Hydrangea-Shaped 3D Hierarchical Porous Magnesium Hydride

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Hydrangea-Shaped 3D Hierarchically Porous Magnesium HydrideCarbon Framework with High-Rate Performance for Lithium Storage Baoping Zhang, Yinsong Si, Qinfen Gu, Min Chen, and Xuebin Yu ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b10527 • Publication Date (Web): 17 Jul 2019 Downloaded from pubs.acs.org on July 17, 2019

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ACS Applied Materials & Interfaces

Hydrangea-Shaped 3D Hierarchically Porous Magnesium Hydride-Carbon Framework with HighRate Performance for Lithium Storage Baoping Zhang, † ‡ # Yinsong Si, † # Qinfen Gu, § Min Chen, †* and Xuebin Yu †* †

Department of Materials Science, Fudan University, Shanghai 200433, China

§

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

‡

School of Energy and Environment, City University of Hong Kong, Hong Kong, China

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

ACS Paragon Plus Environment

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ABSTRACT: Magnesium hydride (MgH2) is a promising anode material for lithium ion batteries (LIBs) by virtue of its high theoretical specific capacity, suitable potentials, and abundant source. However, the electrochemical performance of MgH2 electrode is still far from satisfactory due to its poor electronic conductivity and fast capacity decay. In this paper, a hydrangea-shaped hierarchically 3D magnesium hydride-carbon framework (MH@HyC) consisted of MgH2 nanoparticles (NPs) uniformly self-assembled on hierarchical porous carbon (HyC) is fabricated for advanced lithium storage. Featuring by high surface area and the well-defined macro-mesomicropore structure, HyC plays an ideal structure-directing role for the growth of MgH2 NPs with size-control, high loading, and hydrangea-shape array. Taking advantage of the robust 3D hierarchically porous structure and the derived interactions, MH@HyC not only provides sufficient electrochemically active sites, enhances the electronic conductivity and channels for rapid transfer of electrons/Li ions, but also relieves the agglomeration and accommodate the volumetric effects during cycling, leading to high capacity utilization, fast electrochemical kinetics, and well-sustained structural integrity. As a result, MH@HyC delivers a high reversible capacity of 554 mAh g−1 after 1000 cycles at a high current rate of 2 A g–1, enabling it a potential anode candidate for LIBs.

KEYWORDS: magnesium hydride; 3D hierarchically porous structure; self-assembly; porous carbon spheres; lithium storage


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INTRODUCTION In the past decades, we have witnessed the remarkable progress and wide application of lithiumion batteries (LIBs) in the field of portable electronic devices.

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To satisfy the ever-increasing

demands for high performance and safety, constant efforts have been devoted to developing various electrode materials with excellent rate performance, high specific capacity, as well as longtime cycling stability.4-6 Magnesium hydride (MgH2), by virtue of its high specific theoretical gravimetric and volumetric capacities (2038 mAh g−1 and 2878 mAh L−1), suitable voltage (0.58 V vs. Li+/Li0), natural magnesium resource abundance and eco-friendly, has drawn considerable attention as a potential promising conversion-type electrode material for LIBs.7, 8 However, the progress of MgH2 electrode for LIBs has been largely hindered by its low electronic/ionic conductivity and volumetric effects during lithiation, giving rise to a low capacity utilization and poor rate capability. In the last ten years, constant efforts have been devoted to addressing those issues by reducing the particle sizes of MgH2, compositing with carbon materials (graphite9-11, mesoporous carbon, and graphene12, 13) or metal additives (TiH214, 15 and TiO216), introducing effective binder (CMC-f)17, and fabricating solid state batteries.18-21 Among these efforts, to obtain MgH2 at nanoscale has demonstrated to be a relatively effective route to enhance the electrochemical performance of MgH2 due to the fast transport kinetics and sufficient accessible active sites of nanosized MgH2 compared with that of bulk materials. Unfortunately, Mg-based active nanograins are easily tend to self-grow, agglomerate and pulverize upon cycling, thus leading to a poor reversibility. For instance, Oumellal et al. 22 firstly prepared a series of MgH2 NPs on commercial graphite (HSAG-500) by a bottom up method for lithium storage. The obtained composite contained a max amount of 50% MgH2 with controlled size distribution. Nevertheless, it delivered a rapid capacity decay after a few cycles due to the agglomeration of MgH 2 particles


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and thus the fast loss of electronic contact with graphite. Recently, we have demonstrated that graphene with high surface area could be employed as an ideal conductive scaffold for anchoring MgH2 NPs and play an effective role of improving the electrode reversibility12, 13. Therefore, it is of great interest to develop other attractive conductive hosts with high surface area for the tailored self-assembly, uniform distribution, and confinement of MgH2 NPs to fully realize its lithium storage performance. Compared with common carbon materials, 3D hierarchically porous carbon materials (HPCs) demonstrated great superiorities in terms of low density, high surface area, superior conductivity, and excellent electrochemical stability, either as electrode materials itself or efficient scaffolds for LIBs and beyond.

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Particularly, as the desirable host for loading active materials, the

hierarchical porous networks ensure sufficient electrolyte penetration and efficient charge delivery to active materials, leading to the superior Li/electronic transport kinetics, and the full utilization of lithium storage capability. Moreover, the mechanically robust HPCs architectures are beneficial for alleviating the volumetric effects within electrode, thus enabling a long-term cycling stability. Specifically, HPCs spheres with micro-size and spherical morphology present perfect particle mobility and potential high packing density for the high volumetric energy and power density. 30, 31

For example, Zhu et al.,32 realized the uniformly distribution of ultrafine MOx (M=Sn, Mn) NPs

on 3D hierarchically porous carbon matrix as electrode materials for LIBs, which demonstrated high active materials utilization, impressive power density, and superior cycling stability. Therefore, inspired by both the superiorities of MgH2 and HPCs, it is highly desirable to construct composite which not only consist of MgH2 at nanoscale, but also involve robust 3D hierarchically porous framework so as to fully utilize the potential storage capability in terms of high specific capacity, high-rate performance, and enhanced cycling stability.


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In this paper, well-defined hierarchical porous carbon spheres (HyC), featuring by macro– meso–micropores and high surface areas, were synthesized and functioned as an ideal conductive scaffold for the uniform self-assembly of MgH2 NPs. Under the structure-directing role of HyC, MgH2 NPs with size-control, high loading, associated with hydrangea-shaped 3D hierarchically porous frameworks were fabricated (MH@HyC) and applied for lithium storage. Taking advantage of the robust 3D hierarchically porous structure and the interactions derived at nanoscale, the hydrangea-shaped MH@HyC exhibits high conductivity, sufficient accessible electrochemically active sites and abundant channels for rapid transfer of electrons/Li ions, which is largely beneficial for the fast electrochemical reaction and high capacity utilization. Moreover, the 3D hierarchically porous structure plays an effective role in accommodating the volumetric effects of active materials and alleviating its agglomeration, thus leading to a sustained structural stability. As a result, the obtained hydrangea-shaped MH@HyC was endowed with high reversible capacity, intriguing rate capability, as well as enhanced cycling stability, delivering a high specific capacity of 554 mAh g−1 at 2 A g–1 after 1000 cycles. RESULTS AND DISCUSSION Figure 1 schematically demonstrates the fabrication process of MH@HyC framework. Firstly, HyC was prepared using the structure-directing role of HyC. By adjusting the content of HyC, MH@HyC with different loading ratios of MgH2 (60, 70 and 80 wt %) were obtained and denoted as MH@HyC60, 70 and 80, respectively. The morphology and structure of the obtained HyC and MH@HyC were characterized by the field-emission transmission electron microscopy (FETEM), and the scanning electron microscopy (FESEM). Figure 2a-c show the well-defined 3D hydrangea-shaped porous structure of HyC with an average diameter of 400 nm (Figure S1). It consists of the radial large macropores (~200 nm), with coexist of mesopores (3.8 nm) and


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micropores (2.0 nm) decorated on the petal framework of HyC, as verified by the N2 adsorption/desorption tests (Figure S2). The Brunauer-Emmett-Teller specific surface area of the hierarchical porous HyC is about 623.4 m2 g−1, which makes it an ideal skeleton for the uniform distribution of MgH2 NPs. Taking MH@HyC60 for example, the obtained composite inherits the overall morphology of HyC with MgH2 NPs homogenously decorated on the surface of the hierarchical porous structure (Figure 2d-g). No obvious structural collapse or aggregation of MgH2 NPs can be observed, indicating a stable interaction between HyC and MgH2. Moreover, the MgH2 NPs mostly display a tailored and controlled particle size of ~10.8 nm (Figure 2h). As shown in the high-resolution TEM image (Figure 2i), the nanocrystals demonstrate the lattice fringes of MgH2 planes (200) with spacing of ~2.26 Å. This result is consistent with the selected-area electron diffraction (SAED) pattern of MgH2 (Figure 2j), indicating its polycrystalline structure. In addition, the scanning transmission electron microscopy (STEM) images, as well as the elemental mapping images of C and Mg (Figure 2k-m) further verify the homogeneous distribution of MgH2 on HyC. Besides, as the mass percentage of MgH2 rises from 70 wt% to 80 wt%, the density of MgH2 NPs anchored on HyC increases gradually and agglomerates severely (Figure S3), leading to a formation of composite without obvious porous framework. The crystallographic structure of MH@HyC was further identified by the X-ray power diffraction (XRPD). Compared with the typical XRPD pattern of HyC (Figure S4), MH@HyC60 (Figure 3a) confirms the generation of MgH2, with all diffraction peaks corresponding to the tetragonal MgH2 phase (JCPDS no. 12–0697). However, it can be noted that the presence of carbon largely decreases the cystalline size of MgH2 NPs, which exhibts nanocrystalline size and high purity. Further more, the surface composition of MH@HyC was tested by the X‐ray photoelectron


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spectroscopy (XPS). As shown, the obtained XPS spectrum (Figure S5a) and the high-resolution XPS spectrum of Mg 2p (Figure S5b) demonstrate the characteristic peak of MgH2 at 50.1 eV, which further confirms the formation of MgH2. There is also part of MgO and Mg(OH)2 detected due to the inevitable exposure to air during testing. Besides, the porosity traits of MH@HyC resulting from the 3D assembly of MgH2 NPs on the HyC was revealed by the N2 adsorption/desorption tests (Figure 3b). As shown, it exhibits a type IV curve which confirms both the mesoporous and macroporous characteristics. There is a relatively board pore size distribution, featured by coexist of microporous (1 nm), small mesopores (4 nm) and large mesopores (32 nm). The above results further verify the hierarchical porous structure of MH@HyC, which is consistent with the observation of TEM images. Moreover, to illustrate the electronic structure of the MgH2 and MH@HyC, the corresponding density of states (DOS) were calculated based on DFT calculation. According to Figure 3c, the pure MgH2 exhibits a relatively broad band gap (~2.15 eV), indicating a semi-conductor to insulator behavior. On the contrary, the MH@HyC demonstrates a metallic feature (Figure 3d), with no band gap around the Fermi level.33 Specifically, the partial DOS of MH@HyC shows that the conductor characteristic is mainly ascribed to the hybridized orbitals of Mg, H and C atoms, resulting in the change of electron density near Fermi level (rise of HyC and decline of MgH2) though the charge redistribution among Mg-C-H bonding. Therefore, the electrons in MH@HyC framework can easily flow along the HyC and rapid transfer to MgH2 by Mg-C-H bonding, which is largely beneficial for the electrochemical lithium storage process. All in all, the resulting 3D hierarchically porous networks of MH@HyC framework make it a potential promising electrode material for lithium storage.


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Figure 4a is the typical cyclic voltammetry (CV) profiles of MH@HyC60 electrode for the initial three cycles. During the first cathodic scan, the peaks below 0.70 V can be ascribed to the conversion of MgH2 to Mg and LiH (Equation. 1). MgH2 + 2Li+ + 2e- ↔ Mg + 2LiH

(1)

Mg + xLi+ + xe- ↔ LixMg

(2)

The broad peaks at 0.20-0.005 V should be ascribed to the alloying process of Mg with Li (Equation. 2). In the following anodic scan, peaks at low voltages (0.10-0.21 V) correspond to the delithizaiton of LixMg, and the peak centered at 0.62 V confirms the regeneration of MgH2. In the subsequent cycles, pairs of cathodic/anodic board peaks at about 0.21 V/0.20 V, and 0.62 V/0.46 V can be well detected and nearly overlapped, indicating a good reversibility. It is noted that there is a large irreversible peak at ~ 0.9 V in the first scan for the generation of stable solid electrolyte interphase (SEI) layer at the present of HyC (Figure S6), which can play an effective role in enhancing the cycle stability of the MgH2 electrode. Correspondingly, the obtained dischargecharge voltage profiles of MH@HyC are in good consistent with the CV results. As shown, Figure 4b exhibits pairs of slanted charge/discharge plateaus which can be ascribed to the nanostructured MgH2 particles. To evaluate the superiority of the MH@HyC, the galvanostatic charge/discharge performances were explored. Figure 4c displays the cycling performances of different MH@HyC electrodes at 100 mA g−1. The specific capacities were measured by weight of active materials. The performance of the bulk MgH2 (Bulk MH) and the blank HyC electrodes were also measured under the same conditions as comparison. According to Figure 4c, the initial discharge capacity of the bulk MgH2 is 1000 mAh g−1, while its capacity fades rapidly with a low capacity (172 mAh g−1) retains in 20 cycles. On the contrary, MH@HyC60 electrode delivers a high initial discharge/charge capacities


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of 2370/1137 mAh g−1. The capacity loss can be mainly ascribed to the irreversible side-reactions and the generation of SEI layer. After 100 cycles, it maintains a reversible capacity of 684 mA h g−1, corresponding to a Coulombic efficiency of 99.5%. In addition, the MH@HyC70 electrode with higher weight percentage of MgH2 displays a slightly decreased specific capacity and cycling stability. It is mainly attributed to the increased density of MgH2 NPs, which promotes the agglomeration of active particles as well as reduces the electronic conductivity of overall electrode. It is also noted that the pristine HyC delivers a low discharge capacity of 341 mAh g−1 at 100 mA g−1 after 100 cycles, and the capacity contribution of HyC on composite can be neglected, particularly at high current rates. Rate capability is an important index of merit of electrode for LIBs, most of which is limited by the slow Li/electron transport kinetics. Benefiting from the 3D hierarchically porous channels and the interpenetrating networks of the conductive carbon skeleton, the MH@HyC electrode can not only ensure highly efficient electrochemical reaction, but also boost electron /ion transport so as to promote the rate performance. Indeed, the MH@HyC electrode at variable current densities demonstrate impressive rate capability. Taking MH@HyC60 for example (Figure 4d), when cycled at different current rates of 100, 200, 500, 1000 and 2000 mA g−1, MH@HyC electrode delivers high reversible capacities of 1062, 964, 810, 687 and 578 mAh g−1, respectively. When the current rate recovers to 100 mA g−1, the corresponding capacity can restore to 1011 mAh g−1, suggesting a decent recoverability and rate capability. Moreover, even at a high current rate, MH@HyC60 also delivers desirable long-term cycling stability. According to Figure 4e, when cycled at 2 A g−1 for 1000 cycles, a high reversible capacity of 554 mAh g−1 can be retained, corresponding to a Coulombic efficiency of 98.5%. As a comparsion, the blank HyC only maintained a specific capacity of 360 mA h-1 at 2 A g-1 after 1000 cycles (Figure S7).


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The intriguing cycling stability and rate capability can be largely ascribed to the synergistic interactions of the 3D hierarchically porous MgH2-carbon framework at nanoscale. For one thing, the uniform distribution of MgH2 NPs on HyC matrix provides sufficient active sites for lithium storage and efficient electrons/Li-ions transfer. For another, the 3D hierarchically porous frameworks possess abundant interconnected conductive channels for sufficient penetration of electrolyte into the active sites across the entire network. The synergistic effect largely promotes the dynamics of the electrochemical reaction kinetic of MgH2, and enables the reactions that hardly occur for MgH2 in bulk state. Moreover, the robust frameworks play effective roles of alleviating volumetric effect and mitigating agglomeration of active components during cycling, which results in well-sustained structure integrity. The electrochemical impedance spectra (EIS) of MH@HyC60 electrodes before and after cycling were shown in Figure S8. The Nyquist plots of MH@HyC electrode contains two depressed semicircles at the medium-high frequencies, and an inclined straight line at lowfrequencies. The high-frequency semicircle and the middle-frequency semicircle corresponds to the interface resistances (Rf) formation on the surface of electrode and the charge-transfer resistance (Rct) of electrode reaction with Li ions, respectively. According to the equivalent circuits (inset of Figure S8), MH@HyC delivers relatively small values of Rf and Rct, which also decreases slightly after the activation process. For example, the fitting values of Rf before and after 100 cycles are 90.6 Ω and 75.7 Ω, indicating a stable SEI film and good electronic/ionic conductivity. Moreover, the inclined straight line at low-frequencies corresponding to the Warburg impedance resistance (Wo) further reflects a good solid-state of lithium ions diffusion kinetics within MH@HyC electrodes.


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The in operando synchrotron XRPD was applied to elucidate the structure and phase variation during the lithium storage process of MgH2 electrodes. However, due to the nanocrystalline size and poor cystalline of MH@HyC electrode, its diffraction peaks cannot be clearly identified in the cycling battery under the given condition (Figure S9). As a comparison, the in operando synchrotron XRPD in the cycling battery for bulk MgH2 was recorded. Figure 5a shows that the diffraction peaks corresponding to the (110), (101), (200) and (211) planes of MgH2 can be clearly identified, as well as the characteristic peaks of current collector (Cu). During discharge, the peak intensity of MgH2 represented by 17.5

o

with (200) plane gradually decreases, indicating a

conversion reaction process. Besides, the appeared Bragg peak at about 15.9 o can be ascribed to the (101) planes of the generated Mg based phase (d = 2.46 Å) in the electrode. However, it can be noted that quite amount of active materials remains unreacted with lithium, which is likely due to the poor electrochemical accessibility of bulk MgH2 electrode at the testing condition. Moreover, the peak intensity cannot recover fully in the charge process, indicating a poor electrochemical reversibility of bulk MgH2. The LiH phase was not detected in the XRPD patterns due to its small crystal size or amorphous phase at the cycling cell. The obtained results well explain the low capacity utilization, fast capacity fading and poor rate capacity of bulk MgH2 electrode aforementioned. To further evaluate the interactions between MH@HyC and Li ions, the ab-initio molecular dynamics (AIMD) simulation and DFT calculation was performed on simplified model MgH2 on single carbon layer. Combine with the AIMD, DFT simulation and Li adsorption results, it can be seen that H atoms of MgH2 can diffuse fast and form bonding with Li ions firstly (Figure 5b-c and Video S1). That is, during the lithiation process, Li ions can be absorbed onto MgH2 surface and forms Li-H bonding fast and firstly. It is much stronger than Mg-H bonding, thus leading to the


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formation of LiH. The process is self-efficacy, with the reaction energy calculated to be -1.0516 eV. Therefore, constructing MgH2-carbon framework with hierarchical porous at nanoscale can be an efficient way to promote the diffuse kinetics of H atoms, and thus beneficial for the electrochemical process of MH@HyC (Figure 5d). Besides, the electronic structure of MH@HyC with Li ions (Li+MH@HyC) were analysed based on the valence electron localization function (VELF). As shown in Figure 5e, the ELF maxima between C atoms in carbon layer with covalent bonding, and the ELF maxima localized electrons between H and Mg atoms indicates ionic bonding feature of MgH2. It is noteworthy that the ELF maxima between Li and H atoms distorts towards H atoms, suggesting strong ionic bonding interaction of Li-H bonds. Moreover, the charge density difference maps results of MH@HyC (Figure 5f) show that a charge accumulation regions appear at the interface between MgH2 and carbon layer. Besides, after adsorbing Li ions (Figure 5g), there is a pattern of strong charge accumulation between MgH2 layer and Li ions, whereas a charge depletion regions appear within MgH2 slab and near the top of Li ions. This suggested that a built can be established in electric field directed from Li ions to carbon based on the charge transfer from Li ions to MgH2, and then to carbon layer. Therefore, the HyC can actually serve as electron collector and promote the delivery of electron within the whole electrode. The results based on the simulation and calculation further verify the admirable interactions of Li+MH@HyC frameworks and the synergistic effect on the optimization of electrochemical lithium storage. The structural stability of MH@HyC electrode during cycling was characterized using TEM. According to Fig. 6a, after the first lithiation process, MH@HyC sample well retains its hierarchical porous framework, with active nanograins uniformly anchored on HyC. Besides, the HRTEM image (Figure 6b) demonstrates a series of multiple or intermediate nanocrystalline phases generated after the lithiation process. As shown, the detected d-spacing of 2.45~2.46 Å can


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be assigned to (110) planes of LixMg, resulting from the slightly expansion of lattice fringes after lithiation with wide composition ranges. Moreover, after long-time cycles at 200 mA g−1, MH@HyC electrode still preserves its overall hydrangea-shape without structural collapse. According to the SEM/TEM images (Figure 6c-d), MH@HyC is well wrapped by a SEI layer, with active nanograins homogeneously decorated on the surface. The corresponding STEM image and the elemental mapping (Figure 6e-g) further verify the homogeneous distribution of Mg element on C element (HyC), indicating a good structural integrity. Therefore, the robust 3D hierarchically porous frameworks, the derived admirable interactions, as well as the spatially confined lithiation/delithiation process play the synergistic roles of improving the lithium storage performance of MH@HyC. CONCLUSION In summary, hydrangea-shaped 3D magnesium hydride-carbon framework (MH@HyC) with hierarchical porous was fabricated by the self-assembly of MgH2 NPs on the hierarchical porous carbon spheres (HyC). Characterized by the high surface areas and periodic macro–meso– micropores distribution, HyC plays a favorable role of structure-directing for the bottom-up array of MgH2 NPs with size-control and high loading. When evaluated as anode material for LIBs, MH@HyC demonstrates high capacity utilization, enhanced electrochemical kinetic, and good structural stability. It is largely benefiting from the rational 3D hierarchically porous frameworks and the admirable derived interactions, which largely improves its electronic conductivity, provides sufficient active sites, and ensures confined lithiation/delithiation process spatially. As a result, the MH@HyC demonstrates high lithium storage capacity (a reversible capacity of 684 mAh g–1 at 100 mA g−1 after 100 cycles), admirable rate capability, and enhanced long-term cycling stabiltiy (a reversible capacity of 554 mAh g−1 at 2 A g–1 after 1000 cycles). Generally,


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this work reports an effective strategy towards the construction of MgH2 array with 3D framework for high-rate performance of lithium storage. And also, it is believed that the rational designed 3D hierarchically porous carbon spheres can serve as attractive and general scaffold for other nanomaterials in numerous applications. EXPERIMENTAL METHODS Material Preparation Synthesis of hierarchical porous carbon spheres (HyC): HyC were synthesized using hierarchical porous silica spheres (HP-SiO2) as hard templates. Firstly, HP-SiO2 were synthesized according to our earlier report.34 Typically, tetraethyl orthosilicate (0.5 mL) was added into the mixture of cetyltrimethylammonium bromide (0.32 g), absolute alcohol (20 mL), deionized water (45 mL), p-xylene (3.0 mL) and concentrated ammonia aqueous solution (4 mL 25-28%) under the mechanical stirring rate of 700 rpm. After reacting at 25 °C for 2 h, the HP-SiO2 were obtained by separating and washing with absolute alcohol for three times. The obtained silica spheres were then redistributed to 50 mM Tris-HCl buffer solution (50 mL) and 0.30 g dopamine hydrochloride was added later. After stirring at 25 °C for 12 h, the sample was separated and dried in vacuum, followed by calcination at 800 °C for 2 h with a ramping rate of 5 °C min-1 in Ar flow getting a silica/carbon hybrid. HyC were then obtained by treating the hybrid with 4 wt% HF solution (2 mL) for 24 h and washing with water for 5 times. Synthesis of MH@HyC: MgH2 NPs self-assembled on HyC were fabricated via the hydrogenation process of dibutyl magnesium. In a typical experiment, dibutyl magnesium solution (1.6 mL, 1 M in heptane) and HyC (27.7 mg) were mixed in in cyclohexane (40 mL) with a hydrogen pressure of 45 atm. The mixture was heated at 200 °C for 24 h. Then the obtained mixture was washed, centrifuged and dried, leading to the formation of HyC-supported MgH2 NPs. In addition, MH@HyC composites with different loading ratios (60 wt%, 70 wt % and 80 wt %, denoted as


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MH@HyC70 and MH@HyC80, respectively) were fabricated. As a comparison, the bulk or commercial MgH2 was used after treated by ball milling. Computational Methods. Density functional theory (DFT) calculations were performed using the VASP 5.4.4 code on Australian Synchrotron Compute Infrastructure (ASCI).35 ASSOCIATED CONTENT Supporting Information Available: Details about the computational methods, the additional XRD, SEM images, XPS, the electrochemical performances of relevant samples, and the ab-initio molecular dynamics (AIMD) simulation. 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] Author Contributions # These authors contributed equally. 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) and the Science and Technology Commission of Shanghai Municipality (17XD1400700). Part of experimental test was conducted at PD beamline, Australian Synchrotron (ANSTO).


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FIGURES

Figure 1. Schematic illustration of the fabrication process of MH@HyC framework.


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Figure 2. (a, b) SEM images and (c) TEM image of black HyC. (d-f) SEM images, (g, h) TEM images, (i) HRTEM image, (j) SAED pattern, and (k-m) STEM image and the corresponding elemental mapping of C and Mg for MH@HyC60.


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Figure 3. (a) The typical XRPD patterns of MgH@HyC. Inset of (a) is the tetragonal structure of MgH2. (b) The typical pore-size distribution of MH@HyC. Inset of (b) is the nitrogen adsorptiondesorption isotherm. The calculated Density of States (DOS) of bulk MgH2 (c) and MH@HyC (d).


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Figure 4. (a) Cyclic voltammetry profiles of MH@HyC60 electrode at a scan rate of 0.05 mV s– 1

. (b) Discharge and charge voltage curves of MH@HyC60 at a current density of 100 mA g–1. (c)

Cycling performances of MH@HyC, black HyC, and bulk MgH2 electrodes at a current rate of 100 mA g–1. (d) Rate capabilities of MH@HyC and black HyC. (e) Cycling performance of MH@HyC60 at a high current density of 2000 mA g−1.


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Figure 5. (a) The in operando synchrotron XRPD patterns for the cycling cell of bulk MgH2 electrode. (b-c) The ab-initio molecular dynamics (AIMD) simulation of MgH2 and Li ions. Colours of atoms: green (Li), dark yellow (Mg), pink (H), brown (C). (d) Schematic illustration of the lithiation/delithiation process of MH@HyC. (e) The valence electron localization functions (VELF) of MH@HyC with the adsorbed Li ions. The charge density difference map of (f) MH@HyC and (g) Li+MH@HyC system. Nudy blue and yellow regions represent the isosurfaces of electron depletion and accumulation.


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Figure 6. The structural evolution of MH@HyC electrode during cycling. (a) TEM image and (b) HRTEM image of MH@HyC at the first lithiation state. (c) SEM image, (d) TEM image, and (eg) the corresponding elemental mapping of MH@HyC electrode after 100 cycles at a current density of 200 mA g−1.


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Table of Contents


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Hydrangea-Shaped 3D Hierarchical Porous Magnesium Hydride

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