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Reduced graphene oxide aerogel as stable host for dendrite-free sodium metal anode
Energy Storage Materials 22 (2019) 376–383
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Reduced graphene oxide aerogel as stable host for dendrite-free sodium metal anode Feng Wu a, b, Jiahui Zhou a, Rui Luo a, Yongxin Huang a, Yang Mei a, Man Xie a, **, Renjie Chen a, b, * a b
School of Materials Science & Engineering, Beijing Key Laboratory of Environmental Science and Engineering, Beijing Institute of Technology, Beijing 100081, China Collaborative Innovation Center of Electric Vehicles in Beijing, Beijing 100081, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Sodium metal anode Graphene aerogel Oriented freeze-drying Long-term stability High current density
Sodium (Na) metal has attracted great attention as a promising anode for next-generation energy storage systems because of its abundant resources, potentially low cost, and high theoretical capacity. However, severe dendrite growth, large volume expansion during plating/stripping lead to poor cycle performance and limit the practical application of sodium metal anodes. Here, oriented freeze-drying and molten infusion are used to synthesize a Nainfused reduced graphene oxide aerogel (Na@rGa) composite anode using a reduced graphene oxide aerogel (rGa) as a stable host. This unique host shows ultra-light weight which guarantees high capacity (1064 mAh g1), and the large specific surface area and uniform pores effectively lower the local current density and uniform electrolyte distribution from the inner to the outside of electrode. The Na@rGa anode exhibits low overpotential (50 mV) and stable cycle performance for 1000 cycles at 5 mA cm2 in carbonate electrolyte system. The electrochemical performance of a full cell using the Na@rGa composite anode is clearly superior to that of a reference cell. This work provides a new horizon for construction of three-dimensional stable hosts and safe sodium metal anodes.
1. Introduction With the increasing demand for large-scale energy storage, sodiumion batteries (SIBs) have attracted tremendous attention as alternatives to lithium-ion batteries (LIBs) because of their abundant resources, low cost, and analogous battery chemistry to that of LIBs [1,2]. Progress has been achieved in terms of cathode materials for SIBs including layered transition metal oxides, Prussian blue, organic polymers, and polyanions [3–6]. However, research on anode materials for SIB still faces great challenges, mainly because the commercial graphite used in LIBs cannot be applied to SIBs [7,8]. Considerable effort has been devoted to investigating high-capacity materials for SIBs, including carbon materials [9, 10], phosphorus [11,12], and transition metal oxides and alloys [13,14]. Compared to these anode candidates, metallic Na possesses high theoretical capacity (1165 mAh g1) and low potential (2.714 V vs. the standard hydrogen electrode) [15]. Furthermore, the investigation of Na metal anodes is of great importance for batteries with high energy density such as Na–O2 [16,17] and Na–S batteries [18].
Despite all the advantages mentioned above, secondary batteries that use Na metal as anodes still face severe problems, especially in carbonate electrolyte [19]. Uneven nucleation, dendrite growth of Na during the plating process and uncontrollable side reactions between fresh Na with the liquid electrolyte and formation of “dead Na” in the following stripping process, leading to short cycle life and poor electrochemical performance [20–22]. Moreover, non-skeleton Na metal shows a huge volume change during plating and stripping, leading to fracture of solid electrolyte interface (SEI) and severe side reactions. The development of sodium metal batteries (SMBs) is still in its initial stage. At present, dendrite growth and volume expansion of the Na metal anode during cycling are the principal barriers limiting the further development of the SMBs [23]. Several strategies used to realize lithium (Li) metal anodes with long cycle lives have been adapted to solving the above-mentioned problems of Na metal anodes, mainly including fabricating organic and inorganic thin layer [24], adding liquid electrolyte additives [25] and using highly concentrated Sodium bis (fluorosulfonyl) imide in 1,2-dimeth-oxyethane
* Corresponding author. School of Materials Science & Engineering, Beijing Key Laboratory of Environmental Science and Engineering, Beijing Institute of Technology, Beijing 100081, China. ** Corresponding author. School of Materials Science ＆ Engineering, Beijing Key Laboratory of Environmental Science and Engineering, Beijing Institute of Technology, Beijing 100081, China. E-mail addresses: [email protected] (M. Xie), [email protected] (R. Chen). https://doi.org/10.1016/j.ensm.2019.02.015 Received 25 December 2018; Received in revised form 18 February 2019; Accepted 18 February 2019 Available online 20 February 2019 2405-8297/© 2019 Elsevier B.V. All rights reserved.
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deionized (DI) water by probe sonication for 60 min. Ascorbic acid (50 mg) was added to the solution, which was then continuously stirred for 2 h at room temperature. The solution was then transferred into a cylindrical bottle, sealed and heated in a water bath at 60 C for 4 h. After solvothermal reaction, the resulting cylindrical-shaped hydrogel was collected, washed with a mixture of ethanol and DI water (1:10 v/v) several times to remove unreacted chemicals, and then immersed in ethanol/DI water (1:10 v/v) for 12 h. The collected hydrogel was hanged in refrigerator about 6 h to orient freeze and obtained by freeze-drying overnight in order to remove absorbed water and ethanol. The obtain aerogel was cut into thin slices and placed in a glove box. For comparison, an aerogel foam was also synthesized with the same method by extending the reaction time in the water bath from 2 to 24 h under the same conditions. Na infusion into rGa: Na infusion process was carried out in a glove box. Solid Na cube was melted on a hot plate, and as-prepared rGa slices were then immersed into molten Na. The infusion process was finished in 20 s and 3D Na@rGa composite anode was obtained. Synthesis of P2–Na0.67Ni0.25Mn0.75O2 (P2-NNM) electrode: The P2NNM cathode material was synthesized by the solid reaction described in our previous report [3] and summarized as follows. An appropriate ratio of precursor powders (Na2CO3, NiO, and Mn2O3) was thoroughly mixed in a high-energy ball mill for 6 h. The mixture was then calcined twice at 900 C for 16 h in air atmosphere with grinding between the calcination steps. The calcined powder was then ground and stored in an Ar-filled glovebox until use. P2-NNM electrodes were prepared by casting a slurry of active material, acetylene black, and PVDF binder (75:15:10 w/w/w) on an aluminum foil collector.
electrolyte [26]. These methods effectively inhibit the growth of Na dendrites by forming artificial SEI films on the Na surface. However, in most cases, surface protection alone cannot inhibit the large volume expansion of the Na metal anode during electrochemical cycling, especially at high current density (>3 mA cm2). Meanwhile, the artificial SEI film inevitably increases interface resistance and restricts the migration of Na ions. Several attempts to develop various matrixes as hosts for Na/Li metal with the aim of alleviating both volume changes and dendrite growth have been explored [27–31]. Fan's group developed a three-dimensional (3D) carbon felt as a host for Na metal, which displayed relatively uniform and stable Na plating/stripping behavior [27]. Hu's group fabricated a carbonized wood composite as a high-surface-area, conductive, stable Na metal anode that tackled nearly all the problems associated with SMBs [28]. Chen and co-workers reported an Na/graphene film composite anode that exhibited great stability in Na–CO2 batteries [38]. These results suggest that confining Na metal in a 3D host could greatly enhance the stability of Na metal anodes. However, most available matrixes have poor wettability towards molten Na metal. Researchers have tried to decrease the contact angle of matrixes and solve the wetting issue between molten alkali metal and matrixes by coating with a metal oxide (ZnO [32], Al2O3 [33], and SnO2 [34]) and adding alloying elements (Ge [35], Al [36], and Mg [37]). However, these approaches inevitably caused a loss of capacity, in addition, the adsorption of Na in the host reduces the contact area between the Na and the electrolyte. The electrolyte does not sufficiently wet the entire electrode which reduces the cycle stability and enhances the possibility of dendrite formation after long cycling. Developing a “sodiophilic” host structure with high-loading and good wettability with electrolyte is highly desirable to achieve high-performance SMBs. An ideal host for Na should possess flexibility, ultra-light weight to obtain a composite electrode with high capacity, large specific surface area to lower the local current density, good chemical stability to avoid side reactions with liquid electrolyte, good Na and electrolyte wettability. rGa is a typical 3D porous framework that has attracted extensive attention because of its interconnected hierarchical porous system, large specific surface area, excellent adsorption in organic solvents, and ultralight weight [39–41]. Furthermore, the porous framework of rGa aids fast ion transfer inside the aerogel and enables capillary action during the melt infusion process [32,42]. These unique features make rGa promising as a carbon host for encapsulating Na metal. Notably, the self-assembly process of rGa is accompanied by the reduction of oxygen functional groups, which is of important for Na metal infusion [43]. Herein, we used a modified hydrothermal reduction reaction and subsequent oriented freeze-drying process to form rGa host with ultralight weight and good wettability. Abundant oxygen functional groups of the rGa aid infusion of molten Na to form Na@rGa composite, and oriented freeze-drying ensures that rGa has a highly homogeneous pore structure. Benefiting from these unique structure and chemical property, the obtained Na@rGa composite anode shows an ultra-high loading of Na metal of 98.75 wt% and capacity of 1064 mAh g1. Moreover, the homogeneous pores enhance wettability between Na@rGa and electrolyte, which is crucial for high ionic mobility and uniform electrolyte distribution from the inner to the outside of electrode, the large specific surface area reduce local current density and inhibit the nucleation of Na dendrite. The Na@rGa anode exhibits a low overpotential of about 50 mV at 5 mA cm2 in a traditional carbonate electrolyte without any additives over 1000 cycles. The electrochemical performance of a full battery with the Na@rGa composite anode is also investigated.
2.2. Structural characterization and electrochemical measurements Structural Characterization: The lattice structure of Na@rGa was recorded using a D8 Advance diffractometer (XRD, Bruker AXS, WI, USA) with a scanning range from 10 to 70 . The morphologies images of the sample were taken using a Hitachi SU-70 field-emission scanning electron microscope (SEM) with an energy-dispersive X-ray (EDX) spectroscopy attachment and accelerating voltage of 5.0 kV. To observe the surface morphology of the Na@rGa composite electrode and bare Na foil after cycling, the cycled cells were disassembled in a glove box and then washed with dimethyl carbonate several times to remove surface impurities. Fourier transform infrared (FTIR) spectra were recorded on a Nicolet iS50 FTIR spectrometer (Thermo Scientific). X-ray photoelectron spectra (XPS) measurements were recorded on a spectrometer with Al Kα radiation (hv ¼ 1486.6 eV). Electrochemical Measurement: To investigate the electrochemical plating/stripping of the Na@rGa composite electrode, 2032-type coin cells were assembled with two identical Na@rGa electrodes. For comparison, symmetric cells were fabricated with two bare Na metal foil electrodes. For full cell measurements, P2-NNM electrodes were paired with Na@rGa and bare Na. All the coin cells were assembled using 1.0 M NaClO4 in EC/DEC (1:1 v/v) solution as the electrolyte without any additives and glass fiber as the separator. Galvanostatic cycling tests were carried out at room temperature using an Arbin BT2000 potentiostat. EIS was performed using a CHI 660e electrochemical workstation (ChenHua Instruments Co.) between 0.01 Hz and 100 kHz. 3. Results and discussion The synthesis process of the Na@rGa composite anode is illustrated in Fig. 1a. A 3D rGa host with ordered microporous structure and abundant oxygen functional groups was formed via a mild hydrothermal method and next oriented freeze-drying. The mild water bath made 2D graphene sheets preliminary self-assemble into porous 3D structures and retained most oxygen-containing functional groups of GO precursor, these oxygen-containing functional groups raised the wettability of rGa towards molten Na metal. Next, the hydrogel was transformed into a
2. Experimental section 2.1. Synthesis Synthesis of the rGa nanostructure: rGa was synthesized via a simple mild reduction followed by freeze-drying. In a typical process, GO (50 mg, Tan Feng Tech. Inc.) was thoroughly dispersed into 20 mL 377
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Fig. 1. (a) Synthetic illustration of Na@rGa composite anode: rGO hydrogel-rGa-Na@rGa composite sample; (b) Photographs of graphene oxide solution after hydrothermal 4 h reduction; (c) Photographs of the ultra-light weight graphene aerogel; (d) Good flexibility of a Na@rGa composite strip, which is bended by two tweezers; (e, f) SEM images of rGa-4h and Na@rGa composite anode; (g) SEM images and elemental mapping images after pressing of the Na@rGa composite anode.
the intensity of O peak decreased, confirming the further reduction of rGa. It should be noted that oxygen functional groups are of great importance for the infusion of molten Na metal. Based on the above results, we chose 60 C and 4 h as the best hydrothermal conditions to achieve a balance between porous structure and wettability towards molten Na metal. Fig. 1b presents a photograph of cylindrical GO hydrogels that self-assembled during hydrothermal treatment for 4 h. As shown in Fig. 1c, the obtained rGa-4h sample was extremely light. After being deformed by pressure (200 g), the rGa-4h sample returned to its original shape (Fig. S4). Moreover, the Na@rGa composite material was highly flexible, as illustrated in Fig. 1d, which is caused by the high flexibility of the rGa-4h precursor. We used a melt infusion strategy to obtain an Na@rGa composite anode. With the large number of pore structures and excellent mechanical stability, the as-prepared rGa materials with a thickness about 5 mm was compressed to form ~300 μm-thick flexible 3D rGa film (Fig. S5). Owing to the strong cross-linking of graphene sheets in the rGa, the pressed rGa film still has excellent flexibility (Fig. S6). When the rGa film was partially immersed in molten Na metal at 300 C, molten Na was rapidly adsorbed into whole rGa host; the entire infusion process occurred within 20 s (Fig. S7). In this step, a flexible Na@rGa composite anode was quickly obtained. The anode surface exhibited metallic luster, indicating the uptake of metallic Na into the 3D porous structure of the aerogel. The successfully infusion could be explained by the capillary force derived from the microcellular structure and oxygen functional groups on the surface of the graphene sheets [9,32,42]. SEM image (Fig. 1f) and corresponding EDX of Na and C (Fig. 1g) for Na@rGa clearly show that Na metal was uniformly adsorbed in the pores of rGa, and the 3D rGa structure was well preserved even if pressed. It is worth noting that the host material only accounted for 1.75 wt% of the rGa@Na composite anode (Table S1), which is beneficial for improving the capacity of the composite anode. The chemical composition and structural characteristics of the rGa are shown in Fig. 2. Fig. 2a presents the XRD patterns of GO and rGa. In the XRD pattern obtained for rGa, the sharp peak of GO located at 12
homogeneous porous rGa by oriented freeze-drying (Fig. S1). In the freezing process, the aqueous solution in the hydrogel was fast crystallization to form acicular ice crystals which grow from the outside to the inside of hydrogel, the 3D network structure of graphene in hydrogel was reoriented under the action of ice crystal boundary press and etching. Ice crystals grew along the transverse and longitudinal directions and squeezed each other after contacting, and then compromised each other to extend to the unfrozen hydrogel region inside, eventually forming the 3D uniform porous structure [44]. It should be noted that the evaporation of the ice template during the freeze-drying process was important to form a uniform microstructure without introducing impurities. Moreover, the pore size of rGa could be accurately controlled by adjusting freezing rate [41]. When the hydrothermal reaction time of rGa was 2 h (sample denoted as rGa-2h), the self-assembly process of GO formed a hydrogel structure through the driving force of π-π bonding. At this stage, rGa-2h material contained a large number of oxygen functional groups, which is beneficial for adsorption of molten Na, but a loose structure because of insufficient self-assembly. Thus, the rGa-2h sample exhibits severe stacking and uneven pore distribution, as shown in Figs. S2a and b. When the hydrothermal reaction time was extended to 24 h (sample denoted as rGa-24h), most of the oxygen functional groups were reduced compared with the case for rGa-2h, leading to an increased degree of reduction of GO. The aerogel obtained under this condition possessed a firm structure but displayed poor wettability towards molten Na metal. Fig. 1e shows the typical structure of the aerogel with a hydrothermal reaction time of 4 h (denoted as rGa-4h), which clearly exhibits ordered microcavities with a diameter of 5 μm and rGa walls that are connected together to form 3D scaffold. Hierarchical porous structure offers large specific area and is beneficial for wettability between Na@rGa and electrolyte [22,33, 34]. Comparison with Figs. S2c and d revealed that rGa-4h had similar porous morphology to that rGa-24h, indicating that the self-assembly process was almost complete after 4 h. Detailed XPS analysis further confirmed the effect of hydrothermal reaction time on the reduction degree of rGa (Fig. S3). With lengthening hydrothermal reaction time,
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Fig. 2. (a) XRD characterizations of GO and rGa-4h after mild reduction; (b) FTIR spectra of pristine GO film and rGa-4h sample; (c) XPS spectra of C 1s of GO film and rGa-4h sample; (d) XPS spectra of O 1 s of GO film and rGa-4h sample.
disappeared and a broad peak emerged at around 24 , indicating the mild reduction of GO. A sharp peak at about 28 , corresponding to the (110) plane of Na metal, was observed in the XRD pattern which confirms the existence of Na in the rGa host (Fig. S8). Fig. 2b presents FTIR spectra of rGa and GO samples. Before the mild reduction process, a very broad peak centered around 3400 cm1, corresponding to hydroxyl (–OH) groups and surface water, was observed in the FTIR spectrum of GO. – O) at Peaks from oxygen-containing groups, including carbonyl (C– 1729 cm1 and epoxy (C–O–C) at 1053 cm1, were also detected. After mild reduction, the peaks from oxygen functional groups became very weak, confirming the partial reduction of GO. Furthermore, XPS analysis were conducted to further examine the change of GO during the reduction process. As shown in Fig. 2c and d, the intensity of the peak from C–O–C decreased from 46.8 % for GO to 5.8 % for rGa and the intensity – O dropped from 7.9 % for GO to 3.2 % for rGa. These of the peak from C– changes indicate partial reduction of oxygen functional groups, which could effectively increase the wettability of rGa [40]. To evaluate the electrochemical performance of the Na@rGa composite anode, 2032-type coin cells were assembled with two identical Na@rGa composite anodes and 1 M sodium perchlorate (NaClO4) in ethylene carbonate/diethyl carbonate (EC/DEC, 1:1 v/v) as the electrolyte for galvanostatic cycling measurements. It should be noted that using carbonate-based electrolytes in SMBs still faces many challenges because of the formation of an unstable SEI layer and their comparatively positive reduction potential. However, optimization of the SMB electrolyte, although very meaningful, is beyond the scope of this study. For comparison, symmetric cells using bare Na metal were assembled. Fig. 3 compares the time-dependent voltage profiles of cells with Na@rGa and bare Na electrodes at current densities of 0.5, 3, and 5 mA cm2 within a protective cutoff voltage window of 3.0 to 3.0 V. As shown in Fig. 3a, at a current density of 0.5 mA cm2 with a capacity of 0.5 mAh cm2, the voltage hysteresis of the bare Na cell gradually increased and short circuited after 250 h of cycling. In contrast, the Na@rGa composite anode maintained stable cycling for over 350 h and presented smaller average overpotential (35 mV) than that of the Na reference cell. When the deposition capacity was increased to 1 mAh cm2 (Fig. S9), the cell with
Na@rGa composite electrodes also exhibited flat plating/stripping profiles and a low overpotential of about 35 mV over 350 h. When the current density was raised to 3 mA cm2 with a capacity of 2 mAh cm2 (Fig. 3b), the Na@rGa composite anode exhibited stable cycling with an overpotential of only about 80 mV for over 120 h, whereas the symmetric cell with bare Na electrodes failed after 82.5 h. Compared with the flat voltage plateaus of the Na@rGa cell, the voltage response of the bare Na cell showed an obvious fluctuation, which was attributed to the unstable SEI on the Na surface caused by the large volume changes during cycling of high current density and capacity. The plating/stripping of Na ions inevitably resulted in large volume expansion and contraction. The Na metal newly exposed by the cracks in the SEI reacted with the electrolyte to generate a new SEI, this process leads to continuous consumption of Na metal and the liquid electrolyte during cycling of the reference bare Na cell. When the current density was further increased to 5 mA cm2 with a cycling capacity of 1 mAh cm2 (Fig. 3c), the bare Na cell presented a very short cycle life (about 22 h) and much higher overpotential (>250 mV) compared with those of the Na@rGa composite anode. The main reason for the high voltage hysteresis and short cycle life of the bare Na metal electrode was the inhomogeneous nucleation of Na ions during plating/stripping at 5 mA cm2. By contrast, the rGa host provides abundant electrochemically active sites for Na deposition, which effectively lower the actual current density, and its 3D porous structure can mitigate the volume expansion/contraction of Na metal during the plating/stripping process and enhance wettability between Na@rGa and electrolyte. The Na@rGa composite anode still exhibited a stable voltage profile with a low polarization (110 mV) at 5 mA cm2 with a capacity of 5 mAh cm2 (Fig. S10). These advantages endow the Na@rGa composite anode with superior performance, including a long cycle life of over 1000 cycles and low overpotential of about 50 mV even at a current density of 5 mA cm2, as shown in Fig. 3c. To further investigate cycling stability and the location of the Na plating during cycling, SEM images of the bare Na and Na@rGa composite anodes were collected. Fig. 4c and d show the surface morphology of the Na@rGa composite anode after 80 cycles at a current density of 5 mA cm2 and capacity of 1 mAh cm2. The surface is smooth and
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Fig. 3. Electrochemical performances of symmetric cells using Na@rGa composite anode and bare Na metal anode at (a) 0.5 mA cm2 with capacity of 0.5 mAh cm2, (b) 3 mA cm2 with capacity of 2 mAh cm2 and (c) 5 mA cm2 with capacity of 1 mAh cm2.
(~160 Ω after 20 cycles and ~190 Ω after 50 cycles). Conversely, for the bare Na anode, interfacial resistance obviously increased during plating/ stripping cycles (~210 Ω after 20 cycles and ~780 Ω after 50 cycles). This behavior means that compared with that of the bare Na metal, the Na@rGa composite anode shows a more stable interfacial surface during cycling, which could be ascribed to the rGa host. The homogeneous pores in the host can ensure much better electrode stability and more uniform electrolyte distribution from the inner to the outside of electrode to enhance Na stripping/plating kinetic. While the interfacial resistance of bare Na metal anode increased to a much higher value of ~780 Ω, indicating that thick SEI layer grown on the surface and huge volume changes during repeated stripping/plating. Overall, the SEM and EIS results, along with the illustrations in Fig. 4a and b, clearly show that the Na@rGa composite anode with a porous 3D network has much higher stability and more favorable Na stripping/plating kinetics than those of bare Na. Infinite volume expansion and dendrite growth are catastrophic problems that currently limit the development of alkali metal anodes. Here, we assembled symmetric cells of Na@rGa and bare Na foil (initial thickness of 320 μm) to investigate stripping/plating behaviors at a capacity of 5 mAh cm2 and current density of 1 mA cm2. Fig. S11a shows the thickness variation of bare Li and bare Na during plating/stripping. Because of their different densities and theoretical capacities, when
homogenous without Na dendrites or obvious protuberances. This desirable surface morphology after cycling is ascribed to the 3D aerogel host with a homogeneous Na ion distribution on its surface, which resulted from the lower local current density and more uniform distribution of the ionic flux during Na plating/stripping than was the case for the bare Na electrode. As a contrast, Fig. 4e and f show the surface morphology of the bare Na anode after cycling under the same conditions as the Na@rGa composite anode. The bare Na anode possessed an uneven surface and mossy dendrites, which are attributed to the non-uniform local current density. This leads to uneven Na plating/stripping and high reaction kinetics on the bare Na surface. As expected, the bare Na foil surface has a rough surface with an unstable SEI layer and dendrite formation. These results further explain the improved cycling stability of the Na@rGa composite anode compared with that of the bare Na anode (Fig. 3). Fig. 4g and h present the EIS spectra of the Na@rGa composite and bare Na anode after 20 and 50 cycles of stripping/plating at a current density of 5 mA cm2 and capacity of 1 mAh cm2. The SEI and charge transfer resistance at the Na surface are associated with the high frequency semicircles in the Nyquist plots. As illustrated in Fig. 4g, the Nyquist plots for the Na@rGa composite anode showed no obvious increase of interfacial impedance after 50 cycles of N stripping/plating 380
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Fig. 4. (a, b) The schematic illustration of stripping and plating behavior of dendritefree Na@rGa composite anode and bare Na metal anode. (c–f) SEM images of Na@rGa and bare Na metal anode after 80 charge/ discharge cycles at 5 mA cm2 current density. (g, h) Nyquist plots of Na@rGa composite anode and bare Na after 20 cycles and 50 cycles at a current density of 5 mA cm2. (i) Top view (i) and cross-sectional (j) of the Na@rGa composite anode after plating 5 mAh cm2 at 1 mA cm2 (k) Top views of the Na@rGa composite anode after stripping 5 mAh cm2 at 1 mA cm2.
composite materials remained dense structure and the total thickness keeps at 320 μm after plating/stripping 5 mAh cm2 (Fig. S11d). These results confirm the Na@rGa composite anode can effectively inhibit volume expansion during plating/stripping. More important, Na@rGa composite anode also exhibits competitive high capacity compared with bare Na foil. As shown in Fig. 5a, a capacity of 1064 mAh g1 (based on the total weight of the electrode) was extracted when the Na@rGa composite anode was charged to 1 V. This capacity is very close to the theoretical capacity of bare Na (1165 mAh g1), suggesting that rGa host ensures stable cycling of composite anode and without loss of capacity. The high specific capacity can be attributed to the ultra-light weight and porous structure of rGa, which ensured high loading of Na metal in the whole composite anode (98.25 wt%) and large storage space. Therefore, aerogel hosts offer an exciting possibility of fabricating high-performance composite anodes with almost no adverse effect on capacity. In addition, we employed Na0.67Ni0.25Mn0.75O2 (NNM) as cathode and the Na@rGa as anode to assemble a full cell (NNM-Na@rGa) using 1 M NaClO4 in EC/DEC (1:1 in volume) as electrolyte. For comparison, the NNM cathode was also assembled with a bare Na anode as a reference (NNM-Na). Fig. 5b showed the galvanostatic charge-discharge profiles of NNM-Na@rGa and NNM-Na coin cells at 2 C rate for 300 cycle. NNMNa@rGa cell exhibits higher capacity compared with NNM-Na cell (62
plating/stripping the same areal capacity, the thickness variation of Na is larger than that of bare Li. Bare Na metal exhibits obvious rupture in a top-views SEM image (Fig. S11b) and thickness variation (about 42 μm) during plating at 5 mAh cm2, this behavior is roughly in accordance with previous calculation results, where volume variation caused by “hostless” nature of Na metal leads to rupture of the SEI and a large amount of “dead Na” during cycling (Fig. S11c). The “dead Na” hinders the contact between electrolyte and electrode and impede the ion transport rate, increase the overpotential, and cause serious fluctuation during cycling. As shown in Fig. 3i, the surface of the composite electrode is smoother and denser than the Na anode. The Na@rGa anode possesses uniform pores and a large specific surface area, which effectively lower the local current density and promote uniform Na nucleation. The surface morphology of the Na@rGa electrode is maintained well during cycling. Fig. 3j show the thickness variation of the Na@rGa anode during plating at a capacity of 5 mAh cm2. The thickness of the Na@rGa electrode increases by only 10 μm after plating. This is accordance with the results in Fig. 3i and suggests the Na@rGa composite anode can accommodate the volume change of Na electrode when the plating capacity is about 5 mAh cm2. After 1 stripping/plating cycle, the rGa structure is gradually exposed and shows obvious porous structure (Fig. 3k). Although the Na@rGa anode surface exhibited a different morphology compared with the initial stage, no obvious dendritic Na can be observed. Moreover, the 381
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Fig. 5. Electrochemical characterization of the Na@rGa electrodes: (a) The full Na stripping curve of the Na@rGa electrode to 1 V versus Naþ/Na at a current density of 25 mA g1, which shows a specific capacity of ~1064 mAh g1; (b) Galvanostatic discharging and charging curve of NNMNa@rGa anode at a rate of 2 C for 300 cycle. (c) Rate capability of the NNMNa@rGa cell at different rates from 0.2 to 10 C, and coulombic efficiency of NNMNa@rGa electrodes. (d) Cycling performance and coulombic efficiency of NNMNa@rGa electrode at a rate of 0.1 C.
mAh g1 vs 47 mAh g1). As can be seen from Fig. 5c, the rate capability of NNM-Na@rGa also presents reversible capacity of 73, 71, 68, 62, 51, 23 mAh g1 at the rates of 0.2, 0.5, 1, 2, 4, 10 C. Respectively. A capacity of 70 mAh g1 at 0.5 C and a high coulombic efficiency of 99.8 % are still obtained after 110 cycles. Meanwhile, NNM-Na@rGa cell exhibits more stable electrochemical performance over 100 cycles at the current density of 0.1 C with reversible capacity of 79 mAh g1 and coulombic efficiency of 99.5%. These results further illustrate that Na@rGa electrode still has a good adaptability in the full battery.
Appendix A. Supplementary data Supplementary data to this article can be found online at https://do i.org/10.1016/j.ensm.2019.02.015. References [1] J.B. Goodenough, K.S. Park, The Li-ion rechargeable battery: a perspective, J. Am. Chem. Soc. 135 (2013) 1167–1176. [2] K. Xu, Electrolytes and interphases in Li-ion batteries and beyond, Chem. Rev. 114 (2014) 11503–11618. [3] R. Luo, F. Wu, M. Xie, Y. Ying, J. Zhou, Y. Huang, Y. Ye, L. Li, R. Chen, Habit planedriven P2-type manganese-based layered oxide as long cycling cathode for Na-ion batteries, J. Power Sources 383 (2018) 80–86. [4] Y. Huang, M. Xie, J. Zhang, Z. Wang, Y. Jiang, G. Xiao, S. Li, L. Li, F. Wu, R. Chen, A novel border-rich Prussian blue synthetized by inhibitor control as cathode for sodium ion batteries, Nano Energy 39 (2017) 273–283. [5] M. Lee, J. Hong, J. Lopez, Y. Sun, D. Feng, K. Lim, W.C. Chueh, M.F. Toney, Y. Cui, Z. Bao, High-performance sodium–organic battery by realizing four-sodium storage in disodium rhodizonate, Nat. Energy 2 (2017) 861–868. [6] N. Nitta, F. Wu, J. Lee, Y. Gleb, Li-ion battery materials: present and future, Mater. Today 5 (2018) (2015) 252–264. [7] A. Metrot, D. Guerard, D. Billaud, A. Herold, New results about the sodium-graphite system, Synth. Met. 1 (1980) 363–369. [8] M. Camelia Ghimbeu, Joanna G orka, V. Simone, Insights on the Naþ ion storage mechanism in hard carbon: discrimination between the porosity, surface functional groups and defects, Nano Energy 44 (2018) 327–335. [9] Y. Cao, L. Xiao, L. Maria Sushko, W. Wang, B. Schwenzer, J. Xiao, Z. Nie, V. Laxmikant Saraf, Z. Yang, J. Liu, Sodium ion insertion in hollow carbon nanowires for battery applications, Nano Lett. 12 (2016) 3783–3787. [10] J. Xu, M. Wang, N.P. Wickramaratne, M. Jaroniec, S. Dou, L. Dai, High-performance sodium ion batteries based on a 3D anode from nitrogen-doped graphene foams, Adv. Mater. 27 (2015) 2042–2048. [11] Y. Kim, Y. Park, A. Choi, N.S. Choi, J. Kim, J. Lee, H.R. Ji, S.M. Oh, K.T. Lee, Composites: an amorphous red phosphorus/carbon composite as a promising anode material for sodium ion batteries, Adv. Mater. 25 (2013), 3010-3010. [12] J. Sun, H.W. Lee, M. Pasta, H. Yuan, G. Zheng, Y. Sun, Y. Li, Y. Cui, A phosphorene–graphene hybrid material as a high-capacity anode for sodium-ion batteries, Nat. Nanotechnol. 10 (2015) 980–985. [13] Z. Li, J. Ding, D. Mitlin, Tin and tin compounds for sodium ion battery anodes: phase transformations and performance, Acc. Chem. Res. 48 (2015) 1657–1665. [14] B. Kong, L. Zu, C. Peng, Y. Zhang, W. Zhang, J. Tang, C. Selomulya, L. Zhang, H. Chen, Y. Wang, Direct superassemblies of freestanding metal–carbon frameworks featuring reversible crystalline-phase transformation for electrochemical sodium storage, J. Am. Chem. Soc. 138 (2016) 16533–16541. [15] A.P. Cohn, N. Muralidharan, R. Carter, K. Share, C.L. Pint, Anode-free sodium battery through in situ plating of sodium metal, Nano Lett. 17 (2017) 1296–1301. [16] C. Xia, R. Fernandes, F.H. Cho, N. Sudhakar, B. Buonacorsi, S. Walker, M. Xu, J. Baugh, L.F. Nazar, Direct evidence of solution-mediated superoxide transport and organic radical formation in sodium-oxygen batteries, J. Am. Chem. Soc. 138 (2016) 11219–11226.
4. Conclusion In summary, we successfully constructed Na@rGa composite anode via infusing molten Na metal into 3D graphene aerogel. Herein, we employed low-temperature hydrothermal to control the self-assembly of GO accurately. Because of its advantages of uniform porous structure, ultra-light quality and good wettability with both Na and electrolyte, rGa is an ideal choice as host for further application. Based on the above characteristics, the ultra-light host ensures a high energy density of Na@rGa composite anode, and the Na@rGa anode present excellent cycling performance in carbonate-based electrolyte without using any additives. At high current density of 5 mA cm2, the Na@rGa composite anode exhibited a highly cycling stability with no dendrite morphology and low hysteresis (50 mV) even after 1000 cycles which show an excellent electrochemical performance (Table S2). In addition, full-cell battery (NNM-Na@rGa) presents enhanced cycling stability. Our work provides new idea for the construction of 3D wettable scaffolds and principle for the facile applications of safe Na metal anodes. Meanwhile, it could be a viable strategy for other alkali metallic anode systems. Data availability The raw/processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations. Acknowledgements This work was supported by the National Key Research and Development Program (2016YFB0901501), the National Natural Science Foundation of China (51772030). 382
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Contents lists available at ScienceDirect
Energy Storage Materials journal homepage: www.elsevier.com/locate/ensm
Reduced graphene oxide aerogel as stable host for dendrite-free sodium metal anode Feng Wu a, b, Jiahui Zhou a, Rui Luo a, Yongxin Huang a, Yang Mei a, Man Xie a, **, Renjie Chen a, b, * a b
School of Materials Science & Engineering, Beijing Key Laboratory of Environmental Science and Engineering, Beijing Institute of Technology, Beijing 100081, China Collaborative Innovation Center of Electric Vehicles in Beijing, Beijing 100081, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Sodium metal anode Graphene aerogel Oriented freeze-drying Long-term stability High current density
Sodium (Na) metal has attracted great attention as a promising anode for next-generation energy storage systems because of its abundant resources, potentially low cost, and high theoretical capacity. However, severe dendrite growth, large volume expansion during plating/stripping lead to poor cycle performance and limit the practical application of sodium metal anodes. Here, oriented freeze-drying and molten infusion are used to synthesize a Nainfused reduced graphene oxide aerogel (Na@rGa) composite anode using a reduced graphene oxide aerogel (rGa) as a stable host. This unique host shows ultra-light weight which guarantees high capacity (1064 mAh g1), and the large specific surface area and uniform pores effectively lower the local current density and uniform electrolyte distribution from the inner to the outside of electrode. The Na@rGa anode exhibits low overpotential (50 mV) and stable cycle performance for 1000 cycles at 5 mA cm2 in carbonate electrolyte system. The electrochemical performance of a full cell using the Na@rGa composite anode is clearly superior to that of a reference cell. This work provides a new horizon for construction of three-dimensional stable hosts and safe sodium metal anodes.
1. Introduction With the increasing demand for large-scale energy storage, sodiumion batteries (SIBs) have attracted tremendous attention as alternatives to lithium-ion batteries (LIBs) because of their abundant resources, low cost, and analogous battery chemistry to that of LIBs [1,2]. Progress has been achieved in terms of cathode materials for SIBs including layered transition metal oxides, Prussian blue, organic polymers, and polyanions [3–6]. However, research on anode materials for SIB still faces great challenges, mainly because the commercial graphite used in LIBs cannot be applied to SIBs [7,8]. Considerable effort has been devoted to investigating high-capacity materials for SIBs, including carbon materials [9, 10], phosphorus [11,12], and transition metal oxides and alloys [13,14]. Compared to these anode candidates, metallic Na possesses high theoretical capacity (1165 mAh g1) and low potential (2.714 V vs. the standard hydrogen electrode) [15]. Furthermore, the investigation of Na metal anodes is of great importance for batteries with high energy density such as Na–O2 [16,17] and Na–S batteries [18].
Despite all the advantages mentioned above, secondary batteries that use Na metal as anodes still face severe problems, especially in carbonate electrolyte [19]. Uneven nucleation, dendrite growth of Na during the plating process and uncontrollable side reactions between fresh Na with the liquid electrolyte and formation of “dead Na” in the following stripping process, leading to short cycle life and poor electrochemical performance [20–22]. Moreover, non-skeleton Na metal shows a huge volume change during plating and stripping, leading to fracture of solid electrolyte interface (SEI) and severe side reactions. The development of sodium metal batteries (SMBs) is still in its initial stage. At present, dendrite growth and volume expansion of the Na metal anode during cycling are the principal barriers limiting the further development of the SMBs [23]. Several strategies used to realize lithium (Li) metal anodes with long cycle lives have been adapted to solving the above-mentioned problems of Na metal anodes, mainly including fabricating organic and inorganic thin layer [24], adding liquid electrolyte additives [25] and using highly concentrated Sodium bis (fluorosulfonyl) imide in 1,2-dimeth-oxyethane
* Corresponding author. School of Materials Science & Engineering, Beijing Key Laboratory of Environmental Science and Engineering, Beijing Institute of Technology, Beijing 100081, China. ** Corresponding author. School of Materials Science ＆ Engineering, Beijing Key Laboratory of Environmental Science and Engineering, Beijing Institute of Technology, Beijing 100081, China. E-mail addresses: [email protected] (M. Xie), [email protected] (R. Chen). https://doi.org/10.1016/j.ensm.2019.02.015 Received 25 December 2018; Received in revised form 18 February 2019; Accepted 18 February 2019 Available online 20 February 2019 2405-8297/© 2019 Elsevier B.V. All rights reserved.
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deionized (DI) water by probe sonication for 60 min. Ascorbic acid (50 mg) was added to the solution, which was then continuously stirred for 2 h at room temperature. The solution was then transferred into a cylindrical bottle, sealed and heated in a water bath at 60 C for 4 h. After solvothermal reaction, the resulting cylindrical-shaped hydrogel was collected, washed with a mixture of ethanol and DI water (1:10 v/v) several times to remove unreacted chemicals, and then immersed in ethanol/DI water (1:10 v/v) for 12 h. The collected hydrogel was hanged in refrigerator about 6 h to orient freeze and obtained by freeze-drying overnight in order to remove absorbed water and ethanol. The obtain aerogel was cut into thin slices and placed in a glove box. For comparison, an aerogel foam was also synthesized with the same method by extending the reaction time in the water bath from 2 to 24 h under the same conditions. Na infusion into rGa: Na infusion process was carried out in a glove box. Solid Na cube was melted on a hot plate, and as-prepared rGa slices were then immersed into molten Na. The infusion process was finished in 20 s and 3D Na@rGa composite anode was obtained. Synthesis of P2–Na0.67Ni0.25Mn0.75O2 (P2-NNM) electrode: The P2NNM cathode material was synthesized by the solid reaction described in our previous report [3] and summarized as follows. An appropriate ratio of precursor powders (Na2CO3, NiO, and Mn2O3) was thoroughly mixed in a high-energy ball mill for 6 h. The mixture was then calcined twice at 900 C for 16 h in air atmosphere with grinding between the calcination steps. The calcined powder was then ground and stored in an Ar-filled glovebox until use. P2-NNM electrodes were prepared by casting a slurry of active material, acetylene black, and PVDF binder (75:15:10 w/w/w) on an aluminum foil collector.
electrolyte [26]. These methods effectively inhibit the growth of Na dendrites by forming artificial SEI films on the Na surface. However, in most cases, surface protection alone cannot inhibit the large volume expansion of the Na metal anode during electrochemical cycling, especially at high current density (>3 mA cm2). Meanwhile, the artificial SEI film inevitably increases interface resistance and restricts the migration of Na ions. Several attempts to develop various matrixes as hosts for Na/Li metal with the aim of alleviating both volume changes and dendrite growth have been explored [27–31]. Fan's group developed a three-dimensional (3D) carbon felt as a host for Na metal, which displayed relatively uniform and stable Na plating/stripping behavior [27]. Hu's group fabricated a carbonized wood composite as a high-surface-area, conductive, stable Na metal anode that tackled nearly all the problems associated with SMBs [28]. Chen and co-workers reported an Na/graphene film composite anode that exhibited great stability in Na–CO2 batteries [38]. These results suggest that confining Na metal in a 3D host could greatly enhance the stability of Na metal anodes. However, most available matrixes have poor wettability towards molten Na metal. Researchers have tried to decrease the contact angle of matrixes and solve the wetting issue between molten alkali metal and matrixes by coating with a metal oxide (ZnO [32], Al2O3 [33], and SnO2 [34]) and adding alloying elements (Ge [35], Al [36], and Mg [37]). However, these approaches inevitably caused a loss of capacity, in addition, the adsorption of Na in the host reduces the contact area between the Na and the electrolyte. The electrolyte does not sufficiently wet the entire electrode which reduces the cycle stability and enhances the possibility of dendrite formation after long cycling. Developing a “sodiophilic” host structure with high-loading and good wettability with electrolyte is highly desirable to achieve high-performance SMBs. An ideal host for Na should possess flexibility, ultra-light weight to obtain a composite electrode with high capacity, large specific surface area to lower the local current density, good chemical stability to avoid side reactions with liquid electrolyte, good Na and electrolyte wettability. rGa is a typical 3D porous framework that has attracted extensive attention because of its interconnected hierarchical porous system, large specific surface area, excellent adsorption in organic solvents, and ultralight weight [39–41]. Furthermore, the porous framework of rGa aids fast ion transfer inside the aerogel and enables capillary action during the melt infusion process [32,42]. These unique features make rGa promising as a carbon host for encapsulating Na metal. Notably, the self-assembly process of rGa is accompanied by the reduction of oxygen functional groups, which is of important for Na metal infusion [43]. Herein, we used a modified hydrothermal reduction reaction and subsequent oriented freeze-drying process to form rGa host with ultralight weight and good wettability. Abundant oxygen functional groups of the rGa aid infusion of molten Na to form Na@rGa composite, and oriented freeze-drying ensures that rGa has a highly homogeneous pore structure. Benefiting from these unique structure and chemical property, the obtained Na@rGa composite anode shows an ultra-high loading of Na metal of 98.75 wt% and capacity of 1064 mAh g1. Moreover, the homogeneous pores enhance wettability between Na@rGa and electrolyte, which is crucial for high ionic mobility and uniform electrolyte distribution from the inner to the outside of electrode, the large specific surface area reduce local current density and inhibit the nucleation of Na dendrite. The Na@rGa anode exhibits a low overpotential of about 50 mV at 5 mA cm2 in a traditional carbonate electrolyte without any additives over 1000 cycles. The electrochemical performance of a full battery with the Na@rGa composite anode is also investigated.
2.2. Structural characterization and electrochemical measurements Structural Characterization: The lattice structure of Na@rGa was recorded using a D8 Advance diffractometer (XRD, Bruker AXS, WI, USA) with a scanning range from 10 to 70 . The morphologies images of the sample were taken using a Hitachi SU-70 field-emission scanning electron microscope (SEM) with an energy-dispersive X-ray (EDX) spectroscopy attachment and accelerating voltage of 5.0 kV. To observe the surface morphology of the Na@rGa composite electrode and bare Na foil after cycling, the cycled cells were disassembled in a glove box and then washed with dimethyl carbonate several times to remove surface impurities. Fourier transform infrared (FTIR) spectra were recorded on a Nicolet iS50 FTIR spectrometer (Thermo Scientific). X-ray photoelectron spectra (XPS) measurements were recorded on a spectrometer with Al Kα radiation (hv ¼ 1486.6 eV). Electrochemical Measurement: To investigate the electrochemical plating/stripping of the Na@rGa composite electrode, 2032-type coin cells were assembled with two identical Na@rGa electrodes. For comparison, symmetric cells were fabricated with two bare Na metal foil electrodes. For full cell measurements, P2-NNM electrodes were paired with Na@rGa and bare Na. All the coin cells were assembled using 1.0 M NaClO4 in EC/DEC (1:1 v/v) solution as the electrolyte without any additives and glass fiber as the separator. Galvanostatic cycling tests were carried out at room temperature using an Arbin BT2000 potentiostat. EIS was performed using a CHI 660e electrochemical workstation (ChenHua Instruments Co.) between 0.01 Hz and 100 kHz. 3. Results and discussion The synthesis process of the Na@rGa composite anode is illustrated in Fig. 1a. A 3D rGa host with ordered microporous structure and abundant oxygen functional groups was formed via a mild hydrothermal method and next oriented freeze-drying. The mild water bath made 2D graphene sheets preliminary self-assemble into porous 3D structures and retained most oxygen-containing functional groups of GO precursor, these oxygen-containing functional groups raised the wettability of rGa towards molten Na metal. Next, the hydrogel was transformed into a
2. Experimental section 2.1. Synthesis Synthesis of the rGa nanostructure: rGa was synthesized via a simple mild reduction followed by freeze-drying. In a typical process, GO (50 mg, Tan Feng Tech. Inc.) was thoroughly dispersed into 20 mL 377
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Fig. 1. (a) Synthetic illustration of Na@rGa composite anode: rGO hydrogel-rGa-Na@rGa composite sample; (b) Photographs of graphene oxide solution after hydrothermal 4 h reduction; (c) Photographs of the ultra-light weight graphene aerogel; (d) Good flexibility of a Na@rGa composite strip, which is bended by two tweezers; (e, f) SEM images of rGa-4h and Na@rGa composite anode; (g) SEM images and elemental mapping images after pressing of the Na@rGa composite anode.
the intensity of O peak decreased, confirming the further reduction of rGa. It should be noted that oxygen functional groups are of great importance for the infusion of molten Na metal. Based on the above results, we chose 60 C and 4 h as the best hydrothermal conditions to achieve a balance between porous structure and wettability towards molten Na metal. Fig. 1b presents a photograph of cylindrical GO hydrogels that self-assembled during hydrothermal treatment for 4 h. As shown in Fig. 1c, the obtained rGa-4h sample was extremely light. After being deformed by pressure (200 g), the rGa-4h sample returned to its original shape (Fig. S4). Moreover, the Na@rGa composite material was highly flexible, as illustrated in Fig. 1d, which is caused by the high flexibility of the rGa-4h precursor. We used a melt infusion strategy to obtain an Na@rGa composite anode. With the large number of pore structures and excellent mechanical stability, the as-prepared rGa materials with a thickness about 5 mm was compressed to form ~300 μm-thick flexible 3D rGa film (Fig. S5). Owing to the strong cross-linking of graphene sheets in the rGa, the pressed rGa film still has excellent flexibility (Fig. S6). When the rGa film was partially immersed in molten Na metal at 300 C, molten Na was rapidly adsorbed into whole rGa host; the entire infusion process occurred within 20 s (Fig. S7). In this step, a flexible Na@rGa composite anode was quickly obtained. The anode surface exhibited metallic luster, indicating the uptake of metallic Na into the 3D porous structure of the aerogel. The successfully infusion could be explained by the capillary force derived from the microcellular structure and oxygen functional groups on the surface of the graphene sheets [9,32,42]. SEM image (Fig. 1f) and corresponding EDX of Na and C (Fig. 1g) for Na@rGa clearly show that Na metal was uniformly adsorbed in the pores of rGa, and the 3D rGa structure was well preserved even if pressed. It is worth noting that the host material only accounted for 1.75 wt% of the rGa@Na composite anode (Table S1), which is beneficial for improving the capacity of the composite anode. The chemical composition and structural characteristics of the rGa are shown in Fig. 2. Fig. 2a presents the XRD patterns of GO and rGa. In the XRD pattern obtained for rGa, the sharp peak of GO located at 12
homogeneous porous rGa by oriented freeze-drying (Fig. S1). In the freezing process, the aqueous solution in the hydrogel was fast crystallization to form acicular ice crystals which grow from the outside to the inside of hydrogel, the 3D network structure of graphene in hydrogel was reoriented under the action of ice crystal boundary press and etching. Ice crystals grew along the transverse and longitudinal directions and squeezed each other after contacting, and then compromised each other to extend to the unfrozen hydrogel region inside, eventually forming the 3D uniform porous structure [44]. It should be noted that the evaporation of the ice template during the freeze-drying process was important to form a uniform microstructure without introducing impurities. Moreover, the pore size of rGa could be accurately controlled by adjusting freezing rate [41]. When the hydrothermal reaction time of rGa was 2 h (sample denoted as rGa-2h), the self-assembly process of GO formed a hydrogel structure through the driving force of π-π bonding. At this stage, rGa-2h material contained a large number of oxygen functional groups, which is beneficial for adsorption of molten Na, but a loose structure because of insufficient self-assembly. Thus, the rGa-2h sample exhibits severe stacking and uneven pore distribution, as shown in Figs. S2a and b. When the hydrothermal reaction time was extended to 24 h (sample denoted as rGa-24h), most of the oxygen functional groups were reduced compared with the case for rGa-2h, leading to an increased degree of reduction of GO. The aerogel obtained under this condition possessed a firm structure but displayed poor wettability towards molten Na metal. Fig. 1e shows the typical structure of the aerogel with a hydrothermal reaction time of 4 h (denoted as rGa-4h), which clearly exhibits ordered microcavities with a diameter of 5 μm and rGa walls that are connected together to form 3D scaffold. Hierarchical porous structure offers large specific area and is beneficial for wettability between Na@rGa and electrolyte [22,33, 34]. Comparison with Figs. S2c and d revealed that rGa-4h had similar porous morphology to that rGa-24h, indicating that the self-assembly process was almost complete after 4 h. Detailed XPS analysis further confirmed the effect of hydrothermal reaction time on the reduction degree of rGa (Fig. S3). With lengthening hydrothermal reaction time,
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Fig. 2. (a) XRD characterizations of GO and rGa-4h after mild reduction; (b) FTIR spectra of pristine GO film and rGa-4h sample; (c) XPS spectra of C 1s of GO film and rGa-4h sample; (d) XPS spectra of O 1 s of GO film and rGa-4h sample.
disappeared and a broad peak emerged at around 24 , indicating the mild reduction of GO. A sharp peak at about 28 , corresponding to the (110) plane of Na metal, was observed in the XRD pattern which confirms the existence of Na in the rGa host (Fig. S8). Fig. 2b presents FTIR spectra of rGa and GO samples. Before the mild reduction process, a very broad peak centered around 3400 cm1, corresponding to hydroxyl (–OH) groups and surface water, was observed in the FTIR spectrum of GO. – O) at Peaks from oxygen-containing groups, including carbonyl (C– 1729 cm1 and epoxy (C–O–C) at 1053 cm1, were also detected. After mild reduction, the peaks from oxygen functional groups became very weak, confirming the partial reduction of GO. Furthermore, XPS analysis were conducted to further examine the change of GO during the reduction process. As shown in Fig. 2c and d, the intensity of the peak from C–O–C decreased from 46.8 % for GO to 5.8 % for rGa and the intensity – O dropped from 7.9 % for GO to 3.2 % for rGa. These of the peak from C– changes indicate partial reduction of oxygen functional groups, which could effectively increase the wettability of rGa [40]. To evaluate the electrochemical performance of the Na@rGa composite anode, 2032-type coin cells were assembled with two identical Na@rGa composite anodes and 1 M sodium perchlorate (NaClO4) in ethylene carbonate/diethyl carbonate (EC/DEC, 1:1 v/v) as the electrolyte for galvanostatic cycling measurements. It should be noted that using carbonate-based electrolytes in SMBs still faces many challenges because of the formation of an unstable SEI layer and their comparatively positive reduction potential. However, optimization of the SMB electrolyte, although very meaningful, is beyond the scope of this study. For comparison, symmetric cells using bare Na metal were assembled. Fig. 3 compares the time-dependent voltage profiles of cells with Na@rGa and bare Na electrodes at current densities of 0.5, 3, and 5 mA cm2 within a protective cutoff voltage window of 3.0 to 3.0 V. As shown in Fig. 3a, at a current density of 0.5 mA cm2 with a capacity of 0.5 mAh cm2, the voltage hysteresis of the bare Na cell gradually increased and short circuited after 250 h of cycling. In contrast, the Na@rGa composite anode maintained stable cycling for over 350 h and presented smaller average overpotential (35 mV) than that of the Na reference cell. When the deposition capacity was increased to 1 mAh cm2 (Fig. S9), the cell with
Na@rGa composite electrodes also exhibited flat plating/stripping profiles and a low overpotential of about 35 mV over 350 h. When the current density was raised to 3 mA cm2 with a capacity of 2 mAh cm2 (Fig. 3b), the Na@rGa composite anode exhibited stable cycling with an overpotential of only about 80 mV for over 120 h, whereas the symmetric cell with bare Na electrodes failed after 82.5 h. Compared with the flat voltage plateaus of the Na@rGa cell, the voltage response of the bare Na cell showed an obvious fluctuation, which was attributed to the unstable SEI on the Na surface caused by the large volume changes during cycling of high current density and capacity. The plating/stripping of Na ions inevitably resulted in large volume expansion and contraction. The Na metal newly exposed by the cracks in the SEI reacted with the electrolyte to generate a new SEI, this process leads to continuous consumption of Na metal and the liquid electrolyte during cycling of the reference bare Na cell. When the current density was further increased to 5 mA cm2 with a cycling capacity of 1 mAh cm2 (Fig. 3c), the bare Na cell presented a very short cycle life (about 22 h) and much higher overpotential (>250 mV) compared with those of the Na@rGa composite anode. The main reason for the high voltage hysteresis and short cycle life of the bare Na metal electrode was the inhomogeneous nucleation of Na ions during plating/stripping at 5 mA cm2. By contrast, the rGa host provides abundant electrochemically active sites for Na deposition, which effectively lower the actual current density, and its 3D porous structure can mitigate the volume expansion/contraction of Na metal during the plating/stripping process and enhance wettability between Na@rGa and electrolyte. The Na@rGa composite anode still exhibited a stable voltage profile with a low polarization (110 mV) at 5 mA cm2 with a capacity of 5 mAh cm2 (Fig. S10). These advantages endow the Na@rGa composite anode with superior performance, including a long cycle life of over 1000 cycles and low overpotential of about 50 mV even at a current density of 5 mA cm2, as shown in Fig. 3c. To further investigate cycling stability and the location of the Na plating during cycling, SEM images of the bare Na and Na@rGa composite anodes were collected. Fig. 4c and d show the surface morphology of the Na@rGa composite anode after 80 cycles at a current density of 5 mA cm2 and capacity of 1 mAh cm2. The surface is smooth and
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Fig. 3. Electrochemical performances of symmetric cells using Na@rGa composite anode and bare Na metal anode at (a) 0.5 mA cm2 with capacity of 0.5 mAh cm2, (b) 3 mA cm2 with capacity of 2 mAh cm2 and (c) 5 mA cm2 with capacity of 1 mAh cm2.
(~160 Ω after 20 cycles and ~190 Ω after 50 cycles). Conversely, for the bare Na anode, interfacial resistance obviously increased during plating/ stripping cycles (~210 Ω after 20 cycles and ~780 Ω after 50 cycles). This behavior means that compared with that of the bare Na metal, the Na@rGa composite anode shows a more stable interfacial surface during cycling, which could be ascribed to the rGa host. The homogeneous pores in the host can ensure much better electrode stability and more uniform electrolyte distribution from the inner to the outside of electrode to enhance Na stripping/plating kinetic. While the interfacial resistance of bare Na metal anode increased to a much higher value of ~780 Ω, indicating that thick SEI layer grown on the surface and huge volume changes during repeated stripping/plating. Overall, the SEM and EIS results, along with the illustrations in Fig. 4a and b, clearly show that the Na@rGa composite anode with a porous 3D network has much higher stability and more favorable Na stripping/plating kinetics than those of bare Na. Infinite volume expansion and dendrite growth are catastrophic problems that currently limit the development of alkali metal anodes. Here, we assembled symmetric cells of Na@rGa and bare Na foil (initial thickness of 320 μm) to investigate stripping/plating behaviors at a capacity of 5 mAh cm2 and current density of 1 mA cm2. Fig. S11a shows the thickness variation of bare Li and bare Na during plating/stripping. Because of their different densities and theoretical capacities, when
homogenous without Na dendrites or obvious protuberances. This desirable surface morphology after cycling is ascribed to the 3D aerogel host with a homogeneous Na ion distribution on its surface, which resulted from the lower local current density and more uniform distribution of the ionic flux during Na plating/stripping than was the case for the bare Na electrode. As a contrast, Fig. 4e and f show the surface morphology of the bare Na anode after cycling under the same conditions as the Na@rGa composite anode. The bare Na anode possessed an uneven surface and mossy dendrites, which are attributed to the non-uniform local current density. This leads to uneven Na plating/stripping and high reaction kinetics on the bare Na surface. As expected, the bare Na foil surface has a rough surface with an unstable SEI layer and dendrite formation. These results further explain the improved cycling stability of the Na@rGa composite anode compared with that of the bare Na anode (Fig. 3). Fig. 4g and h present the EIS spectra of the Na@rGa composite and bare Na anode after 20 and 50 cycles of stripping/plating at a current density of 5 mA cm2 and capacity of 1 mAh cm2. The SEI and charge transfer resistance at the Na surface are associated with the high frequency semicircles in the Nyquist plots. As illustrated in Fig. 4g, the Nyquist plots for the Na@rGa composite anode showed no obvious increase of interfacial impedance after 50 cycles of N stripping/plating 380
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Fig. 4. (a, b) The schematic illustration of stripping and plating behavior of dendritefree Na@rGa composite anode and bare Na metal anode. (c–f) SEM images of Na@rGa and bare Na metal anode after 80 charge/ discharge cycles at 5 mA cm2 current density. (g, h) Nyquist plots of Na@rGa composite anode and bare Na after 20 cycles and 50 cycles at a current density of 5 mA cm2. (i) Top view (i) and cross-sectional (j) of the Na@rGa composite anode after plating 5 mAh cm2 at 1 mA cm2 (k) Top views of the Na@rGa composite anode after stripping 5 mAh cm2 at 1 mA cm2.
composite materials remained dense structure and the total thickness keeps at 320 μm after plating/stripping 5 mAh cm2 (Fig. S11d). These results confirm the Na@rGa composite anode can effectively inhibit volume expansion during plating/stripping. More important, Na@rGa composite anode also exhibits competitive high capacity compared with bare Na foil. As shown in Fig. 5a, a capacity of 1064 mAh g1 (based on the total weight of the electrode) was extracted when the Na@rGa composite anode was charged to 1 V. This capacity is very close to the theoretical capacity of bare Na (1165 mAh g1), suggesting that rGa host ensures stable cycling of composite anode and without loss of capacity. The high specific capacity can be attributed to the ultra-light weight and porous structure of rGa, which ensured high loading of Na metal in the whole composite anode (98.25 wt%) and large storage space. Therefore, aerogel hosts offer an exciting possibility of fabricating high-performance composite anodes with almost no adverse effect on capacity. In addition, we employed Na0.67Ni0.25Mn0.75O2 (NNM) as cathode and the Na@rGa as anode to assemble a full cell (NNM-Na@rGa) using 1 M NaClO4 in EC/DEC (1:1 in volume) as electrolyte. For comparison, the NNM cathode was also assembled with a bare Na anode as a reference (NNM-Na). Fig. 5b showed the galvanostatic charge-discharge profiles of NNM-Na@rGa and NNM-Na coin cells at 2 C rate for 300 cycle. NNMNa@rGa cell exhibits higher capacity compared with NNM-Na cell (62
plating/stripping the same areal capacity, the thickness variation of Na is larger than that of bare Li. Bare Na metal exhibits obvious rupture in a top-views SEM image (Fig. S11b) and thickness variation (about 42 μm) during plating at 5 mAh cm2, this behavior is roughly in accordance with previous calculation results, where volume variation caused by “hostless” nature of Na metal leads to rupture of the SEI and a large amount of “dead Na” during cycling (Fig. S11c). The “dead Na” hinders the contact between electrolyte and electrode and impede the ion transport rate, increase the overpotential, and cause serious fluctuation during cycling. As shown in Fig. 3i, the surface of the composite electrode is smoother and denser than the Na anode. The Na@rGa anode possesses uniform pores and a large specific surface area, which effectively lower the local current density and promote uniform Na nucleation. The surface morphology of the Na@rGa electrode is maintained well during cycling. Fig. 3j show the thickness variation of the Na@rGa anode during plating at a capacity of 5 mAh cm2. The thickness of the Na@rGa electrode increases by only 10 μm after plating. This is accordance with the results in Fig. 3i and suggests the Na@rGa composite anode can accommodate the volume change of Na electrode when the plating capacity is about 5 mAh cm2. After 1 stripping/plating cycle, the rGa structure is gradually exposed and shows obvious porous structure (Fig. 3k). Although the Na@rGa anode surface exhibited a different morphology compared with the initial stage, no obvious dendritic Na can be observed. Moreover, the 381
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Fig. 5. Electrochemical characterization of the Na@rGa electrodes: (a) The full Na stripping curve of the Na@rGa electrode to 1 V versus Naþ/Na at a current density of 25 mA g1, which shows a specific capacity of ~1064 mAh g1; (b) Galvanostatic discharging and charging curve of NNMNa@rGa anode at a rate of 2 C for 300 cycle. (c) Rate capability of the NNMNa@rGa cell at different rates from 0.2 to 10 C, and coulombic efficiency of NNMNa@rGa electrodes. (d) Cycling performance and coulombic efficiency of NNMNa@rGa electrode at a rate of 0.1 C.
mAh g1 vs 47 mAh g1). As can be seen from Fig. 5c, the rate capability of NNM-Na@rGa also presents reversible capacity of 73, 71, 68, 62, 51, 23 mAh g1 at the rates of 0.2, 0.5, 1, 2, 4, 10 C. Respectively. A capacity of 70 mAh g1 at 0.5 C and a high coulombic efficiency of 99.8 % are still obtained after 110 cycles. Meanwhile, NNM-Na@rGa cell exhibits more stable electrochemical performance over 100 cycles at the current density of 0.1 C with reversible capacity of 79 mAh g1 and coulombic efficiency of 99.5%. These results further illustrate that Na@rGa electrode still has a good adaptability in the full battery.
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4. Conclusion In summary, we successfully constructed Na@rGa composite anode via infusing molten Na metal into 3D graphene aerogel. Herein, we employed low-temperature hydrothermal to control the self-assembly of GO accurately. Because of its advantages of uniform porous structure, ultra-light quality and good wettability with both Na and electrolyte, rGa is an ideal choice as host for further application. Based on the above characteristics, the ultra-light host ensures a high energy density of Na@rGa composite anode, and the Na@rGa anode present excellent cycling performance in carbonate-based electrolyte without using any additives. At high current density of 5 mA cm2, the Na@rGa composite anode exhibited a highly cycling stability with no dendrite morphology and low hysteresis (50 mV) even after 1000 cycles which show an excellent electrochemical performance (Table S2). In addition, full-cell battery (NNM-Na@rGa) presents enhanced cycling stability. Our work provides new idea for the construction of 3D wettable scaffolds and principle for the facile applications of safe Na metal anodes. Meanwhile, it could be a viable strategy for other alkali metallic anode systems. Data availability The raw/processed data required to reproduce these findings cannot be shared at this time due to technical or time limitations. Acknowledgements This work was supported by the National Key Research and Development Program (2016YFB0901501), the National Natural Science Foundation of China (51772030). 382
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