# In Situ Plating of Mg Sodiophilic Seeds and Evolving Sodium Fluoride Protective Layers for Superior Sodium Metal Anodes

In Situ Plating of Mg Sodiophilic Seeds and Evolving Sodium Fluoride Protective Layers for... IntroductionGrid‐scale energy storage devices are essential for renewable energy utilization to guarantee continuous energy harvesting and steady energy output.[1] Rechargeable Li‐ion batteries (LIBs) have been well developed for portable electronic devices and electric vehicles, but their ever‐increasing material price limits their applications for grid‐scale energy storage.[2] Sodium‐based batteries have been regarded as affordable alternatives to LIBs since Na exhibits similar electrochemical properties but much higher resource abundance as compared with Li.[3] With a high energy density of 1166 mAh g−1 and a low electrochemical potential of −2.714 V, sodium metal is the premium choice of anode material for various sodium‐based batteries with high energy density including sodium metal batteries, sodium‐sulfur batteries, and sodium air batteries.[4] Na metal anode, however, suffers from a few challenging problems which impede their practical applications. For example, the plating/stripping of Na metal with large volume changes would lead to the breaking up of solid‐electrolyte interphase (SEI) layers and exposure of fresh Na, which would generate new SEI layers with continuous consumption of Na metal as well as the electrolyte, together with increasing of the cell impedance due to the gradually thickening SEI.[5] The repeated exposure of fresh Na is also the cause of low Coulombic efficiency (CE), which is another obstacle to the practical application of Na metal anodes. Besides, the uneven plating/stripping of Na metal would lead to the formation of Na metal dendrites, which would eventually pierce the separator, causing serious safety hazards due to internal short circuits with thermal runaways, and even fires or explosions.[6] The Na metal dendrites are likely to be detached to produce “dead Na,” which is inactive for further electrochemical reactions and leads to loss of active Na.[7]Great efforts have been made to develop stable Na metal anodes.[8] As a result of the huge volume changes during cycling, tiny fractures on the interfaces would gradually propagate into large cracks, which could eventually destroy the interface and leave the Na metal without protection.[9] It was reported that constructing stiff interfaces on Na metal anodes that could survive the huge volume changes during cycling is a valid means to enhance the stability of Na metal anodes, and such stable interfaces can be achieved either by surface coating or through electrolyte engineering.[10] It was also reported that the metals exhibiting some solubility in Na, such as Be, Mg, and Sn could serve as sodiophilic nucleation seeds for Na metal to reduce the nucleation overpotential and realize uniform plating.[11] Li et al. reported Sn nanoparticles embedded in a carbon framework as preferable nucleation sites, which enabled stable cycling of Na metal anodes with high Coulombic efficiency. While the metal seeded plating had little effect on the properties of the SEI layers, the interface degradation could not be inhibited. Another feasible approach toward stable Na metal anodes is applying porous host materials, which could alleviate the huge volume changes derived from the hostless nature of Na.[12] The plating and stripping of Na metal within the porous hosts are significantly different from the hostless process. The porous hosts could sustain some voids to accommodate Na metal, which does not shrink after the thorough stripping of Na metal.[13] As a result, the SEI films can persist on the surfaces of the hosts. The conductive porous hosts also offer large surface areas as nucleation sites, which could dramatically reduce the local current density and thus realize more stable Na metal plating.[14]A variety of porous materials have been developed as hosts for Na metal anodes, among which, porous metals (Cu, Al, Zn, and Ni) with good electrical conductivity have been proven to be promising candidates.[15] For instance, porous Al was reported as a host for Na metal anodes, which could generate smooth interfaces, as well as stable cycling.[16] The metal hosts are generally costly and exhibit high mass density, which would weaken the advantages of Na metal anodes. In comparison, various carbon‐based materials with good electrical conductivity, rich porous structures, and low mass density are preferable as host materials for Na metal anodes.[17] Porous carbons, carbon nanotubes (CNTs), and graphene are all well‐investigated host materials for Na metal anodes.[10c,13a,15b,18] It was demonstrated that Na metal plating on CNT electrodes could generate a smooth surface, which enabled stable cycling of the Na metal anodes.[18a,19] Chen and co‐workers have reported a porous reduced graphene oxide (RGO) as a host for Na metal anodes and successfully assembled Na‐CO2 batteries, which sustained stable cycling for more than 50 cycles.[20] Although approaches through interface engineering, metal seeded plating, and host material design all have been proved viable for stable Na metal anodes, reasonable material design with synergistic effects to incorporate these merits in one multifunctional host has rarely been achieved.Herein, we report a simple approach via hydrothermal assembly and thermal reduction to prepare RGO aerogel anchored with MgF2 nanocrystals (Figure 1a). The aerogel structure with its 3D conductive skeleton and rich porous structure reduces the local current density and offers interconnected void spaces to accommodate the deposition of Na metal. Moreover, the MgF2 component plays a key role in achieving non‐dendritic and highly reversible Na plating/stripping. The MgF2 nanocrystals can be converted into Mg (nucleation sites) and NaF (solid electrolyte interphase) during the initial stage of the first Na plating. The Mg sodiophilic seeds guarantee uniform Na nucleation and growth. Meanwhile, the NaF‐rich protective layer evolves over plating/stripping cycles, which suppresses the Na dendrite growth and prevents continuous electrolyte depletion. As a result of the synergistic effect, the MgF2@RGO aerogel serves as a multifunctional host to regulate the uniform deposition of Na metal, enabling the Na anode with highly stable cycling stability. Furthermore, the sodium metal full cells coupled with the Na metal anode confined in the MgF2@RGO aerogel host and Na3V2(PO4)3 (NVP) cathode deliver enhanced cycling stability and rate capability, suggesting its great potential for practical applications.1Figurea) Illustration of the synthesis of the MgF2@RGO aerogel; b) illustration of the working mechanism of MgF2@RGO aerogel as a host material for Na metal anodes.ResultsThe aerogels were synthesized by the hydrothermal method, followed by a freezing‐dry and a calcination process. The aerogels as‐synthesized were ≈10 mm in diameter and ≈5 mm in thickness, and the weight for each aerogel cake was ≈1.3 mg (Figure S1, Supporting Information). The lightweight aerogel host could accommodate sufficient Na metal without sacrificing the gravimetric energy density of the composite Na metal anode, as compared with the heavy porous metal‐based hosts.[16] The aerogels exhibit an interconnected porous structure (Figure 2a), which was most likely derived from the self‐assembly of the graphene layers during the hydrothermal treatment. The ethanediamine and sodium borate could interact with the functional groups on the graphene oxide (GO) layers during the hydrothermal treatment to prevent the layers from stacking with each other (Figure 1a).[21] The high‐resolution scanning electron microscopy (SEM) image of the aerogel (Figure 2b) demonstrates that nanocrystals with ≈20 nm in diameter were uniformly distributed on the graphene layers. The transmission electron microscope (TEM) image (Figure S2, Supporting Information) confirms the structure of the nanocrystals grown on thin graphene layers. The high‐resolution TEM image (Figure 2c) of the nanocrystals on the layers exhibits distinctive lattice fringes with a lattice spacing of 0.327 nm, which corresponds to the d‐spacing of (110) lattice planes for the MgF2 crystals. The energy‐dispersive X‐ray spectroscopy (EDS) mapping images (Figure 2d) indicate that the aerogel layers feature uniformly distributed C, Mg, and F elements, suggesting the uniform distribution of the MgF2 nanocrystals on the graphene layers. The X‐ray diffraction (XRD) pattern (Figure 2e) of the sample can be indexed to MgF2 (JCPDS No. 41–1443), and the broad peak at ≈28° is similar to that of the RGO sample, which should be derived from the graphene layers. Therefore, it is reasonable to deduce that the as‐synthesized aerogel samples consisted of a porous RGO matrix with anchored MgF2 nanocrystals (MgF2@RGO). The nitrogen adsorption isotherms (Figure 2f) indicate that the Brunauer–Emmett–Teller (BET) surface area of the MgF2@RGO aerogel (≈153 m2 g−1) is slightly smaller than that for the RGO aerogel (≈206 m2 g−1), which could have resulted from the reduced surface area due to the loading of MgF2 nanocrystals. The pore size distributions in Figure S3a, Supporting Information, also indicate that there were numerous nanopores of 5–10 nm in size, which are favorable for the penetration of electrolytes and transportation of the Na+ ions. The high surface area of the MgF2@RGO and RGO aerogels would decrease slightly (Figure S3b, Supporting Information) after compression due to battery assembly, while the amount of nanopores in the compressed aerogels would increase as a result of the enhanced contacting of the graphene layers (Figure S3c, Supporting Information). The X‐ray photoelectron spectroscopy (XPS) survey spectrum (Figure 2g) and high‐resolution C 1s, Mg 1s, and F 1s spectra (Figure S4, Supporting Information) confirm the presence of Mg, F, C, and O elements in the MgF2@RGO aerogel sample, as compared with the RGO sample, which is mainly composed of C and O elements. The content of the MgF2 on the aerogel was determined to be ≈22.5% by thermogravimetric analysis (TGA) (Figure 2h). The Raman spectra (Figure S5, Supporting Information) demonstrate that the intensity ratios (based on peak area) of the defect‐induced D band and crystalline graphite‐derived G band (ID/IG) for the MgF2@RGO and RGO aerogels were 1.48 and 1.37, respectively. The high ID/IG ratios indicate that the aerogels were highly defect‐rich, which should be attributed to the exfoliation of the graphene layers by water molecules via shear force from the ultrasonication treatment during the synthesis. The exfoliation of graphene to produce more graphene layers is favorable for the construction of the aerogel architecture and the loading of MgF2 nanocrystals. These characterizations indicate that the as‐synthesized MgF2@RGO aerogel exhibits thin and large‐area graphene layers, interconnected porous structures, and uniformly distributed MgF2 nanocrystals.2Figurea) SEM image, b) high‐resolution SEM image, c) high‐resolution TEM image, and d) EDS mapping images for the MgF2@RGO samples; e) XRD pattern, f) nitrogen adsorption isotherms, and g) XPS survey spectra for the RGO and MgF2@RGO samples; h) TGA curve for the MgF2@RGO sample.The MgF2@RGO and RGO aerogels were used as free‐standing electrodes to evaluate their viability as host materials for Na metal anode. The average Coulombic efficiencies (ACE) for the Na metal anodes with the bare Cu, RGO, and MgF2@RGO aerogel hosts (denoted as Na/Cu, Na/RGO, and Na/MgF2@RGO) at the current density of 0.5 mA cm−2 are presented in Figure 3a. The ACE in the first 20 cycles for the Na/MgF2@RGO electrode was 96.24%, which was higher than the ACE for the Na/RGO electrode (92.62%) or the Na/Cu electrode (84.73%), and also better than many of the state‐of‐the‐art reports.[9] The ACE for the Na metal anodes with the MgF2@RGO and RGO aerogel hosts (Figure S7a, Supporting Information) at the current density of 1 mA cm−2 was 94.08% and 91.30%, which were slightly lower than that at 0.5 mA cm−2. The decreased ACE indicates deteriorated cycling stability at the higher current density. As for the Na metal anode with the bare Cu, the ACE at 1 mA cm−2 was decreased to 82.32%, suggesting much inferior cycling stability of the Na metal anode. The CE for the Na/Cu electrode at 0.5 mA cm−2 (Figure 3b) was only around 80% in the first 100 cycles, but it seriously deteriorated afterward, with randomly scattered values ranging from 10% to 120%. The CE for the Na/RGO electrode at 0.5 mA cm−2 was initially lower than 60% and gradually increased to ≈92% after more than 40 cycles, which is likely due to the establishment of relatively stable SEI layers. The CE became unstable, however, with a clear trend of degradation after 100 cycles and even decreased to ≈20% after 300 cycles. In comparison, the CE for the Na/MgF2@RGO electrode quickly reached more than 96% within 20 cycles, and it remained steady for more than 300 cycles. The voltage curves (Figure S6, Supporting Information) at different cycles of the Na metal anodes with the MgF2@RGO, RGO, and Cu hosts during the CE testings are in accordance with the CE results in Figure 3b, demonstrating that the Na metal anodes with the MgF2@RGO aerogel host exhibited excellent cycling stability. When the current density was increased to 1 mA cm−2, the CE of the Na metal anodes with the MgF2@RGO aerogel host (Figure S7b, Supporting Information) remained stable for ≈100 cycles, which was dramatically superior to that for the RGO and bare Cu hosts. It can be inferred that the enhanced CE for the Na/MgF2@RGO electrode must have resulted from the uniform Na plating with more stable SEI layers, as a result of the introduced MgF2 nanocrystals. The CEs of the Na metal anodes at the current density of 0.5 mA cm−2 and areal capacity of 2 mAh cm−2 deliver similar results (Figure S8, Supporting Information), suggesting the superior cycling stability of the Na metal anodes with the MgF2@RGO aerogel host. The rate capabilities of the Na/Cu, Na/RGO, and Na/MgF2@RGO aerogels were also evaluated in symmetrical cells, as presented in Figure 3c. It can be seen that the Na/Cu electrode could achieve stable cycling at the current density of 0.1 to 2 mA cm−2, while the voltage hysteresis at 2 mA cm−2 was as large as 0.4 V, which is much higher than those for the Na/RGO and Na/MgF2@RGO electrodes at the current density of 5 mA cm−2. This indicates that the 3D interconnected conductive structures of the aerogels could effectively decrease the local current density during Na plating/stripping at high current rates, so as to deliver smaller voltage hysteresis as compared to the 2D planar bare Cu. The Na/RGO and Na/MgF2@RGO electrodes exhibited stable cycling with the current density ranging from 0.1 to 5 mA cm−2, and they also delivered stable cycling for ≈100 cycles when the current density was recovered from 5 to 0.5 mA cm−2. It is noticeable that the voltage hysteresis for the Na/MgF2@RGO electrode was slightly smaller than that for the Na/RGO electrode, which is similar to the results in previous reports, suggesting that the SEI layers on the Na/MgF2@RGO electrode may be thinner with enhanced interface stability.[7,19]3Figurea) Average Coulombic efficiencies at 0.5 mA cm−2, b) Coulombic efficiencies at 0.5 mA cm−2 and 0.5 mAh cm−2, and c) cycling performances at various areal current densities of the Na metal anodes with the Cu, RGO, and MgF2@RGO hosts; nucleation potentials of Na metal plating on the d) MgF2@RGO, e) RGO, and f) Cu hosts at various current densities; g) summary of the Na plating overpotentials on the various hosts; h) cycling stabilities of the Na metal anodes on the Cu, RGO, and MgF2@RGO hosts at 0.5 and 0.5 mAh cm−2.The nucleation overpotential for the Na metal plating, which is defined as the voltage gap between the dip and the plateau on the voltage curve of the first plating process, was also investigated for the host materials at various current densities. The nucleation overpotentials (Figure 3d) with the MgF2@RGO host at 0.1, 0.2, 0.5, 1, and 2 mA cm−2 were 27.7, 38.4, 45.2, 72.0, and 92.6 mV, respectively, which are smaller than those for the Na plating on the RGO host (Figure 3e). In comparison, the nucleation overpotentials for the Na plating on the bare Cu electrodes (Figure 3f) at 0.1, 0.2, 0.5, 1, and 2 mA cm−2 were 154.4, 163.8, 223.7, 292.9, and 307.3 mV, respectively. The large nucleation overpotentials for the planar Cu electrodes (Figure 3g) further demonstrate the advantage of the aerogels with 3D conductive porous structures, which could increase the affinity between Na and the aerogels and reduce the local current density. The lowest nucleation overpotentials for the MgF2@RGO host were most likely derived from the synergetic effect of the conductive graphene aerogel and the anchored MgF2 nanoparticles. The latter could be converted into Mg as nucleation seeds for Na metal plating to further decrease the overpotentials. Long‐term cycling of the Na metal anodes with the Cu, RGO, and MgF2@RGO hosts was performed at 0.5 mA cm−2 with an initially plated 5 mAh of Na as the working electrode (Figure 3h). The Na/Cu electrode exhibited a dramatic increase of the voltage hysteresis from ≈0.1 V to more than 1 V after stable cycling for ≈300 h, which may be derived from the complete consumption of the pre‐deposited Na on the working electrode after cycling because of the continuous formation of “dead” Na. The unstable cycling performance of the Na/Cu electrode is in accordance with its CE performance, as presented in Figure 3b, further demonstrating that the Na metal anodes on planar Cu exhibited inferior cycling stability. Although the cycling stability for the Na/RGO electrode was enhanced to ≈650 h, a similar exacerbation of the voltage hysteresis occurred thereafter, which suggests deteriorated interfaces of the Na metal anodes after cycling. In contrast, the Na/MgF2@RGO electrode delivered smaller voltage hysteresis and much more stable cycling performance for more than 1600 h, which is more than two times longer than that for the Na/RGO electrode. The plating/stripping curves of the Na metal anodes with the Cu, RGO, and MgF2@RGO aerogel hosts at different cycling times (Figure S9, Supporting Information) demonstrate that the voltage curves for the Na/MgF2@RGO are stable during cycling, as compared with the inferior stability for the Na/Cu and Na/RGO anodes with gradually increased voltage tips. Furthermore, the Na/MgF2@RGO anodes could deliver stable cycling for ≈500 h at 1 mA cm−2 (Figure S10a, Supporting Information) and ≈300 h at 2 mA cm−2 (Figure S10b, Supporting Information), which are both better than the Na metal anodes with the bare Cu and RGO hosts. The excellent cycling stability of the Na metal anodes with the MgF2@RGO hosts is superior to those in many of the state‐of‐the‐art reports (Figure S11, Supporting Information), which further demonstrates the advantages of this host material for Na metal anodes.The morphology evolution of the Na metal anodes with the bare Cu and aerogel hosts was also investigated. The morphologies for the pristine MgF2@RGO (Figure S12a, Supporting Information) and RGO (Figure S12b, Supporting Information) are similar, showing relatively flat surfaces derived from randomly assembled RGO layers. The bare Cu exhibits an even but rough surface (Figure S12c, Supporting Information). As presented in Figure 4a, the Na metal initially plated on the MgF2@RGO aerogel exhibits a nodule‐like structure, which is similar to that for the RGO aerogel in Figure 4e. However, the Na nodules on the MgF2@RGO aerogel are in intimate contact and form large lumps with reduced surface area, as compared with the much smaller and more separated nodules which present a porous structure (Figure 4e). Similarly, many nodules can be seen in the case of the Na metal plating on the bare Cu (Figure 4i), while there are also some Na whiskers and the surfaces of the nodules were cracked, indicating inferior interface stability. After 200 cycles, the Na metal plating on the MgF2@RGO aerogel (Figure 4b) exhibits a dense and flat surface, while the surface for the Na/RGO (Figure 4f) is rough and uneven. The Na metal on the bare Cu after 200 cycles (Figure 4j) shows much smaller Na nodules and even mossy Na, which is most likely derived from the cracking of large Na nodules, suggesting the poor cycling stability of the Na metal on bare Cu. When the plated Na metal was stripped from the MgF2@RGO aerogel after the 200th cycle, a flat and uniform interface could be achieved (Figure 4c). From this, it can be deduced that the MgF2@RGO aerogel could significantly enhance the cycling stability of the Na metal anodes, as a result of the 3D porous conductive matrix and MgF2‐derived stable interfaces (Figure 4d). As for the surface of the stripped Na/RGO electrode (Figure 4g), irregular‐shaped island‐like Na lumps were retained, which are most likely detached inactive “dead Na” as a result of the poor stability of the SEI layers (Figure 4h). As for the bare Cu (Figure 4k), mossy‐like Na was retained with noticeable cracks on the electrode after the 200th stripping. Its inability to be stripped from the bare Cu indicates that the mossy Na is inactive, implying that it is detached from the current collector as “dead Na.” The continuous accumulation of dead Na on the bare Cu (Figure 4l) indicates that the Na metal on bare Cu exhibits poor cycling stability. The cross‐section SEM images and the digital photographs of the cycled Na metal anodes with the bare Cu, RGO, and MgF2@RGO aerogel hosts are presented in Figures S13 and S14, Supporting Information. It can be seen that the Na metal is preferably deposited in/onto the MgF2@RGO aerogel matrix (Figure S14a, Supporting Information) with dense and compact morphology and little volume change upon cycling (Figure S13a,b, Supporting Information), indicating good sodiophilicity of the aerogel host and superior stability of the composite Na metal anode. In contrast, the Na/RGO (Figure S14b, Supporting Information) presents a little amount of Na metal plated onto the battery shell instead of the host, and the volume change (Figure S13c,d, Supporting Information) upon cycling is larger. As for the Na/Cu electrode (Figure S14c, Supporting Information), a considerable amount of Na metal is deposited on the battery shell, and the thickness of the Na metal on the bare Cu was decreased due to more Na being deposited onto the battery shell upon cycling. Therefore, the morphology evolution clearly demonstrates that the RGO aerogel with the 3D conductive matrix can enhance the stability of Na metal to some extent as compared to the planar Cu. Significantly, the MgF2@RGO aerogel host could achieve dramatically improved Na metal plating, as a result of the synergy of MgF2 nanocrystals and RGO aerogel, which could stabilize the interfaces.4FigureSEM images of the Na metal anodes with MgF2@RGO host after a) the 1st plating, b) the 200th plating, c) the 200th stripping, and d) corresponding illustration of the uniform Na plating on the MgF2@RGO host. SEM images of the Na metal anodes with the RGO host after e) the 1st plating, f) the 200th plating, g) the 200th stripping, and h) corresponding illustration of the Na plating on the RGO host with dead Na growth. SEM images of the Na metal anodes on bare Cu after i) the 1st plating, j) the 200th plating, k) the 200th stripping, and l) corresponding illustration of the Na plating on the bare Cu with mossy dead Na growth. Scale bars: 20 µm.To achieve an in‐depth understanding of the working mechanism of the MgF2@RGO aerogel as host material for Na metal anodes, experimental and theoretical investigations were further conducted. During the first Na plating into the MgF2@RGO aerogel host (Figure S14, Supporting Information), two voltage slopes can be seen on the voltage curve. It is noticed that the electrode delivers a high initial capacity of ≈1000 mAh g−1, which should be ascribed to the Na+ storage into the RGO matrix, the conversion of MgF2 nanocrystals, and also the formation of SEI layers. Phase evolution of the MgF2@RGO aerogel host at various stages of Na metal plating (Figure 5a) was studied via ex situ XRD analysis, as presented in Figure 5b. Initially, at stage a with no Na plating, the host presents a characteristic broad peak that corresponds to the diffraction of RGO and also the diffraction peaks for MgF2. At stage b when the voltage has decreased to 0 V with intercalation of Na+, the diffraction peaks for MgF2 have vanished, and new peaks for Mg and NaF have appeared. The phase change from stages a to b is in line with the voltage features in Figure S15, Supporting Information, suggesting that the MgF2 nanocrystals were converted into Mg and NaF during the Na plating. The voltage dip between stages b and c corresponds to the nucleation of Na on the MgF2@RGO aerogel host, after which, the diffraction peaks for metallic Na gradually emerge and become stronger from stages c to e with continued Na metal plating. When the plated Na metal was stripped from the aerogel host at stage f, the diffractions for metallic Na disappeared, which indicated that most of the Na had been completely stripped, with no remaining “dead” metallic Na. Diffraction peaks for Mg and NaF can also be detected, suggesting that the Mg and NaF can be preserved during the Na plating/stripping process.5Figurea) Various stages on the potential curve for the first Na metal plating/stripping on the MgF2@RGO host, and b) corresponding evolution of the XRD patterns during the plating/stripping process; c) illustration of the binding sites for Na on the Mg/graphene surfaces, and d) corresponding binding energies based on the DFT calculations.The surfaces of the electrodes were also investigated via XPS analysis. As presented in Figure S16, Supporting Information, the Na 2s peaks at 1070 and 1070.4 eV for the cycled Na/MgF2@RGO electrode (Figure S16a, Supporting Information) can be assigned to metallic Na and NaF (and/or Na‐O), respectively.[22] The fitted F 1s peak at 683.3 eV (Figure S16b, Supporting Information) is corresponding to NaF. Similar peaks assigned to NaF can also be observed on the cycled Na/RGO electrode (Figure S16c,d, Supporting Information), which should be originated from the fluoroethylene carbonate (FEC) derived SEI layers. It is noteworthy that the cycled Na/MgF2@RGO electrode shows a much stronger NaF signal than that of the Na/RGO electrode, which implies additional NaF has been generated from the conversion of MgF2 in the cycled Na/MgF2@RGO electrodes. To explore the feasibility of Na plating on the surfaces of Mg and RGO, density functional theory (DFT) calculations were performed to determine the binding energy of Na on the (101) plane for Mg and the graphene plane for RGO. Based on the calculations, the preferred binding sites are illustrated in Figure 5c, and the corresponding binding energies are presented in Figure 5d. The binding energies for Na adsorption on Mg are −0.20, −0.19, and −0.14 eV (Figure S17, Supporting Information), suggesting that the surface of Mg is favorable for Na adsorption, while the binding energies for the graphene layer are 0.38, 0.61, and 0.72 eV, indicating that the process of Na adsorption on the graphene layers requires additional energy, which is unfavorable.Based on these investigations, it is rational to infer that the MgF2 nanocrystals on the MgF2@RGO aerogels can be converted into Mg and NaF during the Na plating process (Figure 1b), and the Mg nanocrystals could alloy with metal Na to serve as nucleation sites and achieve stable Na plating. The feasibility of Mg for nucleation sites has also been proved by recent research, and NaF has also been well recognized as a preferred component of SEI layers to construct NaF‐rich stable interfaces.[11a,12,23] Besides, 3D RGO aerogels significantly reduce the local current density and accommodate the deposited Na. Therefore, the working mechanism of MgF2@RGO aerogel as a host material for Na metal anodes can be summarized as the synergetic effects of the 3D conductive matrix, Mg‐induced sodiophilic plating, and NaF‐generated stable SEI layers. The outstanding performance of the Na metal anodes in the MgF2@RGO aerogel host with enhanced CE, improved interface stability, and ultra‐stable long‐term cycling can be attributed to the synergetic effects derived from the multifunctional aerogel host.To further verify the feasibility of the aerogel hosts for practical Na metal anodes, Na metal full cells with NVP cathode were investigated. It can be seen that, at the low current rate of 0.2 C, the full cells with the Na/RGO and Na/MgF2@RGO anodes delivered a comparable discharge capacity of 110 mAh g−1 (Figure 6a), while at higher rates of 0.5, 1, 2, and 5 C, the full cells with Na/MgF2@RGO anode delivered enhanced rate capability as compared with that for the Na/RGO anode. When the current density was recovered to 1 C, the full cells with Na/MgF2@RGO anode delivered stable cycling for more than 100 cycles with a decent discharge capacity of ≈96 mAh g−1. The full cells with Na/RGO anode exhibited a slightly lower discharge capacity of ≈90 mAh g−1, and the cycling gradually deteriorated after 100 cycles. It is reasonable to deduce that the inferior rate performance and cycling performance should be attributed to the instability of the Na/RGO anode. Charge/discharge curves of the full cells (Figure 6b) suggest that the voltage hysteresis gradually increased at higher rates. The full cells delivered comparable voltage plateaus at various current rates, while the voltage hysteresis of the full cells with the Na/RGO anode occurred at smaller specific capacities than those with the Na/MgF2@RGO anode. Figure 6c shows that the full cell with the Na/RGO anode delivered stable cycling at 1 C for 200 cycles with a high capacity retention of 91.3% (79.88 out of 87.48 mAh g−1). The cycling of the full cell with the Na/RGO anode dramatically deteriorated after ≈90 cycles, with the discharge capacity quickly decreasing from 79.5 to 38.8 mAh g−1. The poor cycling stability of the full cells should be attributed to the worse stability of the Na metal accommodated in the RGO host, which may be exhausted as a result of inferior interface stability. In Figure 6d, electrochemical impedance spectroscopy (EIS) analysis of the cycled full cells indicates that the impedance spectra with the Na/MgF2@RGO anode were stable during cycling, suggesting good electrode stability with constant interfaces. As compared to the full cells with the Na/RGO anodes, the semicircles corresponding to the charge transfer resistance gradually increased in size, implying progressive thickening of the SEI layers on the anodes. Furthermore, second small semicircles were identified on the EIS curves of the full cells with the Na/RGO anodes after 20 and 100 cycles, indicating the build‐up of new electrode interfaces, which could be resulted from the accumulation of “dead Na” on the anode surfaces. The potential of the aerogel host materials for practical Na metal full cells is further investigated with high‐loading NVP cathodes. As demonstrated in Figure 6e, even at the critical condition with a high cathode loading of ≈15 mg cm−2 and a low N/P ratio of ≈2.5, the Na metal full cells with the MgF2@RGO aerogel host could deliver stable cycling for more than 30 cycles, which is significantly superior to that for the RGO host with only 15 cycles of stable cycling. The charge/discharge curves (Figure 6f) for the Na metal full cells with the high loading NVP cathode and Na/MgF2@RGO anode at the 1st, 5th, 10th, and 20th cycles are nearly overlapped, suggesting good cycling stability under practical conditions. Thereby, it can be confirmed that the stability of the Na metal anode was critical for the stable cycling of the Na metal full cells, and the MgF2@RGO aerogel host could significantly increase the stability of the Na metal anode to achieve dramatically enhanced full cell performance.6Figurea) Rate performance, b) charge/discharge curves, c) cycling performance, and d) EIS evolution of the NVP‐based sodium metal full cells with the Na/MgF2@RGO and Na/RGO anodes, with the equivalent circuit shown in the inset. e) Cycling performance and f) charge/discharge curves at 1st, 5th, 10th, 20th, and 30th cycles for the high loading NVP cathode‐based Na metal full cells with the Na/MgF2@RGO and Na/RGO anodes.ConclusionA robust MgF2@RGO aerogel has been developed as a multifunctional host material for Na metal anode. Via in situ electrochemical conversion of MgF2, Mg sodiophilic sites, and NaF protective layers can be simultaneously planted on the RGO matrix, which can synergistically regulate the Na plating and achieve a superior long‐cycling lifetime. The morphology evolutions confirm that the combination of Na metal anode with the MgF2@RGO aerogel suppresses the formation of mossy Na and dendrite growth. The sodium metal full cells delivered enhanced rate and cycling performances, verifying the feasibility of the sodium metal anode for practical applications. The concept of host material design with synergistic effects of a porous conductive matrix, metal‐derived sodiophilic nucleation, and inorganic‐rich stable interfaces has been proved to be effective to achieve stable Na metal anodes, which sheds light on the future development of high‐energy sodium metal batteries.Experimental SectionPreparation of the MgF2@RGO AerogelsAll the reagents were of analytical grade and used as received without further purifications. GO solution was prepared by the modified Hummer's method, as reported previously.[24] Typically, 15 mL of GO solution (concentration ≈5 mg mL−1) was added into a 100 mL beaker with 25 mL of deionized water (DI), and the solution was stirred for 15 min before being treated under ultrasonication for 30 min. Thereafter, 2 mL of Mg(NO3)2 aqueous solution (160 mg mL−1) and 4 mL of NaF aqueous solution (40 mg mL−1) were dropped into the GO solution in sequence, and the mixture was kept under stirring for 1 h. Then, 160 µL of ethanediamine and 400 µL of sodium borate (2 wt.%) were dropped into the above solution and stirred for 15 min. Afterward, 1 mL of the mixture was placed into a 5 mL vial and sealed with a silica gel stopper, and the vials and 5 mL of DI water were then put into a 100 mL Teflon‐lined stainless‐steel autoclave and maintained at 180 °C for 18 h before being cooled to room temperature naturally. The as‐synthesized hydrogel was carefully transferred into a beaker with 490 mL DI water and 10 mL ethanol for distillation, and the distillation was maintained for 12 h and repeated four times. The preparation of GO hydrogel was conducted with the identical procedure, except for the addition of Mg(NO3)2 and NaF aqueous solution. The MgF2@GO and GO hydrogels were frozen in a refrigerator and then freeze‐dried. Thermal treatment was conducted at 500 °C for 2 h under an Ar atmosphere with a temperature ramp of 5 °C min−1 to obtain the MgF2@RGO and RGO aerogels. The aerogels as‐synthesized were ≈10 mm in diameter and ≈5 mm in thickness. The thickness of the aerogels after compression for battery assembly would be decreased to ≈0.5 mm, while the diameter of the aerogels would be retained. The weights for each MgF2@RGO and RGO aerogel were ≈1.3 and ≈1 mg, respectively.Materials CharacterizationsSEM images and EDS maps of the samples were collected with a field‐emission scanning electron microscope (FE‐SEM, Hitachi Regulus‐8230, Japan) equipped with an Ar‐filled sample transfer box. TEM images and high‐resolution TEM images were taken with a TEM Tecnai F30 (FEI) with an accelerating voltage of 300 kV. XRD patterns were recorded on an X‐ray diffractometer (Rigaku SmartLab) with Cu Kα radiation (λ = 0.15418 nm) at a voltage of 45 kV and a current of 200 mA. Raman spectra were recorded on an ARCSP 2558 spectrometer (Princeton Instruments). TGA was performed on a Setaram instrument (SETSYS EVO 18) from 30 to 700 °C under an air atmosphere with a heating rate of 5 °C min−1. The BET surface area was measured using N2 adsorption/desorption via an automated surface area analyzer (ASAP2420‐4MP). XPS spectra were recorded (Thermo Scientific ESCALAB 250Xi) with an Al Kα radiation source. For the XPS analysis of the cycled electrodes, the batteries were cycled at 0.5 mA cm−2 for ten cycles to establish stable SEI layers. Sample preparations for the characterizations of the cycled electrodes (SEM, XRD, and XPS) were carried out in an Ar‐filled glovebox, where the electrodes were obtained from the disassembled batteries, washed with the electrolyte solvents (EC/DEC, v/v = 1:1, with 5% FEC), dried in the chamber of the glovebox, placed in air‐tight chambers (SEM/XPS) with Ar atmosphere or protected with Kapton tape (XRD), and then transferred into the facility for characterizations. Sample preparations for the BET analysis of the compressed aerogel hosts were similar, except for using the electrolyte solvents instead of the electrolyte for the battery assembly.Electrochemical MeasurementsCoin‐type cells were tested at 25 °C on a Neware battery test system (CT‐4008T‐5V10mA‐164, Shenzhen China). The aerogels were directly used as free‐standing electrodes without further treatment. The coin‐type cells (LIR2025) were assembled in an argon‐filled glovebox (moisture and oxygen < 0.1 ppm) with Whatman glass fiber as the separator, Na foils as both counter and reference electrodes, and 100 µL of electrolyte composed of 1 m NaClO4 in ethylene carbonate (EC) and diethyl carbonate (DEC) (1:1 by volume), with 5% FEC as additive. The symmetrical Na cells were tested within a protective voltage range of −4 to 4 V. For the cross‐section morphology analysis of the cycled Na metal anodes with MgF2@RGO or RGO aerogel hosts, the batteries were assembled with an additional layer of Cu foil to support the samples. To evaluate the Coulombic efficiency, the batteries were first pre‐cycled at 0–1 V for five cycles (Figure S18, Supporting Information) to eliminate contaminations and establish stable SEI layers,[18c] then plated for 1 h and stripped to 0.5 V (vs Na+/Na) at the current density of 0.5 or 1 mA cm−2. For symmetrical Na||Na cell tests, 5 mAh cm−2 of Na was pre‐deposited onto the working electrodes (Cu, RGO, and MgF2@RGO), and the diameters of the Na counter/reference electrode and the Cu current collector were 16 and 12 mm, respectively. Both charging and discharging were conducted at the current density of 0.5, 1, or 2 mA cm−2 and cut off at a fixed time of 60 min. The ACEs of the Na metal anodes were tested following protocols similar to those in previous reports.[9] Briefly, the current collectors were first activated at the current density of 0.5 mA cm−2 by plating 5 mAh of Na and then stripping it to 1.2 V. Thereafter, 5 mAh (Qp) of Na was plated on the activated current collectors for working electrodes at the current density of 0.5 mA cm−2, then the working electrodes were cycled at the current density of 0.5 mA cm−2 and capacity of 0.5 mAh cm−2 (Qc) for n cycles (n = 20), and finally the residue Na was stripped to 1.2 V (Qs). The ACE was obtained according to the following equation:1ACE=(Qs+nQc)/(Qp+nQc)$\begin{array}{*{20}{c}}{{\rm{ACE}} = \left( {{{\rm{Q}}_{\rm{s}}} + {\rm{n}}{{\rm{Q}}_{\rm{c}}}} \right)/\left( {{{\rm{Q}}_{\rm{p}}} + {\rm{n}}{{\rm{Q}}_{\rm{c}}}} \right)}\end{array}$The NVP, Super P, and polyvinylidene difluoride (PVDF) were mixed in a weight ratio of 8:1:1 with a proper amount of 1‐methyl‐2‐pyrrolidinone (NMP) to form a uniform slurry, which was coated on Al foil, then dried overnight in a vacuum oven at 60 °C and cut into discs with a diameter of 12 mm. The areal mass loading for the normal and thick NVP cathodes were ≈2 and ≈15 mg cm−2, respectively. To achieve adequate infiltration of the electrolyte, the high‐loading NVP cathodes were fully immersed in the electrolyte for 48 h before battery assembly. The Na||NVP full cells were tested by pairing pre‐deposited Na/MgF2@RGO or Na/RGO anode (5 mAh) with the NVP cathode. The Na metal full‐cell performance with different N/P ratios of 18.9 and 2.5 was investigated, respectively. EIS was conducted with an amplitude of 10 mV over the frequency range of 100 kHz–5 mHz (BioLogic VSP electrochemical workstation).Computational DetailsAll DFT calculations were performed with the Vienna ab initio Simulation Package (VASP).[25] The projector‐augmented wave method was used to describe the electron–ion interactions.[26] The generalized gradient approximation with the Perdew, Burke, and Enzerhof functional was employed to evaluate the exchange‐correlation energy,[27] and cut‐off energy of 520 eV was used. For each configuration, there was a vacuum slab of 10 Å, and a 3 × 3 × 1 k‐point grid was set. The energy convergence tolerance was set to below 1 × 10−5 eV. The atomic coordinates were relaxed until the Hellmann–Feynman force on each atom was reduced to < 0.02 eV Å−1.AcknowledgementsL.Z. and Z.H. contributed equally to this work. This work was supported by the Australian Renewable Energy Agency (ARENA) project (G00849), the Australian Research Council (ARC) (DE170100928 and DP160102627) and the National Natural Science Foundation of China (Grant Nos. 51971124, 52171217, 22179079, U20A20249, and 21972108). The authors would like to thank Dr. Tania Silver for critical revisions of the manuscript, and eceshi for TGA and XPS characterizations.Open access publishing facilitated by University of Wollongong, as part of the Wiley ‐ University of Wollongong agreement via the Council of Australian University Librarians.Conflict of InterestThe authors declare no conflict of interest.Data Availability StatementThe data that support the findings of this study are available from the corresponding author upon reasonable request.Y. Zhao, K. R. Adair, X. L. Sun, Energy Environ. Sci. 2018, 11, 2673.F. Wu, J. Maier, Y. Yu, Chem. Soc. Rev. 2020, 49, 1569.a) B. Lee, E. Paek, D. Mitlin, S. W. Lee, Chem. Rev. 2019, 119, 5416;b) L. 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# In Situ Plating of Mg Sodiophilic Seeds and Evolving Sodium Fluoride Protective Layers for Superior Sodium Metal Anodes

, Volume 12 (32) – Aug 1, 2022
10 pages

/lp/wiley/in-situ-plating-of-mg-sodiophilic-seeds-and-evolving-sodium-fluoride-l5aASQJ7Dz
Publisher
Wiley
ISSN
1614-6832
eISSN
1614-6840
DOI
10.1002/aenm.202200990
Publisher site
See Article on Publisher Site

### Abstract

IntroductionGrid‐scale energy storage devices are essential for renewable energy utilization to guarantee continuous energy harvesting and steady energy output.[1] Rechargeable Li‐ion batteries (LIBs) have been well developed for portable electronic devices and electric vehicles, but their ever‐increasing material price limits their applications for grid‐scale energy storage.[2] Sodium‐based batteries have been regarded as affordable alternatives to LIBs since Na exhibits similar electrochemical properties but much higher resource abundance as compared with Li.[3] With a high energy density of 1166 mAh g−1 and a low electrochemical potential of −2.714 V, sodium metal is the premium choice of anode material for various sodium‐based batteries with high energy density including sodium metal batteries, sodium‐sulfur batteries, and sodium air batteries.[4] Na metal anode, however, suffers from a few challenging problems which impede their practical applications. For example, the plating/stripping of Na metal with large volume changes would lead to the breaking up of solid‐electrolyte interphase (SEI) layers and exposure of fresh Na, which would generate new SEI layers with continuous consumption of Na metal as well as the electrolyte, together with increasing of the cell impedance due to the gradually thickening SEI.[5] The repeated exposure of fresh Na is also the cause of low Coulombic efficiency (CE), which is another obstacle to the practical application of Na metal anodes. Besides, the uneven plating/stripping of Na metal would lead to the formation of Na metal dendrites, which would eventually pierce the separator, causing serious safety hazards due to internal short circuits with thermal runaways, and even fires or explosions.[6] The Na metal dendrites are likely to be detached to produce “dead Na,” which is inactive for further electrochemical reactions and leads to loss of active Na.[7]Great efforts have been made to develop stable Na metal anodes.[8] As a result of the huge volume changes during cycling, tiny fractures on the interfaces would gradually propagate into large cracks, which could eventually destroy the interface and leave the Na metal without protection.[9] It was reported that constructing stiff interfaces on Na metal anodes that could survive the huge volume changes during cycling is a valid means to enhance the stability of Na metal anodes, and such stable interfaces can be achieved either by surface coating or through electrolyte engineering.[10] It was also reported that the metals exhibiting some solubility in Na, such as Be, Mg, and Sn could serve as sodiophilic nucleation seeds for Na metal to reduce the nucleation overpotential and realize uniform plating.[11] Li et al. reported Sn nanoparticles embedded in a carbon framework as preferable nucleation sites, which enabled stable cycling of Na metal anodes with high Coulombic efficiency. While the metal seeded plating had little effect on the properties of the SEI layers, the interface degradation could not be inhibited. Another feasible approach toward stable Na metal anodes is applying porous host materials, which could alleviate the huge volume changes derived from the hostless nature of Na.[12] The plating and stripping of Na metal within the porous hosts are significantly different from the hostless process. The porous hosts could sustain some voids to accommodate Na metal, which does not shrink after the thorough stripping of Na metal.[13] As a result, the SEI films can persist on the surfaces of the hosts. The conductive porous hosts also offer large surface areas as nucleation sites, which could dramatically reduce the local current density and thus realize more stable Na metal plating.[14]A variety of porous materials have been developed as hosts for Na metal anodes, among which, porous metals (Cu, Al, Zn, and Ni) with good electrical conductivity have been proven to be promising candidates.[15] For instance, porous Al was reported as a host for Na metal anodes, which could generate smooth interfaces, as well as stable cycling.[16] The metal hosts are generally costly and exhibit high mass density, which would weaken the advantages of Na metal anodes. In comparison, various carbon‐based materials with good electrical conductivity, rich porous structures, and low mass density are preferable as host materials for Na metal anodes.[17] Porous carbons, carbon nanotubes (CNTs), and graphene are all well‐investigated host materials for Na metal anodes.[10c,13a,15b,18] It was demonstrated that Na metal plating on CNT electrodes could generate a smooth surface, which enabled stable cycling of the Na metal anodes.[18a,19] Chen and co‐workers have reported a porous reduced graphene oxide (RGO) as a host for Na metal anodes and successfully assembled Na‐CO2 batteries, which sustained stable cycling for more than 50 cycles.[20] Although approaches through interface engineering, metal seeded plating, and host material design all have been proved viable for stable Na metal anodes, reasonable material design with synergistic effects to incorporate these merits in one multifunctional host has rarely been achieved.Herein, we report a simple approach via hydrothermal assembly and thermal reduction to prepare RGO aerogel anchored with MgF2 nanocrystals (Figure 1a). The aerogel structure with its 3D conductive skeleton and rich porous structure reduces the local current density and offers interconnected void spaces to accommodate the deposition of Na metal. Moreover, the MgF2 component plays a key role in achieving non‐dendritic and highly reversible Na plating/stripping. The MgF2 nanocrystals can be converted into Mg (nucleation sites) and NaF (solid electrolyte interphase) during the initial stage of the first Na plating. The Mg sodiophilic seeds guarantee uniform Na nucleation and growth. Meanwhile, the NaF‐rich protective layer evolves over plating/stripping cycles, which suppresses the Na dendrite growth and prevents continuous electrolyte depletion. As a result of the synergistic effect, the MgF2@RGO aerogel serves as a multifunctional host to regulate the uniform deposition of Na metal, enabling the Na anode with highly stable cycling stability. Furthermore, the sodium metal full cells coupled with the Na metal anode confined in the MgF2@RGO aerogel host and Na3V2(PO4)3 (NVP) cathode deliver enhanced cycling stability and rate capability, suggesting its great potential for practical applications.1Figurea) Illustration of the synthesis of the MgF2@RGO aerogel; b) illustration of the working mechanism of MgF2@RGO aerogel as a host material for Na metal anodes.ResultsThe aerogels were synthesized by the hydrothermal method, followed by a freezing‐dry and a calcination process. The aerogels as‐synthesized were ≈10 mm in diameter and ≈5 mm in thickness, and the weight for each aerogel cake was ≈1.3 mg (Figure S1, Supporting Information). The lightweight aerogel host could accommodate sufficient Na metal without sacrificing the gravimetric energy density of the composite Na metal anode, as compared with the heavy porous metal‐based hosts.[16] The aerogels exhibit an interconnected porous structure (Figure 2a), which was most likely derived from the self‐assembly of the graphene layers during the hydrothermal treatment. The ethanediamine and sodium borate could interact with the functional groups on the graphene oxide (GO) layers during the hydrothermal treatment to prevent the layers from stacking with each other (Figure 1a).[21] The high‐resolution scanning electron microscopy (SEM) image of the aerogel (Figure 2b) demonstrates that nanocrystals with ≈20 nm in diameter were uniformly distributed on the graphene layers. The transmission electron microscope (TEM) image (Figure S2, Supporting Information) confirms the structure of the nanocrystals grown on thin graphene layers. The high‐resolution TEM image (Figure 2c) of the nanocrystals on the layers exhibits distinctive lattice fringes with a lattice spacing of 0.327 nm, which corresponds to the d‐spacing of (110) lattice planes for the MgF2 crystals. The energy‐dispersive X‐ray spectroscopy (EDS) mapping images (Figure 2d) indicate that the aerogel layers feature uniformly distributed C, Mg, and F elements, suggesting the uniform distribution of the MgF2 nanocrystals on the graphene layers. The X‐ray diffraction (XRD) pattern (Figure 2e) of the sample can be indexed to MgF2 (JCPDS No. 41–1443), and the broad peak at ≈28° is similar to that of the RGO sample, which should be derived from the graphene layers. Therefore, it is reasonable to deduce that the as‐synthesized aerogel samples consisted of a porous RGO matrix with anchored MgF2 nanocrystals (MgF2@RGO). The nitrogen adsorption isotherms (Figure 2f) indicate that the Brunauer–Emmett–Teller (BET) surface area of the MgF2@RGO aerogel (≈153 m2 g−1) is slightly smaller than that for the RGO aerogel (≈206 m2 g−1), which could have resulted from the reduced surface area due to the loading of MgF2 nanocrystals. The pore size distributions in Figure S3a, Supporting Information, also indicate that there were numerous nanopores of 5–10 nm in size, which are favorable for the penetration of electrolytes and transportation of the Na+ ions. The high surface area of the MgF2@RGO and RGO aerogels would decrease slightly (Figure S3b, Supporting Information) after compression due to battery assembly, while the amount of nanopores in the compressed aerogels would increase as a result of the enhanced contacting of the graphene layers (Figure S3c, Supporting Information). The X‐ray photoelectron spectroscopy (XPS) survey spectrum (Figure 2g) and high‐resolution C 1s, Mg 1s, and F 1s spectra (Figure S4, Supporting Information) confirm the presence of Mg, F, C, and O elements in the MgF2@RGO aerogel sample, as compared with the RGO sample, which is mainly composed of C and O elements. The content of the MgF2 on the aerogel was determined to be ≈22.5% by thermogravimetric analysis (TGA) (Figure 2h). The Raman spectra (Figure S5, Supporting Information) demonstrate that the intensity ratios (based on peak area) of the defect‐induced D band and crystalline graphite‐derived G band (ID/IG) for the MgF2@RGO and RGO aerogels were 1.48 and 1.37, respectively. The high ID/IG ratios indicate that the aerogels were highly defect‐rich, which should be attributed to the exfoliation of the graphene layers by water molecules via shear force from the ultrasonication treatment during the synthesis. The exfoliation of graphene to produce more graphene layers is favorable for the construction of the aerogel architecture and the loading of MgF2 nanocrystals. These characterizations indicate that the as‐synthesized MgF2@RGO aerogel exhibits thin and large‐area graphene layers, interconnected porous structures, and uniformly distributed MgF2 nanocrystals.2Figurea) SEM image, b) high‐resolution SEM image, c) high‐resolution TEM image, and d) EDS mapping images for the MgF2@RGO samples; e) XRD pattern, f) nitrogen adsorption isotherms, and g) XPS survey spectra for the RGO and MgF2@RGO samples; h) TGA curve for the MgF2@RGO sample.The MgF2@RGO and RGO aerogels were used as free‐standing electrodes to evaluate their viability as host materials for Na metal anode. The average Coulombic efficiencies (ACE) for the Na metal anodes with the bare Cu, RGO, and MgF2@RGO aerogel hosts (denoted as Na/Cu, Na/RGO, and Na/MgF2@RGO) at the current density of 0.5 mA cm−2 are presented in Figure 3a. The ACE in the first 20 cycles for the Na/MgF2@RGO electrode was 96.24%, which was higher than the ACE for the Na/RGO electrode (92.62%) or the Na/Cu electrode (84.73%), and also better than many of the state‐of‐the‐art reports.[9] The ACE for the Na metal anodes with the MgF2@RGO and RGO aerogel hosts (Figure S7a, Supporting Information) at the current density of 1 mA cm−2 was 94.08% and 91.30%, which were slightly lower than that at 0.5 mA cm−2. The decreased ACE indicates deteriorated cycling stability at the higher current density. As for the Na metal anode with the bare Cu, the ACE at 1 mA cm−2 was decreased to 82.32%, suggesting much inferior cycling stability of the Na metal anode. The CE for the Na/Cu electrode at 0.5 mA cm−2 (Figure 3b) was only around 80% in the first 100 cycles, but it seriously deteriorated afterward, with randomly scattered values ranging from 10% to 120%. The CE for the Na/RGO electrode at 0.5 mA cm−2 was initially lower than 60% and gradually increased to ≈92% after more than 40 cycles, which is likely due to the establishment of relatively stable SEI layers. The CE became unstable, however, with a clear trend of degradation after 100 cycles and even decreased to ≈20% after 300 cycles. In comparison, the CE for the Na/MgF2@RGO electrode quickly reached more than 96% within 20 cycles, and it remained steady for more than 300 cycles. The voltage curves (Figure S6, Supporting Information) at different cycles of the Na metal anodes with the MgF2@RGO, RGO, and Cu hosts during the CE testings are in accordance with the CE results in Figure 3b, demonstrating that the Na metal anodes with the MgF2@RGO aerogel host exhibited excellent cycling stability. When the current density was increased to 1 mA cm−2, the CE of the Na metal anodes with the MgF2@RGO aerogel host (Figure S7b, Supporting Information) remained stable for ≈100 cycles, which was dramatically superior to that for the RGO and bare Cu hosts. It can be inferred that the enhanced CE for the Na/MgF2@RGO electrode must have resulted from the uniform Na plating with more stable SEI layers, as a result of the introduced MgF2 nanocrystals. The CEs of the Na metal anodes at the current density of 0.5 mA cm−2 and areal capacity of 2 mAh cm−2 deliver similar results (Figure S8, Supporting Information), suggesting the superior cycling stability of the Na metal anodes with the MgF2@RGO aerogel host. The rate capabilities of the Na/Cu, Na/RGO, and Na/MgF2@RGO aerogels were also evaluated in symmetrical cells, as presented in Figure 3c. It can be seen that the Na/Cu electrode could achieve stable cycling at the current density of 0.1 to 2 mA cm−2, while the voltage hysteresis at 2 mA cm−2 was as large as 0.4 V, which is much higher than those for the Na/RGO and Na/MgF2@RGO electrodes at the current density of 5 mA cm−2. This indicates that the 3D interconnected conductive structures of the aerogels could effectively decrease the local current density during Na plating/stripping at high current rates, so as to deliver smaller voltage hysteresis as compared to the 2D planar bare Cu. The Na/RGO and Na/MgF2@RGO electrodes exhibited stable cycling with the current density ranging from 0.1 to 5 mA cm−2, and they also delivered stable cycling for ≈100 cycles when the current density was recovered from 5 to 0.5 mA cm−2. It is noticeable that the voltage hysteresis for the Na/MgF2@RGO electrode was slightly smaller than that for the Na/RGO electrode, which is similar to the results in previous reports, suggesting that the SEI layers on the Na/MgF2@RGO electrode may be thinner with enhanced interface stability.[7,19]3Figurea) Average Coulombic efficiencies at 0.5 mA cm−2, b) Coulombic efficiencies at 0.5 mA cm−2 and 0.5 mAh cm−2, and c) cycling performances at various areal current densities of the Na metal anodes with the Cu, RGO, and MgF2@RGO hosts; nucleation potentials of Na metal plating on the d) MgF2@RGO, e) RGO, and f) Cu hosts at various current densities; g) summary of the Na plating overpotentials on the various hosts; h) cycling stabilities of the Na metal anodes on the Cu, RGO, and MgF2@RGO hosts at 0.5 and 0.5 mAh cm−2.The nucleation overpotential for the Na metal plating, which is defined as the voltage gap between the dip and the plateau on the voltage curve of the first plating process, was also investigated for the host materials at various current densities. The nucleation overpotentials (Figure 3d) with the MgF2@RGO host at 0.1, 0.2, 0.5, 1, and 2 mA cm−2 were 27.7, 38.4, 45.2, 72.0, and 92.6 mV, respectively, which are smaller than those for the Na plating on the RGO host (Figure 3e). In comparison, the nucleation overpotentials for the Na plating on the bare Cu electrodes (Figure 3f) at 0.1, 0.2, 0.5, 1, and 2 mA cm−2 were 154.4, 163.8, 223.7, 292.9, and 307.3 mV, respectively. The large nucleation overpotentials for the planar Cu electrodes (Figure 3g) further demonstrate the advantage of the aerogels with 3D conductive porous structures, which could increase the affinity between Na and the aerogels and reduce the local current density. The lowest nucleation overpotentials for the MgF2@RGO host were most likely derived from the synergetic effect of the conductive graphene aerogel and the anchored MgF2 nanoparticles. The latter could be converted into Mg as nucleation seeds for Na metal plating to further decrease the overpotentials. Long‐term cycling of the Na metal anodes with the Cu, RGO, and MgF2@RGO hosts was performed at 0.5 mA cm−2 with an initially plated 5 mAh of Na as the working electrode (Figure 3h). The Na/Cu electrode exhibited a dramatic increase of the voltage hysteresis from ≈0.1 V to more than 1 V after stable cycling for ≈300 h, which may be derived from the complete consumption of the pre‐deposited Na on the working electrode after cycling because of the continuous formation of “dead” Na. The unstable cycling performance of the Na/Cu electrode is in accordance with its CE performance, as presented in Figure 3b, further demonstrating that the Na metal anodes on planar Cu exhibited inferior cycling stability. Although the cycling stability for the Na/RGO electrode was enhanced to ≈650 h, a similar exacerbation of the voltage hysteresis occurred thereafter, which suggests deteriorated interfaces of the Na metal anodes after cycling. In contrast, the Na/MgF2@RGO electrode delivered smaller voltage hysteresis and much more stable cycling performance for more than 1600 h, which is more than two times longer than that for the Na/RGO electrode. The plating/stripping curves of the Na metal anodes with the Cu, RGO, and MgF2@RGO aerogel hosts at different cycling times (Figure S9, Supporting Information) demonstrate that the voltage curves for the Na/MgF2@RGO are stable during cycling, as compared with the inferior stability for the Na/Cu and Na/RGO anodes with gradually increased voltage tips. Furthermore, the Na/MgF2@RGO anodes could deliver stable cycling for ≈500 h at 1 mA cm−2 (Figure S10a, Supporting Information) and ≈300 h at 2 mA cm−2 (Figure S10b, Supporting Information), which are both better than the Na metal anodes with the bare Cu and RGO hosts. The excellent cycling stability of the Na metal anodes with the MgF2@RGO hosts is superior to those in many of the state‐of‐the‐art reports (Figure S11, Supporting Information), which further demonstrates the advantages of this host material for Na metal anodes.The morphology evolution of the Na metal anodes with the bare Cu and aerogel hosts was also investigated. The morphologies for the pristine MgF2@RGO (Figure S12a, Supporting Information) and RGO (Figure S12b, Supporting Information) are similar, showing relatively flat surfaces derived from randomly assembled RGO layers. The bare Cu exhibits an even but rough surface (Figure S12c, Supporting Information). As presented in Figure 4a, the Na metal initially plated on the MgF2@RGO aerogel exhibits a nodule‐like structure, which is similar to that for the RGO aerogel in Figure 4e. However, the Na nodules on the MgF2@RGO aerogel are in intimate contact and form large lumps with reduced surface area, as compared with the much smaller and more separated nodules which present a porous structure (Figure 4e). Similarly, many nodules can be seen in the case of the Na metal plating on the bare Cu (Figure 4i), while there are also some Na whiskers and the surfaces of the nodules were cracked, indicating inferior interface stability. After 200 cycles, the Na metal plating on the MgF2@RGO aerogel (Figure 4b) exhibits a dense and flat surface, while the surface for the Na/RGO (Figure 4f) is rough and uneven. The Na metal on the bare Cu after 200 cycles (Figure 4j) shows much smaller Na nodules and even mossy Na, which is most likely derived from the cracking of large Na nodules, suggesting the poor cycling stability of the Na metal on bare Cu. When the plated Na metal was stripped from the MgF2@RGO aerogel after the 200th cycle, a flat and uniform interface could be achieved (Figure 4c). From this, it can be deduced that the MgF2@RGO aerogel could significantly enhance the cycling stability of the Na metal anodes, as a result of the 3D porous conductive matrix and MgF2‐derived stable interfaces (Figure 4d). As for the surface of the stripped Na/RGO electrode (Figure 4g), irregular‐shaped island‐like Na lumps were retained, which are most likely detached inactive “dead Na” as a result of the poor stability of the SEI layers (Figure 4h). As for the bare Cu (Figure 4k), mossy‐like Na was retained with noticeable cracks on the electrode after the 200th stripping. Its inability to be stripped from the bare Cu indicates that the mossy Na is inactive, implying that it is detached from the current collector as “dead Na.” The continuous accumulation of dead Na on the bare Cu (Figure 4l) indicates that the Na metal on bare Cu exhibits poor cycling stability. The cross‐section SEM images and the digital photographs of the cycled Na metal anodes with the bare Cu, RGO, and MgF2@RGO aerogel hosts are presented in Figures S13 and S14, Supporting Information. It can be seen that the Na metal is preferably deposited in/onto the MgF2@RGO aerogel matrix (Figure S14a, Supporting Information) with dense and compact morphology and little volume change upon cycling (Figure S13a,b, Supporting Information), indicating good sodiophilicity of the aerogel host and superior stability of the composite Na metal anode. In contrast, the Na/RGO (Figure S14b, Supporting Information) presents a little amount of Na metal plated onto the battery shell instead of the host, and the volume change (Figure S13c,d, Supporting Information) upon cycling is larger. As for the Na/Cu electrode (Figure S14c, Supporting Information), a considerable amount of Na metal is deposited on the battery shell, and the thickness of the Na metal on the bare Cu was decreased due to more Na being deposited onto the battery shell upon cycling. Therefore, the morphology evolution clearly demonstrates that the RGO aerogel with the 3D conductive matrix can enhance the stability of Na metal to some extent as compared to the planar Cu. Significantly, the MgF2@RGO aerogel host could achieve dramatically improved Na metal plating, as a result of the synergy of MgF2 nanocrystals and RGO aerogel, which could stabilize the interfaces.4FigureSEM images of the Na metal anodes with MgF2@RGO host after a) the 1st plating, b) the 200th plating, c) the 200th stripping, and d) corresponding illustration of the uniform Na plating on the MgF2@RGO host. SEM images of the Na metal anodes with the RGO host after e) the 1st plating, f) the 200th plating, g) the 200th stripping, and h) corresponding illustration of the Na plating on the RGO host with dead Na growth. SEM images of the Na metal anodes on bare Cu after i) the 1st plating, j) the 200th plating, k) the 200th stripping, and l) corresponding illustration of the Na plating on the bare Cu with mossy dead Na growth. Scale bars: 20 µm.To achieve an in‐depth understanding of the working mechanism of the MgF2@RGO aerogel as host material for Na metal anodes, experimental and theoretical investigations were further conducted. During the first Na plating into the MgF2@RGO aerogel host (Figure S14, Supporting Information), two voltage slopes can be seen on the voltage curve. It is noticed that the electrode delivers a high initial capacity of ≈1000 mAh g−1, which should be ascribed to the Na+ storage into the RGO matrix, the conversion of MgF2 nanocrystals, and also the formation of SEI layers. Phase evolution of the MgF2@RGO aerogel host at various stages of Na metal plating (Figure 5a) was studied via ex situ XRD analysis, as presented in Figure 5b. Initially, at stage a with no Na plating, the host presents a characteristic broad peak that corresponds to the diffraction of RGO and also the diffraction peaks for MgF2. At stage b when the voltage has decreased to 0 V with intercalation of Na+, the diffraction peaks for MgF2 have vanished, and new peaks for Mg and NaF have appeared. The phase change from stages a to b is in line with the voltage features in Figure S15, Supporting Information, suggesting that the MgF2 nanocrystals were converted into Mg and NaF during the Na plating. The voltage dip between stages b and c corresponds to the nucleation of Na on the MgF2@RGO aerogel host, after which, the diffraction peaks for metallic Na gradually emerge and become stronger from stages c to e with continued Na metal plating. When the plated Na metal was stripped from the aerogel host at stage f, the diffractions for metallic Na disappeared, which indicated that most of the Na had been completely stripped, with no remaining “dead” metallic Na. Diffraction peaks for Mg and NaF can also be detected, suggesting that the Mg and NaF can be preserved during the Na plating/stripping process.5Figurea) Various stages on the potential curve for the first Na metal plating/stripping on the MgF2@RGO host, and b) corresponding evolution of the XRD patterns during the plating/stripping process; c) illustration of the binding sites for Na on the Mg/graphene surfaces, and d) corresponding binding energies based on the DFT calculations.The surfaces of the electrodes were also investigated via XPS analysis. As presented in Figure S16, Supporting Information, the Na 2s peaks at 1070 and 1070.4 eV for the cycled Na/MgF2@RGO electrode (Figure S16a, Supporting Information) can be assigned to metallic Na and NaF (and/or Na‐O), respectively.[22] The fitted F 1s peak at 683.3 eV (Figure S16b, Supporting Information) is corresponding to NaF. Similar peaks assigned to NaF can also be observed on the cycled Na/RGO electrode (Figure S16c,d, Supporting Information), which should be originated from the fluoroethylene carbonate (FEC) derived SEI layers. It is noteworthy that the cycled Na/MgF2@RGO electrode shows a much stronger NaF signal than that of the Na/RGO electrode, which implies additional NaF has been generated from the conversion of MgF2 in the cycled Na/MgF2@RGO electrodes. To explore the feasibility of Na plating on the surfaces of Mg and RGO, density functional theory (DFT) calculations were performed to determine the binding energy of Na on the (101) plane for Mg and the graphene plane for RGO. Based on the calculations, the preferred binding sites are illustrated in Figure 5c, and the corresponding binding energies are presented in Figure 5d. The binding energies for Na adsorption on Mg are −0.20, −0.19, and −0.14 eV (Figure S17, Supporting Information), suggesting that the surface of Mg is favorable for Na adsorption, while the binding energies for the graphene layer are 0.38, 0.61, and 0.72 eV, indicating that the process of Na adsorption on the graphene layers requires additional energy, which is unfavorable.Based on these investigations, it is rational to infer that the MgF2 nanocrystals on the MgF2@RGO aerogels can be converted into Mg and NaF during the Na plating process (Figure 1b), and the Mg nanocrystals could alloy with metal Na to serve as nucleation sites and achieve stable Na plating. The feasibility of Mg for nucleation sites has also been proved by recent research, and NaF has also been well recognized as a preferred component of SEI layers to construct NaF‐rich stable interfaces.[11a,12,23] Besides, 3D RGO aerogels significantly reduce the local current density and accommodate the deposited Na. Therefore, the working mechanism of MgF2@RGO aerogel as a host material for Na metal anodes can be summarized as the synergetic effects of the 3D conductive matrix, Mg‐induced sodiophilic plating, and NaF‐generated stable SEI layers. The outstanding performance of the Na metal anodes in the MgF2@RGO aerogel host with enhanced CE, improved interface stability, and ultra‐stable long‐term cycling can be attributed to the synergetic effects derived from the multifunctional aerogel host.To further verify the feasibility of the aerogel hosts for practical Na metal anodes, Na metal full cells with NVP cathode were investigated. It can be seen that, at the low current rate of 0.2 C, the full cells with the Na/RGO and Na/MgF2@RGO anodes delivered a comparable discharge capacity of 110 mAh g−1 (Figure 6a), while at higher rates of 0.5, 1, 2, and 5 C, the full cells with Na/MgF2@RGO anode delivered enhanced rate capability as compared with that for the Na/RGO anode. When the current density was recovered to 1 C, the full cells with Na/MgF2@RGO anode delivered stable cycling for more than 100 cycles with a decent discharge capacity of ≈96 mAh g−1. The full cells with Na/RGO anode exhibited a slightly lower discharge capacity of ≈90 mAh g−1, and the cycling gradually deteriorated after 100 cycles. It is reasonable to deduce that the inferior rate performance and cycling performance should be attributed to the instability of the Na/RGO anode. Charge/discharge curves of the full cells (Figure 6b) suggest that the voltage hysteresis gradually increased at higher rates. The full cells delivered comparable voltage plateaus at various current rates, while the voltage hysteresis of the full cells with the Na/RGO anode occurred at smaller specific capacities than those with the Na/MgF2@RGO anode. Figure 6c shows that the full cell with the Na/RGO anode delivered stable cycling at 1 C for 200 cycles with a high capacity retention of 91.3% (79.88 out of 87.48 mAh g−1). The cycling of the full cell with the Na/RGO anode dramatically deteriorated after ≈90 cycles, with the discharge capacity quickly decreasing from 79.5 to 38.8 mAh g−1. The poor cycling stability of the full cells should be attributed to the worse stability of the Na metal accommodated in the RGO host, which may be exhausted as a result of inferior interface stability. In Figure 6d, electrochemical impedance spectroscopy (EIS) analysis of the cycled full cells indicates that the impedance spectra with the Na/MgF2@RGO anode were stable during cycling, suggesting good electrode stability with constant interfaces. As compared to the full cells with the Na/RGO anodes, the semicircles corresponding to the charge transfer resistance gradually increased in size, implying progressive thickening of the SEI layers on the anodes. Furthermore, second small semicircles were identified on the EIS curves of the full cells with the Na/RGO anodes after 20 and 100 cycles, indicating the build‐up of new electrode interfaces, which could be resulted from the accumulation of “dead Na” on the anode surfaces. The potential of the aerogel host materials for practical Na metal full cells is further investigated with high‐loading NVP cathodes. As demonstrated in Figure 6e, even at the critical condition with a high cathode loading of ≈15 mg cm−2 and a low N/P ratio of ≈2.5, the Na metal full cells with the MgF2@RGO aerogel host could deliver stable cycling for more than 30 cycles, which is significantly superior to that for the RGO host with only 15 cycles of stable cycling. The charge/discharge curves (Figure 6f) for the Na metal full cells with the high loading NVP cathode and Na/MgF2@RGO anode at the 1st, 5th, 10th, and 20th cycles are nearly overlapped, suggesting good cycling stability under practical conditions. Thereby, it can be confirmed that the stability of the Na metal anode was critical for the stable cycling of the Na metal full cells, and the MgF2@RGO aerogel host could significantly increase the stability of the Na metal anode to achieve dramatically enhanced full cell performance.6Figurea) Rate performance, b) charge/discharge curves, c) cycling performance, and d) EIS evolution of the NVP‐based sodium metal full cells with the Na/MgF2@RGO and Na/RGO anodes, with the equivalent circuit shown in the inset. e) Cycling performance and f) charge/discharge curves at 1st, 5th, 10th, 20th, and 30th cycles for the high loading NVP cathode‐based Na metal full cells with the Na/MgF2@RGO and Na/RGO anodes.ConclusionA robust MgF2@RGO aerogel has been developed as a multifunctional host material for Na metal anode. Via in situ electrochemical conversion of MgF2, Mg sodiophilic sites, and NaF protective layers can be simultaneously planted on the RGO matrix, which can synergistically regulate the Na plating and achieve a superior long‐cycling lifetime. The morphology evolutions confirm that the combination of Na metal anode with the MgF2@RGO aerogel suppresses the formation of mossy Na and dendrite growth. The sodium metal full cells delivered enhanced rate and cycling performances, verifying the feasibility of the sodium metal anode for practical applications. The concept of host material design with synergistic effects of a porous conductive matrix, metal‐derived sodiophilic nucleation, and inorganic‐rich stable interfaces has been proved to be effective to achieve stable Na metal anodes, which sheds light on the future development of high‐energy sodium metal batteries.Experimental SectionPreparation of the MgF2@RGO AerogelsAll the reagents were of analytical grade and used as received without further purifications. GO solution was prepared by the modified Hummer's method, as reported previously.[24] Typically, 15 mL of GO solution (concentration ≈5 mg mL−1) was added into a 100 mL beaker with 25 mL of deionized water (DI), and the solution was stirred for 15 min before being treated under ultrasonication for 30 min. Thereafter, 2 mL of Mg(NO3)2 aqueous solution (160 mg mL−1) and 4 mL of NaF aqueous solution (40 mg mL−1) were dropped into the GO solution in sequence, and the mixture was kept under stirring for 1 h. Then, 160 µL of ethanediamine and 400 µL of sodium borate (2 wt.%) were dropped into the above solution and stirred for 15 min. Afterward, 1 mL of the mixture was placed into a 5 mL vial and sealed with a silica gel stopper, and the vials and 5 mL of DI water were then put into a 100 mL Teflon‐lined stainless‐steel autoclave and maintained at 180 °C for 18 h before being cooled to room temperature naturally. The as‐synthesized hydrogel was carefully transferred into a beaker with 490 mL DI water and 10 mL ethanol for distillation, and the distillation was maintained for 12 h and repeated four times. The preparation of GO hydrogel was conducted with the identical procedure, except for the addition of Mg(NO3)2 and NaF aqueous solution. The MgF2@GO and GO hydrogels were frozen in a refrigerator and then freeze‐dried. Thermal treatment was conducted at 500 °C for 2 h under an Ar atmosphere with a temperature ramp of 5 °C min−1 to obtain the MgF2@RGO and RGO aerogels. The aerogels as‐synthesized were ≈10 mm in diameter and ≈5 mm in thickness. The thickness of the aerogels after compression for battery assembly would be decreased to ≈0.5 mm, while the diameter of the aerogels would be retained. The weights for each MgF2@RGO and RGO aerogel were ≈1.3 and ≈1 mg, respectively.Materials CharacterizationsSEM images and EDS maps of the samples were collected with a field‐emission scanning electron microscope (FE‐SEM, Hitachi Regulus‐8230, Japan) equipped with an Ar‐filled sample transfer box. TEM images and high‐resolution TEM images were taken with a TEM Tecnai F30 (FEI) with an accelerating voltage of 300 kV. XRD patterns were recorded on an X‐ray diffractometer (Rigaku SmartLab) with Cu Kα radiation (λ = 0.15418 nm) at a voltage of 45 kV and a current of 200 mA. Raman spectra were recorded on an ARCSP 2558 spectrometer (Princeton Instruments). TGA was performed on a Setaram instrument (SETSYS EVO 18) from 30 to 700 °C under an air atmosphere with a heating rate of 5 °C min−1. The BET surface area was measured using N2 adsorption/desorption via an automated surface area analyzer (ASAP2420‐4MP). XPS spectra were recorded (Thermo Scientific ESCALAB 250Xi) with an Al Kα radiation source. For the XPS analysis of the cycled electrodes, the batteries were cycled at 0.5 mA cm−2 for ten cycles to establish stable SEI layers. Sample preparations for the characterizations of the cycled electrodes (SEM, XRD, and XPS) were carried out in an Ar‐filled glovebox, where the electrodes were obtained from the disassembled batteries, washed with the electrolyte solvents (EC/DEC, v/v = 1:1, with 5% FEC), dried in the chamber of the glovebox, placed in air‐tight chambers (SEM/XPS) with Ar atmosphere or protected with Kapton tape (XRD), and then transferred into the facility for characterizations. Sample preparations for the BET analysis of the compressed aerogel hosts were similar, except for using the electrolyte solvents instead of the electrolyte for the battery assembly.Electrochemical MeasurementsCoin‐type cells were tested at 25 °C on a Neware battery test system (CT‐4008T‐5V10mA‐164, Shenzhen China). The aerogels were directly used as free‐standing electrodes without further treatment. The coin‐type cells (LIR2025) were assembled in an argon‐filled glovebox (moisture and oxygen < 0.1 ppm) with Whatman glass fiber as the separator, Na foils as both counter and reference electrodes, and 100 µL of electrolyte composed of 1 m NaClO4 in ethylene carbonate (EC) and diethyl carbonate (DEC) (1:1 by volume), with 5% FEC as additive. The symmetrical Na cells were tested within a protective voltage range of −4 to 4 V. For the cross‐section morphology analysis of the cycled Na metal anodes with MgF2@RGO or RGO aerogel hosts, the batteries were assembled with an additional layer of Cu foil to support the samples. To evaluate the Coulombic efficiency, the batteries were first pre‐cycled at 0–1 V for five cycles (Figure S18, Supporting Information) to eliminate contaminations and establish stable SEI layers,[18c] then plated for 1 h and stripped to 0.5 V (vs Na+/Na) at the current density of 0.5 or 1 mA cm−2. For symmetrical Na||Na cell tests, 5 mAh cm−2 of Na was pre‐deposited onto the working electrodes (Cu, RGO, and MgF2@RGO), and the diameters of the Na counter/reference electrode and the Cu current collector were 16 and 12 mm, respectively. Both charging and discharging were conducted at the current density of 0.5, 1, or 2 mA cm−2 and cut off at a fixed time of 60 min. The ACEs of the Na metal anodes were tested following protocols similar to those in previous reports.[9] Briefly, the current collectors were first activated at the current density of 0.5 mA cm−2 by plating 5 mAh of Na and then stripping it to 1.2 V. Thereafter, 5 mAh (Qp) of Na was plated on the activated current collectors for working electrodes at the current density of 0.5 mA cm−2, then the working electrodes were cycled at the current density of 0.5 mA cm−2 and capacity of 0.5 mAh cm−2 (Qc) for n cycles (n = 20), and finally the residue Na was stripped to 1.2 V (Qs). The ACE was obtained according to the following equation:1ACE=(Qs+nQc)/(Qp+nQc)$\begin{array}{*{20}{c}}{{\rm{ACE}} = \left( {{{\rm{Q}}_{\rm{s}}} + {\rm{n}}{{\rm{Q}}_{\rm{c}}}} \right)/\left( {{{\rm{Q}}_{\rm{p}}} + {\rm{n}}{{\rm{Q}}_{\rm{c}}}} \right)}\end{array}$The NVP, Super P, and polyvinylidene difluoride (PVDF) were mixed in a weight ratio of 8:1:1 with a proper amount of 1‐methyl‐2‐pyrrolidinone (NMP) to form a uniform slurry, which was coated on Al foil, then dried overnight in a vacuum oven at 60 °C and cut into discs with a diameter of 12 mm. The areal mass loading for the normal and thick NVP cathodes were ≈2 and ≈15 mg cm−2, respectively. To achieve adequate infiltration of the electrolyte, the high‐loading NVP cathodes were fully immersed in the electrolyte for 48 h before battery assembly. The Na||NVP full cells were tested by pairing pre‐deposited Na/MgF2@RGO or Na/RGO anode (5 mAh) with the NVP cathode. The Na metal full‐cell performance with different N/P ratios of 18.9 and 2.5 was investigated, respectively. EIS was conducted with an amplitude of 10 mV over the frequency range of 100 kHz–5 mHz (BioLogic VSP electrochemical workstation).Computational DetailsAll DFT calculations were performed with the Vienna ab initio Simulation Package (VASP).[25] The projector‐augmented wave method was used to describe the electron–ion interactions.[26] The generalized gradient approximation with the Perdew, Burke, and Enzerhof functional was employed to evaluate the exchange‐correlation energy,[27] and cut‐off energy of 520 eV was used. For each configuration, there was a vacuum slab of 10 Å, and a 3 × 3 × 1 k‐point grid was set. The energy convergence tolerance was set to below 1 × 10−5 eV. The atomic coordinates were relaxed until the Hellmann–Feynman force on each atom was reduced to < 0.02 eV Å−1.AcknowledgementsL.Z. and Z.H. contributed equally to this work. This work was supported by the Australian Renewable Energy Agency (ARENA) project (G00849), the Australian Research Council (ARC) (DE170100928 and DP160102627) and the National Natural Science Foundation of China (Grant Nos. 51971124, 52171217, 22179079, U20A20249, and 21972108). The authors would like to thank Dr. Tania Silver for critical revisions of the manuscript, and eceshi for TGA and XPS characterizations.Open access publishing facilitated by University of Wollongong, as part of the Wiley ‐ University of Wollongong agreement via the Council of Australian University Librarians.Conflict of InterestThe authors declare no conflict of interest.Data Availability StatementThe data that support the findings of this study are available from the corresponding author upon reasonable request.Y. Zhao, K. R. Adair, X. L. Sun, Energy Environ. Sci. 2018, 11, 2673.F. Wu, J. Maier, Y. Yu, Chem. Soc. Rev. 2020, 49, 1569.a) B. Lee, E. Paek, D. Mitlin, S. W. Lee, Chem. Rev. 2019, 119, 5416;b) L. 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