) or composites with conductive carbon (graphene, CNTs, etc.) have been explored to improve the electrode kinetics, buffer the volumetric strains and prevent the agglomeration, leading to largely enhanced Li-storage performances.12-17 Unfortunately, owing to the high sensitivity of MgS to moisture, most of the conventional sulfuration


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methods for electrode optimizing (e.g., hydrothermal methods, solvothermal methods, electrospinning and spray) are not adoptable for MgS,18-23 leading to formidable challenges in the construction of MgS electrode in a controlled manner with high purity. Therefore, seeking feasible techniques to achieve the rational design and fabrication of MgS with favorable morphology and optimized structure is essential to unlock its potential electrochemical performances. Herein, we design a facial metal hydride framework (MHF) strategy for the fabrication of controlled-size monodisperse hollow MgS NCs uniformly anchored on flexible graphene (MgS@G) as an advanced anode material for LIBs. In the MHF strategy, graphene-supported MgH2 NCs (MgH2@G) plays a vital role as nanoreactors and structure directing during sulfuration, which gives rise to the formation of hollow structure based on the Kirkendall effect and the escape of H atoms. Moreover, this strategy is highly controllable and a series of graphene-supported hollow MgS NCs with different crystallite sizes can be realized by tailoring the sizes of starting MgH2 NCs. Taking advantage of the robust nanoarchitecture and admirable interactions derived, MgS@G is endowed with sufficient active sites, rapid transfer for electrons/Li and spatially confined


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lithiation/delithiation process when evaluated as anode material for Li-storage. More importantly, it largely optimizes the dynamics of the two-steps conversion/alloying reactions and induces a synergetic pseudocapacitive storage contribution, leading to substantially enhanced capacity utilizations, excellent rate capability and prolonged stability. As a result, the representative MgS@G delivers a high capacity of >1208 mAh g−1 at a current density of 100 mA g−1 and long-term cycle life even at a high current density of 5 A g−1 over 3000 cycles. This work presents an effective route towards the development of high-performance magnesium-based material for energy storage. RESULTS AND DISCUSSION Figure 1a illustrates the synthetic process of MgS@G schematically. In a typical synthesis, monodisperse MgH2 NCs uniformly anchored on graphene (MgH2@G) was initially fabricated by a self-assembly of di-n-butyl magnesium on graphene based on our earlier studies. 24 Subsequently, the graphene-supported MgH2 NCs were converted to MgS NCs via a solid-gas reaction with sulfur: MgH2 (s) + S (g) → MgS (s) + H2 (g). During sulfuration, MgH2@G plays a vital role as nanoreactors and structure directing, which largely facilitates the conversion process by providing sufficient contact and reduced


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mass transport distance, resulting in well inherited overall morphology. Interestingly, the Kirkendall effect caused by the disparate diffusion rates of Mg and S atoms during sulfuration, 25-30 accompanied with the escape/diffusion of H atoms from MgH2, gives rise to an outward growth of MgS shell and finally the formation of MgS NCs with hollow structure (Figure 1b). Moreover, a series of hollow MgS NCs with controlled-sizes can be achieved by tailoring the starting particle sizes of MgH2 NCs on graphene. Based on the density functional theory (DFT) calculations, MgS rock salt structure shows the most stable exposed {100} face with the lowest surface formation energy of 0.412J m-2. This is in agreement with the TEM observation of most MgS (200) planes in the following section. The calculated binding energy further indicates the formation of MgS@G with favourable binding energy of -1.143 eV. Moreover, the calculated charge density difference map (Figure 1c) displays a significant charge accumulation between bottom Mg atoms and graphene, and small negative charge regions are found between bottom S atoms and graphene. Notably, there are small positive charge isosurfaces in the second MgS layer of the structure from the interface due to charge redistribution when MgS interacts with graphene. Therefore, the electron transfer between MgS NCs and graphene derived at


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nanoscale assures a robust structure of MgS@G and makes it an ideal candidate for lithium storage. Figure 2 presents the morphological and structural characteristic of MgS@G with representative average diameters of 8.6, 14.8, 41.8 and 52.9 nm (denoted as MgS@GA, B, C, and D, respectively), resulting from the corresponding sulfuration of MgH2@G with average sizes of 8.5, 13.3, 36.7 and 45.1 nm (Figure S1). It can be seen that the crystallite sizes growth of MgS NCs on graphene are largely controlled by the tailored starting particle of MgH2 NCs, the disparate diffusion rates of Mg and S atoms at the reaction temperature, as well as the confinement of graphene during sulfuration. Taking MgS@G-C sample for example, the SEM/TEM images in Figure 2 (c1-c2) reveal that the flexible graphene sheets provide stable structural support for hollow MgS NCs. According to the high-resolution TEM (HRTEM) image of Figure 2(c3), the nanocrystals clearly display lattice fringes of 2.6 and 3.0 Å, corresponding to the (200) and (111) planes of MgS, respectively, which is also consistent with the selected-area electron diffraction (SAED) patterns (inset of Figure 2(c2)). Furthermore, the scanning transmission electron microscopy (STEM) images and the elemental mapping images of Mg, C, and S provide


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more evidences for the uniform distribution of hollow MgS NCs throughout the graphene (Figure 2(c4)). Specifically, the hollow structure of MgS NCs on graphene can be further verified by the elemental line-scan profile (Figure 3a-b), which reveals that the content of Mg and S in the fringe area are higher than that in the centre area. The typical X-ray power diffraction (XRPD) pattern of MgS@G (Figure 3c) shows that all the diffraction peaks can be matched to MgS cubic phase, indicating the formation of single MgS phase with high purity. In addition, the obtained MgS@G all exhibit low peak intensity and broad diffraction peak width, indicating the nanocrystalline size nature (Figure S2). Figure 3d shows the X-ray photoelectron spectroscopy (XPS) of MgS@G-C with the presence of C, Mg and S. The corresponding high-resolution XPS in Figure 3e displays the Mg 2p spectrum assigned to MgS (51.1 eV), and the S 2p spectrum with two characteristic peaks of S 2p3/2 (162.5 eV) and S 2p1/2 (165.1 eV).

31

The porosity structure of MgS@G resulting

from the assembly of hollow MgS NCs and graphene can be revealed by the nitrogen adsorption/desorption test. According to Figure 3f, it can be categorized as type IV with distinctive hysteresis loop between relative pressures (P/P0) of 0.5 and 1.0 that confirms


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the mesoporous characteristics of this sample. The Brunauer-Emmett-Teller (BET) surface of MgS@G-C is calculated to be 133.5 m2 g−1, with a pore size distribution from 2 to 20 nm, further indicating its mesoporous structure. Thermogravimetry analysis (TGA) reveals a high loading ratio (72.0 wt. %) of MgS in the composite (Figure S3). As comparison, pure MgS (in absence of graphene) was also prepared by the same procedure using ball milled MgH2 particles as magnesium source (Figure S4-S5). According to the cyclic voltammetry (CV) curves (Figure 4a), in the initial cathodic scan, a peak centred at about 0.4 V is close to the conversion from MgS to metallic Mg and Li2S (Equation.1).8,

32

The broad peaks below 0.2 V can be related to the multi-steps

process of Mg alloying with Li-ions (Equation.2). 33-35 In the anodic scan, Li-ions can be reversibly extracted from LixMg alloys in low voltage (0.2 V), and the peak at about 0.6 V well confirms the regeneration of MgS. The low lithiation/delithiation potentials of MgS are consistent with the value calculated based on the theoretical Gibbs free energy of the conversion reaction, which is beneficial for a low and safe potential window for practical application. For the subsequent cycles, the pairs of anodic/cathodic peaks are nearly overlapped, except the broad peak ascribed to the formation of SEI layer (~0.8 V) in the


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first scan, indicating a good reversibility. Correspondingly, the representative dischargecharge voltage profiles of MgS@G-C (Figure S6) share similar slanted charge/discharge plateaus, which are in perfect fit with the above CV results. The cycling performances of MgS@G, pure MgS and pure graphene electrodes were tested based on the whole mass of electrodes (Figure S6-S7), and the corresponding capacities contribution of MgS NCs were calculated. As shown in Figure 4b, MgS@G in different nanoparticle sizes demonstrate nearly similar cycling performances, which all outperform the pure MgS. Taking MgS@G-C for example, it delivers an initial discharge/charge capacity of 2622/1649 mAh g−1, corresponding to a Coulombic efficiency (CE) of 62.9 %. The irreversible capacity is mainly attributed to the irreversible formation of SEI layer in the first cycle. A stable capacity of 1208 mA h g−1 is maintained over 100 cycles, corresponding to a high CE of 99.5%. As a comparison, the pure MgS exhibits a first discharge capacity of 1544 mAh g−1 with a low CE of 42.6%; and it remains a low capacity of 532 mAh g−1 after 100 cycles. The similar performances of MgS@G electrodes can be mainly ascribed to the hollow nanostructure and monodisperse distribution of MgS NCs on graphene, which demonstrate fast reaction kinetics, high


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capacity utilizations and electrochemical stabilities during cycling. Furthermore, MgS@G electrodes also display comparable excellent rate capabilities particularly when cycling at high current densities (Figure 4c). Specifically, MgS@G-C delivers a decent reversible capacity of 1555, 1291, 1028, 838, and 652 mAh g−1 at current densities of 100, 200, 500, 1000 and 2000 mA g−1, respectively. When the current density reduces back to 100 mA g−1, a high capacity of 1415 mAh g−1 can be restored. When cycled at a high current density of 5 A g−1 after an active process, an admirable capacity of 838 mAh g−1 can be maintained over 3000 cycles (Figure 4d). Noting that the slightly increased capacity during cycling is mainly ascribed to the activation process of MgS nanoparticles caused by the repeated lithization/delihtization processes and the capacitive lithium-storage behavior.36 To the best of our knowledge, the achieved excellent long-time stability at such a high current rate outperforms most of the Mg-based electrodes and MSx electrodes for LIBs (Table S1). The intriguing rate capability and superior cycling stability at high current density imply the fast reaction kinetics and the well-sustained structure integrity of MgS@G.


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Figure 4e shows the electrochemical impedance spectra (EIS) of MgS electrodes after 300 cycles compared with the pure MgS. The Nyquist plots of MgS@G electrodes exhibit much smaller depressed semicircle at the middle-to-high frequency regions, indicating a lower charge transfer resistance and good electronic conductivity. The corresponding fitting values of the charge transfer resistance (RCT) for MgS@G-C and pure MgS are 22.7 Ω and 156.9 Ω, respectively (Figure S8a). Besides, MgS@G displays much inclined straight line at low frequency, representing the Warburg impedance (Wo). It predicts much easier solid-state of Li+ diffusion within electrode, which can be further verified in the relationship between Z' and ω-1/2 (Figure S8b). To further elucidate the electrochemical kinetics of MgS@G, different CV profiles at stepped scan rates from 0.3 to 1.2 mV s−1 were tested (Figure 4f). The relationship of peak current (i) and scan rate (v) can be described by Equation 3: 37-40 1

𝑖 = 𝑎 ∙ 𝑣𝑏 = 𝑘1 ∙ 𝑣 + 𝑘2 ∙ 𝑣2 where i, v, a and b represents the current at a fixed potential, the scan rate, and the adjustable parameters, respectively; k1∙v and k2∙v1/2 correspond to the current


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contribution from the pseudocapacitive behaviour and the diffusion-controlled reactions. Based on the linear relationship between log i and log v plots (Figure 4g), the fitted b value is 0.77, indicating that both diffusion-controlled reaction and capacitive behaviour contribute to the charge-discharge process of MgS@G. Specifically, the inset of Figure 4g displays around 53% capacitive contribution on capacity at the typical scan rate of 0.7 mV s−1, and the capacitive contribution gradually increase upon the increase of scan rate

v (Figure 4h). When v increases to 1.2 mV s−1, the capacitive contribution reaches 60%. The

pseudocapacitive

behaviour

can

be

attributed

to

the

pore-engineered

nanoarchitecture with a large amount of sites on/near the surface areas,41 which partially accounts for the fast charge storage and excellent rate performance of MgS@G. To reveal the lithium storage mechanism, the structure and phase evolution of MgS@GC was elucidated by in operando synchrotron XRPD during the first cycle. As shown in Figure 5a, the diffraction peaks corresponding to the (200), (220) and (222) planes of MgS NCs can be identified along with the diffraction peaks from Cu (current collector). During discharge process, the intensity of MgS peaks gradually decreases, then in the charge process the peak intensity correspondingly recovers, which indicates a good


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reversibility of the electrochemical cycling. In Figure 5b, a new Bragg peak at about 16.0 o

(d = 2.46 Å), appeared during the discharge process, can be ascribed to the (101) peak

of emerging Mg phase from Eq.1. The peak intensity gradually increases, suggestive of the conversion reaction continued. From the further alloying reaction of Mg to LixMg, the peak intensity gradually decreases with the increase of peak width due to the poor crystallinity and internal strain of the crystals, and vice versa in the charging process. The formation of Li2S phase is not observed in the XRPD patterns because of amorphous kind or very small crystal size of this intermediate phase in the cell. DFT calculations further verify the favourable two-steps of conversion/alloying reactions of MgS@G during lithium storage process. It can be speculated that Li-ions are first adsorbed onto MgS NCs and graphene surfaces, and then intercalate or substitute into MgS upon discharging process. Interestingly, low coverage percentage of Li above the void position of MgS indicates a slightly positive adsorption energy of 0.164 eV (Figure 5c), and the energies steadily decrease with the increasing Li coverage percentage, indicating a favourable Li interaction with coordinated ions adsorption effect. While the formation energies of LixMg1-xS are all positive values, which increase with more Li atoms


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replacing Mg atoms in MgS structure (Figure S9a), indicating the formation of LixMg1-xS is energy unfavourable. In addition, Li-ions inserted into MgS will induce structure to expand, decompose and convert into Mg and Li2S with a calculated reaction energy of 1.196 eV, indicating a self-efficacy process. Moreover, based on the Li-Mg binary phase diagram (Figure S9b), LixMg (with Li composition < 30 at. %) alloy is isostructural with pure Mg, and it can form in a wide range of Li composition, which is in agreement with the XRPD results. In the alloying process, Li atoms substitute for Mg atoms or form a mixed occupancy atomic position sharing with Mg atoms in the structure, which shows favourable formation energy (Figure 5d). As the result, during the conversion reaction with Li, part of MgS NCs produce Mg and Li2S, and then, the in-situ formation of ultrasmall metallic Mg particles enables fast alloying reaction with nearby Li-ions, which greatly promote the dynamics of the two-steps conversion/alloying reactions. Particularly, the monodisperse hollow NCs of MgS on graphene provide sufficient active sites, large volume-containable voids, and multi-channels for the fast transfer of Li-ions, which largely enables the reaction that cannot occur in the bulk analogue MgS (Figure S10), thus increasing the reversibility of overall electrode.


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As multiple nanocrystalline phases formed during cycling, the Li-driven structural and morphological evolution of MgS@G-C during cycling were characterized by TEM. Specifically, the TEM/HRTEM images (Figure 5e-g) and the corresponding FFT patterns (Figure S11) reveal the multiple nanocrystalline phases formed during the first lithiation process. As shown in Figure 5f, the lattice fringes with d-spacing of 3.30 Å and 2.86 Å correspond to the (111) and (200) planes of generated Li2S crystals. Besides, the lattice fringes with d-spacing in the range of 2.66 - 2.72 Å can be ascribed to the intermediate phase of LixMg, resulting from the (100) plane of MgS with slightly lattice expansion after insertion of Li-ions. Meanwhile, Figure 5g shows a series of lattice fringes with d-spacing in the range of 2.43 - 2.49 Å, which can be related to the (110) planes of the wide Li composition range of LixMg crystals. In addition, electrode after delithiation (Figure 5h) manifests the reformation of MgS crystals, indicating a good reversibility. The results are in good agreement with the observation from our XRPD patterns. It is noteworthy that during the first lithiation/delithiation process, MgS@G retains its initial overall nanostructure, characterized by hollow globular particles uniformly anchored on graphene layers (Figure 5i-j). The nanostructure is well maintained even after long cycles


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at a high current density. As shown in Figure 5k-l, the STEM image and the corresponding elemental mapping verify the hollow MgS NCs with thin walls and its uniform distribution on flexible graphene sheet, indicating an excellent structural integrity. The rational and robust architecture provides sufficient space for the volumetric expansion and alleviates the strain produced during the lithiation/delithiation process, which effectively prevents the active nanograins from agglomerating and assures a long-term stability. Moreover, such nanoscale pore-engineered structure with thin shell possesses a large amount of sites on the surface areas, which is essential for the pseudocapacitive storage behaviour. The lithiation/delithiation process of MgS@G can be summarized in Figure 6a. The interactions among MgS@G and the absorbed Li-ions (Li+MgS@G system) were simulated by DFT calculations to unveil the synergistic effect on the structural and performance. Figure 6b shows the electronic structure of MgS@G analysed by examining the valence electron localization function (VELF). The highly localized electrons between Mg and S atoms with maximum closer to S atom suggest the ionic bonding interaction between Mg-S bonds. The low VELF value between MgS slab and graphene layer indicates that the Van der Waals interactions are dominated in between interface, while


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the VELF minimum between Li-ions and MgS top layers indicates physisorption of Li-ions on MgS surface. Bader analysis shows that there is clear charge transfer from Mg to S atoms, with average Mg of +1.116 e and S of -2.161 e, respectively. We can see about 0.883 e transferred from each Li atom to MgS@G and 0.044 e received from each C atom in the graphene layer. In addition, calculations show the large value of local states across the Fermi level in density of states (DOS) which reveals the metallic behaviour in the present of Li+MgS@G system (Figure 6c). The predominant feature of hybridization for Li 1s orbital, S 2p orbital, and C 2p orbital is observed in the valence band region, while from the conduction band, the domination is Mg 3s, S 2s, and C 2s states, indicating the hybridization of states among MgS, Li, and graphene. All these evidences suggest the MgS@G composite can largely facilitate both the thermodynamics and dynamics of Listorage process. Moreover, the charge density difference map between MgS@G and the adsorbed Li-ions (Figure 6d) shows that, there is a pattern of charge accumulation at the interface between MgS slab and graphene layer, and a strong charge accumulation between Li-ions and MgS layer, whereas the charge depletion regions appear near the top of Li-ions and within the MgS slab. This indicates that there is a charge transfer from


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Li-ions to MgS layers then to graphene layer, which means a built in electric field directed from Li-ions to graphene can be established. As the results, the graphene is expected to behave as an electron collector to be positively charged with electron from MgS with Liions and improve the electron transfer within electrode. These results further verified that MgS@G can not only confine the lithiation/delithiation process spatially, but also optimize the lithium storage process by admirable Li+MgS@G interactions. CONCLUSION In summary, a series of controlled-size MgS@G NCs with hollow structure and homogenous distribution on flexible graphene have been fabricated by a facile metal hydride framework strategy. The hollow structure of nanocrystals are mainly attributed to the Kirkendall effect and the escape of H atoms from metal hydride during sulfuration. When evaluated as anode material for Li storage, MgS@G shows excellent reaction kinetics, high capacity utilization and electrochemical stability. It is mainly ascribed to the rational nanoarchitecture and the admirable interactions derived, which assure spatially confined

lithiation/delithiation

process,

promote

the

dynamics

of

two-steps

conversion/alloying reactions, and simultaneously induce pseudocapacitive storage


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contribution. The representative MgS@G (41.8 nm) delivers an admirable capacity of >1208 mAh g−1 at a current density of 100 mA g−1, and an excellent long-term cycle stability at a high current density of 5 A g−1 with a capacity of 838 mAh g−1 over 3000 cycles. Overall, this work presents an effective route towards the development of highperformance magnesium-based materials for lithium storage. Also, it is believed that this metal hydride framework strategy based on Kirkendall effect provides a general way of exploring many other nanostructures for numerous applications. EXPERIMENTAL METHODS Synthesis of MgH2@G and MgS@G. MgH2@G samples were prepared via hydrogenation process of di-n-butylmagnesium (MgBu2) in cyclohexane (C6H12). Typically, graphene (21.3 mg) and MgBu2 solution (1.6 mL, 1 M in heptane) were added to cyclohexane (40 mL) in a reactor vessel. The hydrogenation of MgBu2 was performed at 200 oC for 24 h under a H2 pressure of 2.5 MPa. Then, the mixture was centrifuged using cyclohexane. After dried at 100 oC for 12h, the graphene-supported MgH2 NCs was obtained. MgH2 NCs with different diamters (8.5 nm, 13.3 nm, 37.6 nm, 45.1 nm) can be synthsized by adjusting the hydrogenation condition including H2 pressure, temperature


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as well as the usage of cyclohexane (Table S2). The hollow MgS@G were frabricated via one-step thermal MgH2-templated sulfuration based on solid-gas reaction. Typically, a self-made sealed glass boat loaded with MgH2@G and appropriate sulfur powder were put in to tube frunace under a flow of Ar. The temperature was heated to 350°C at a rate of 2 °C min−1 then maintained for 3 h. As comparison, bare MgS sample without graphene was synthesized through a similar procedure except for using the ball milled commercial MgH2 as the source of Mg. It should note that most of the procedures were carried out in an Ar-filled glove box (< 0.1 ppm, O2 and H2O) due to the high reactivity of MgS to moisture. Characterization. Part of the crystalline structure of samples were characterized using Xray powder diffraction (XRPD; D8 Advance, Bruker AXS) equipped with Cu Kα radiation (λ=1.5406). During XRPD measurement, amorphous tape was used to seal the samples to protect samples from humid air. The microstructure and morphology of samples were characterized using transmission electron microscope (TEM, JEOL 2011 F) coupled with an EDX spectrometer, field-emission scanning electron microscope (FESEM, J JEOL 7500FA). The particle size distributions were calculated based on the corrsponding TEM


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images of samples, combined with the descriptive statistical analysis and Gaussian fitting.

The element chemical states was dectected by the X-ray energy disperse

spectrometry equipped with an Mg Kα radiation source (XPS, XSAM-800 spectrometer). Thermogravimetic analysis ( TGA, Netzsch STA 449 F3) connected to a masss spectrometer (MS, Hiden HPR 20) was carried out at a ramp rate of 10 °C min−1 in air. Specific surface area and porosity were determined using an adsorptionmeter (Quadrasorb SI-MP). ASSOCIATED CONTENT

Supporting Information Available: Details about the computational methods, the additional SEM images, TEM images, the analysis of TG, XRD results, and electrochemical performances of relevant samples. This material is available free of charge via the Internet at http://pubs.acs.org.

AUTHOR INFORMATION

Corresponding Author * E-mail: [email protected]; [email protected]


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

ACKNOWLEDGMENT This work was partially supported by the National Science Fund for Distinguished Young Scholars (51625102), the National Natural Science Foundation of China (51471053, 51571063, 51831009), and the Science and Technology Commission of Shanghai Municipality (17XD1400700). Part of the experiment was performed at PD beamline, Australian Synchrotron (ANSTO). DFT calculation was conducted on ASCI.

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FIGURES

Figure 1. (a) Schematic illustration of the fabrication process of MgS@G based on the MHF strategy. (b) The formation of hollow structure of MgS NCs. (c) Top (left), side (middle), and projected (right) views of the charge density difference map of MgS@G. Yellow and skyblue regions depict the isosurfaces of electron accumulation and depletion.


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Figure 2. (a-d) The particle size distributions of MgS@G-A, B, C, and D, respectively, and the corresponding (a1-d1) SEM images, (a2-d2) TEM images, (a3-d3) HRTEM images, (a4-d4) STEM images and the elemental mapping. The insets of (a2-d2) is the SAED patterns. Scale bars in (a1-d1) equal 500 nm and in the insets equal 250 nm; Scale bars in (a2-d2, a4-d4), 100 nm; Scale bars in (a3-d3), 10 nm.


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Figure 3. (a) EDS line scan of hollow MgS NCs on a STEM image of MgS@G-C and (b) the corresponding element distributions of Mg, S, and C, respectively. (c) The typical XRPD patterns of MgS@G-C, MgH2@G-C and Sulfur, respectively. (d) XPS spectrum, (e) high-resolution XPS spectra, and (f) pore-size distribution of MgS@G-C. The inset of (g) is the corresponding nitrogen adsorption-desorption isotherm.


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Figure 4. (a) CV profiles of MgS@G-C at a scan rate of 0.1 mV s–1. (b) Cycling performances at a current density of 100 mA g–1 and (c) rate capabilities of MgS@G electrodes compared with pure MgS. (d) Long-term cycling stability at a high current density of 5 A g–1 for MgS@G-C. (e) The Nyquist plots of MgS@G and pure MgS after 300 cycles at a current density of 2 A g–1. (f) CV profiles at different scan rates, (g) the corresponding log i vs. log v plots, and (h) the normalized contribution ratio at different scan rates of MgS@G-C. The insets of (g): the typical diffusion/capacitive contribution at the scan rate of 0.7 mV s−1.


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Figure 5. (a) XRPD patterns of a cycling battery and (b) the typical Bragg peaks collected by in operando synchrotron XRPD measurement. (c) Li ions adsorption energy as the function of Li ions coverage above the void position of MgS NCs. (d) Formation energy of Li-Mg alloying. (e-g) HRTEM images of MgS@G-C collected at the first lithiation state and h) the delithiation state. (i) TEM images and (j) the EDS line scan on a STEM image of MgS@G-C at the first lithiation state. The inset of (j) is the element distributions of Mg, S, and C, respectively. (k) STEM image and (l) the elemental mapping of MgS@G-C after 300 cycles at 2 A g−1.


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Figure 6. (a) Schematic illustration of the lithiation/delithiation mechanism for MgS@G. (b) Valence Electron localization functions (VELF) of MgS@G with medium coverage ratio of adsorbed Li ions. (c) The calculated Density of States (DOS) and Projected Density of States (PDOS) of Li+MgS@G. (d) Charge density difference map of Li+MgS@G. Yellow and skyblue regions depict the isosurfaces of electron accumulation and depletion.


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Table of Contents


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Controlled-Size Hollow Magnesium Sulfide Nanocrystals Anchored on

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