Solid–Liquid Electrolyte as a Nanoion Modulator for Dendrite-Free

Jun 1, 2018 - Rechargeable lithium (Li) metal batteries are considered the most ..... PE/LE at temperatures ranging from 0 to 100 °C. The solid lines...
5 downloads 0 Views 5MB Size
Download PDF
Research Article www.acsami.org

Cite This: ACS Appl. Mater. Interfaces 2018, 10, 20412−20421

Solid−Liquid Electrolyte as a Nanoion Modulator for Dendrite-Free Lithium Anodes Kaihua Wen,†,‡ Yanlei Wang,† Shimou Chen,*,† Xi Wang,§ Suojiang Zhang,*,† and Lynden A. Archer∥

Downloaded via UNIV OF TOLEDO on June 21, 2018 at 14:29:28 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.

†

Beijing Key Laboratory of Ionic Liquids Clean Process, CAS Key Laboratory of Green Process and Engineering, Institute of Process Engineering, Chinese Academy of Sciences, Beijing 100190, P. R. China ‡ University of Chinese Academy of Sciences, Beijing 100049, P. R. China § School of Sciences, Beijing Jiaotong University, Beijing 100044, P. R. China ∥ Department of Materials Science and Engineering, Cornell University, Ithaca, New York 14850, United States S Supporting Information *

ABSTRACT: Rechargeable lithium (Li) metal batteries are considered the most promising of Li-based energy storage technologies. However, tree-like dendrite produced by irregular Li+ electrodeposition restricts it wide applications. Herein, based on a cation-microphase-regulation strategy, we create solid− liquid electrolytes (SLEs) by absorbing commercial liquid electrolytes into polyethylene glycol (PEG) engineered nanoporous Al2O3 ceramic membranes. By means of molecular dynamics simulations and comprehensive experiments, we show that Li ions are regulated and promoted in the two microphases, the channel phase and nonchannel phase, respectively. The channel phase can achieve homogeneous Li+ flux distribution by multiple mechanisms, including its uniform array of nanochannels and ability to suppress lateral dendrite growth by its high modulus. In the nonchannel phase, PEG chains swollen by electrolyte facilitate desolvation and fast conduction of Li+. As a result, the studied SLEs exhibit high ionic conductivity, low interfacial resistance, and the unique ability to stabilize deposition at the Li anode. By means of galvanostatic cycling studies in symmetric Li cells and Li/Li4Ti5O12 cells, we further show that the materials open a path to Li metal batteries with excellent cycling performance. KEYWORDS: solid−liquid electrolyte, nanochannel confinement, nanoion modulator, stable electrodeposit, lithium metal batteries

■

termed dendrite growth.8 Thickening of the solid electrolyte interface (SEI) on the Li anode and consumption of electrolyte lead to low Coulombic efficiencies and cycling instability that may end in cell failure by either voltage or thermal run-away. Researchers have made a lot of efforts to stabilize uniform Li electrodeposition on metallic anodes. Designing novel electrolytes and Li−electrolyte interfaces are the two main strategies.9−14 Solid electrolytes, such as poly(ethylene oxide) (PEO) based solid polymer electrolytes (SPEs), and solid inorganic electrolytes (SIEs) with high mechanical modulus have both emerged as important platforms for achieving enhanced stability.15−18 However, low ionic conductivities and high interfacial resistance at room temperature currently pose serious barriers for large-scale application.19−21 Methods that take advantage of the beneficial aspects of SPEs and SIEs, but which utilize inorganic fillers and plasticizers, can increase ionic

INTRODUCTION Rechargeable Li ion batteries (LIBs) play important roles in energy storage and conversion technologies.1−3 However, limited by the low storage capacity of the carbon anode (e.g., 360 mAh g−1 for LiC6), the specific energies of current LIBs are at best modest (100−265 Wh kg−1). This makes it difficult to meet the demands of electric vehicles and large-capacity electrical energy storage systems.4 To pursue higher specific energies, a new generation of battery technologies based on Li metal anodes, including Li−S and Li−O2 cells, has been widely studied. Such cells offer specific energies of 2567 Wh kg−1 and 3505 Wh kg−1, respectively.4−6 It is straightforward to show that these high specific energies derive from the high theoretical specific capacity (3860 mA h g−1), low density (0.59 g cm−3), and lowest negative electrochemical potential (−3.04 V vs the standard hydrogen electrode) of the Li metal.7,8 Unfortunately, repeated Li deposition/dissolution lead to parasitic reactions between Li and electrolyte components as well as the volume expansion of the Li anode, consequently resulting in irregular large volume changes, which drive the morphological instability © 2018 American Chemical Society

Received: February 28, 2018 Accepted: June 1, 2018 Published: June 1, 2018 20412

DOI: 10.1021/acsami.8b03391 ACS Appl. Mater. Interfaces 2018, 10, 20412−20421

Research Article

ACS Applied Materials & Interfaces

Figure 1. (a) Synthesis procedure for the solid electrolyte films. Inset: enlarged image of PDA-g-PEG modified layer on AAO film. (b) SEM images of surfaces of the AAO(200) films before (left) and after (right) modification. Inset: digital images of AAO (200 nm) film and A@PP-5 film. (c) SEM images of the cross sections of the AAO(200) films before (left) and after (right) modification. (d) XPS study of the surface of A@PP-5 film. (e) Contact angle image of A@PP-5 film.

membranes with well-aligned nanochannels with pore sizes in the range 20−400 nm. Besides, PEG chains in the nonchannel phase facilitate the ionic transportation on the electrode− electrolyte interface.21 This SLE enables homogeneous Li+ distribution and rapid ion transport in its two contiguous nanophases. The Li metal battery based on this SLE exhibits superior electrochemical performance, which indicates the great possibility of using this SLE with a mechanism of nanoion modulator for Li metal based energy storage system.

conductivity, but even in state-of-the art materials it is difficult to achieve the combination of high bulk and interfacial conductivities that defines the liquid electrolytes (LEs) used in LIB technology.22−26 Additionally, additives for LEs such as LiF,27 Cs+,28 In3+,29 ionic liquids,30−32 nanodiamonds,33 and novel Li salt34 have been reported to facilitate improved ion transport and stability of the SEI by various fundamental processes. Nanoscale Li anodes and current collectors with high specific surface areas have likewise been explored for their potential to reduce local current densities and provide active sites with desired topologies for Li electrodepositon (i.e., 3D Cu/Li composite anode,35 vertically porous Cu electrode,36 polyimide modified Li anode,37 glassy fiber modified Cu electrode,38,39 N-doped graphene/Li anode,40 and microcompartmented anode arrays,41 etc.). The mechanism of the formation of Li nucleation and deposition has been revealed in recent studies.42 Results from both experiments and simulation show that nanoscale Li anode can regulate the electrodeposition of Li+ by uniform Li-ion flux.37,42 Herein, we propose a novel SLE as a nanoion modulator (NIM) that allows facile regulation of Li-ion transport for stable Li deposition. Nanoporous anodic aluminum oxide (AAO) films form the basis of our design. Previously, it is reported that the AAO with sufficiently small pore sizes (e.g., 20, 100 nm) can effectively suppress the nucleation and deposition of Li+ on the Li surface.43−46 A transition from stable to unstable deposition was predicted when Li nuclei size exceeded a value around 200 nm.9,44 In our work, we focus on the AAO

■

RESULTS AND DISCUSSION Figure 1a summarizes the synthesis procedure for the SLE films. After a series of polarization and grafting, a uniform PEGmodified AAO film was prepared.47 Figure 1b shows the surface morphology of AAO films before and after modification by PEG (∼0.5 μm thickness, Figure S1). XPS was used to characterize the surface chemistry of the films, with a scanning range from 0 to 1350 eV and high-resolution scanning near carbon peaks. Due to the surface charging, the raw binding energies were corrected for a constant downshift of 1.62 to 2.0 eV using an internal reference. The spectrum peaks are located in 285.0 eV, 400.83 eV, 532.10 eV, and 74.24 eV corresponding to C 1s, N 1s, O 1s, and Al 2p, respectively. Tables S1 and S2 show surface chemical composition of the unmodified AAO membrane and PDA-coated AAO before and after PEG immobilization. The data were obtained from peak intensities of Figures 1d and S2. The N 1s content of AAO membranes 20413

DOI: 10.1021/acsami.8b03391 ACS Appl. Mater. Interfaces 2018, 10, 20412−20421

Research Article

ACS Applied Materials & Interfaces

Figure 2. (a) Schematic of the Li matal battery of Li|A@PP-5/LE|LTO. The green curves - PEG chains; blue spheres - Li + cation; tawny layer on AAO - PDA layer. (b) Snapshots of the simulation box from MD simulation. (c) The diffusive coefficient D of Li+ in confined space from MD simulations and theoretical model. The red line is the theoretical prediction based on the MD results. (d) The interaction distance between PEG and solvents, which shows that PCs have a close contact with PEG compared to Li+ and TFSI−. (e) Mechanism of Li+ migration in different models. Note: for simplicity, the above double-sided modified AAO represents only the right side.

with a certain range ensure stable Li nucleation and regular deposition, which achieves a homogeneous electrolyte− electrode interface with thin and stable SEI layer. (3) The PEG nanolayer can efficiently reduce the interfacial impendence against Li anode, owning to the desolvation of PEG segments.48,49 We analyzed the transport of Li+ in the two phases using molecular dynamics (MD) simulations (Figure 2a,b). In the channel phase (CP), Li+ flux can be uniformly distributed in well-aligned nanochannels, facilitating homogeneous Li deposition and stable SEI layers. In the simulations, we compare the diffusion coefficients in different pore sizes, from 1 to 15 nm as shown in Figure 2c. We can see that the diffusion coefficient increases with the pore diameter and is always smaller than that in the bulk case. The diffusive coefficient for Li+ in the bulk case is D = 0.966 × 10−11 m2s−1, which is on the same order of LiTFSI as in other solvents from the prior studies,50−54 showing the rationality of our methods. To extend our results from the simulation to the large-scale system, we construct a two-phase model as shown in Figure S4. We divided our CP into two regions, edge layer (∼0.45 nm thickness) and center region. For 1 nm channels, the Li+ diffusion suffered the most from electrostatic force. As the pore sizes increased, the ratios of edge layer decreased, and the diffusion coefficients are observed to increase. Based on the above simulation, an analytical relationship between the diffusion coefficients and

before and after coating PDA is increased from 0.39% to 9.38%, with a ratio of N/C = 0.138 (close to the theoretical value of N/C = 0.125 of dopamine), which can further demonstrate the existence of PDA-coated layers. The PEG-coated layer is evaluated by the content of C−O bond, and 35% of C 1s is cooperated with O 1s to form a C−O bond, which is much more than the pristine AAOs (1.28%) and PDA@AAOs (15.70%). As shown in Figure 1c, cross sections of SLE films basically have the same pore sizes (∼200 nm) as before. Significantly, the grafted PEG is highly surface selective, and the pores are not blocked by PEG. This configuration ensures fast Li+ diffusion in the nanopores with certain apertures. As shown in Figures 1e and S3, the SLE film with a small contact angle of 28.5° can easily absorb LEs. To use the SLE films in Li batteries, enough liquid electrolytes must be loaded inside the nanoscale pores. The SLE films were immersed in LiTFSI/PC for at least 6 h, and the uptake values are ∼60 wt %, which is much higher than the value of polyethylene (PE) separator (48.9 wt %). The liquid electrolyte uptake amount depends on the pore size. A rough calculation shows that a smaller pore AAO membrane exhibits higher porosity (Table S3). We hypothesize that the as-prepared SLEs will enable more stable Li deposition and lower interfacial resistance (Rint) than batteries that employ generic polyolefin separators for at least three reasons: (1) The γ-Al2O3 with high shear modulus can effectively suppress uneven Li dendrites.9 (2) The nanopores 20414

DOI: 10.1021/acsami.8b03391 ACS Appl. Mater. Interfaces 2018, 10, 20412−20421

Research Article

ACS Applied Materials & Interfaces

Figure 3. (a) Temperature-dependent ionic conductivity of A@PP-5/LE, compared with PE/LE, ranging from 0 to 100 °C. (b) Current variation with time during polarization of a Li|A@PP-5/LE|Li symmetrical cell at room temperature. Inset shows the AC impedance spectra of a symmetrical lithium battery. (c) Cyclic voltammetry and linear sweep voltammetry of a Li|A@PP-5/LE|stainless steel cell at room temperature. (d) Impedance spectra of A@PP-5/LE against temperature ranging from 0 to 40 °C. Inset shows impedance spectra of A@PP-5/LE against temperatures ranging from 50 to 100 °C. (e) Resultant Rb and Rint of the Li|A@PP-5/LEs|Li cell fitted by the equivalent circuit model. (f) Time dependence of the interfacial resistance of Li|PE/LEs|Li (left) and Li|A@PP-5/LEs|Li (right) symmetrical cells at room temperature.

AAO/LE system, Li ions are modulated by the well-aligned nanochannels along the normal direction of electrodes, which reduces the influence of “tip effect” on a rough Li surface. Combining its high shear modulus, AAO can achieve modulated Li nucleation and suppress large dendrites. Moreover, to reduce the space charge layer on the Li surface and enhance ion transport on the electrolyte−electrode interface, a PEG-modified AAO is introduced. In our nanoion modulator, the A@PP-5/LE system, the PEG chains get well entangled with PC and TFSI− by the complexing ability of oxygen atoms, which release Li+ from large solvated molecules and eliminate the space charge layer. This configuration facilitates the application of LMBs with stable Li deposition and long lifetime performance. Figure 3a illustrates the conductivity of A@PP-5/LE in comparison to PE/LE at temperatures ranging from 0 to 100 °C. The solid lines represent the Vogel−Tamman−Fulcher (VTF) equation fits of the measured temperature-dependent conductivity data and the activation energy deduced from the fits. Ionic conductivities of PEG-modified AAOs with different thickness (A@PP-2/LE, A@PP-5/LE, A@PP-10/LE, and A@ PP-20/LE) are shown in Figure S8 and Table S3. A@PP-5/LE possesses the highest ionic conductivity (2.85 mS cm−1 at 25 °C). It should be noted that ionic conductivity of AAO200/LE was 2.63 mS cm−1 at 25 °C, which was 1 order of magnitude larger than that of PE/LE (0.66 mS cm−1). As shown in Figure 3b, the tLi+ value of the Li|A@PP-5/LE|Li symmetrical cell was polarized from an initial current value of 29.7 μA to a stable current value of 26.5 μA at a voltage of 3 mV. The interfacial resistance changed from 105.1 Ω to 118.5 Ω after the polarization process. The Li-ion transference number can be calculated by eq 4 to be tLi+ = 0.83, which is much higher than

pore sizes can be proposed (Figure 2c). Results reported in Figure S5 show that the ionic conductivity follows a similar relation. Conversely, when the size of the channel increases, the growth of Li dendrites could not be suppressed. Taken together, these results suggested that channels, with a pore size between 145 and 250 nm, could be considerable candidates (Figure S6). To investigate the interfacial structure of lithium salt and PEG molecules, we also constructed a model consisting of 20 chains of PEG (20 monomers per chain), 43 Li+, 43 TFSI−, and 510 PC molecules. In the nonchannel phase (NCP), ether bonds on PEG chains may associate with Li+, TFSI−, and PC, resulting in a unique solvation structure at the interphase. The radii of structures can be described by the distances between PEG and Li+, TFSI−, or PC. As shown in Figure 2d, Li+ possessed larger coordination distances with both PEG and oxygen atoms in PEG than those of TFSI− and PC. These results indicate that TFSI− and PC form stronger solvation structures with PEG chains than Li+. That is, the intrinsic solvated structures of Li+ in the bulk phase can be broken, which can enhance Li+ transport at the electrode−electrolyte interphase. Additional simulations and parameters are available in Figure S7. According to these analyses, we propose a modulation mechanism of Li + migration that can fundamentally regulate the Li deposition morphology (Figure 2e). In the PE/LE system, random oriented and solvated Li+ are easily converged on the Li surface, due to the “tip effect”.8,55 The space charge layer, formed by the consuming of TFSI−, hinders Li ionic transport between electrolyte and metallic Li electrode.56 Thus, high local current densities promote spare Li+ nucleation toward tips, which finally results in large dendrites. In the 20415

DOI: 10.1021/acsami.8b03391 ACS Appl. Mater. Interfaces 2018, 10, 20412−20421

Research Article

ACS Applied Materials & Interfaces

Figure 4. Li nuclei deposited in different systems. Ex situ SEM images of Li deposition in (a) PE/LE, (b) AAO/LE, and (c) A@PP-5/LE at current densities of 0.5 and 1 mA cm−2, respectively, for a total areal capacity of 0.1 mA h cm−2. Surface SEM images of Li electrodes after galvanostatic cycling at current density of 0.5 and 1 mA cm−2 in (d) PE/LE, (e) AAO/LE, and (f) A@PP-5/LE at current densities of 0.5 and 1 mA cm−2, respectively, for 120 h. (g) Surface SEM image of pristine Li. (h) Schematic of the symmetric cell for the lithium plating/stripping experiment. Nyquist plots of symmetrical lithium cells with a stripping/plating process at a current density of (i) 0.5 mA cm−2 and (j) 1 mA cm−2 before and after 20 cycles.

that measured in a PE separator (Figure S9, tLi+ = 0.39). It is also higher than typical values (tLi+ < 0.5) reported for PEObased polymer electrolytes.57 We conclude that the high tLi+ of A@PP-5/LE makes it possible to be applied in a fast charging/ discharging system. The electrochemical stability was evaluated by cyclic voltammetry and linear sweep voltammetry with Li| A@PP-5/LE|SS cell at room temperature. As shown in Figure 3c, A@PP-5/LE is stable within 4.5 V vs Li+/Li. Figure 3d shows electrochemical impedance spectra (EIS) of A@PP-5/ LE evaluated in a Li|Li cell as a function of temperature. The results were fitted by the equivalent circuit model depicted in Figure S10, and the resultant Rb and Rint are provided in Figure 3e. It seems that the Rint is the main hurdle to deposition. The Rint is related to the passive layer and charge transfer on Li anode. AAO200 shows the lowest Rint (529 Ω cm−2 at 20 °C) among various pore diameters (Figure S11) but is still larger than the PE separator (171 Ω cm−2 at 20 °C, Figure S12). Figure S13 shows the Rb and Rint for AAO200 films with different thickness of the PEG layers. The Rint values for A@PP2/LE, A@PP-5/LE, A@PP-10/LE, and A@PP-20/LE are 139, 134, 157, and 170 Ω cm−2 at 20 °C, respectively. Therefore, A@PP-5/LE is more suitable to be the candidate for further study. The interfacial stability of PE/LE and A@PP-5/LE against Li anodes was studied by EIS of Li|PE/LE|Li and Li|A@ PP-5/LE|Li cells for a period of 30 days. As shown in Figure 3f, the Rint of A@PP-5/LE slightly increases for the first 5 days,

which is attributed to the chemical reaction between LE and the Li anode. On longer time scales, Rint is observed to decrease slightly and finally to approach a constant value for the last 20 days, which verified that A@PP-5/LE makes a stable SEI layer on the Li surface. In contrast, the PE/LE shows an uncertain change with much larger Rint in the same period than that of A@PP-5/LE. To investigate the evolution of Li nucleation and deposition, anodes were observed after a fixed amount of Li (0.1 mAh cm−2) and after galvanostatic cycling for 120 h, respectively. Figure 4a shows the appearance of large, tree-like dendrites in the PE/LE system. In the AAO/LE system, Li deposits as relatively smooth particles (Figure 4b). In the NIM (Figure 4c), uniform spherical nuclei of Li with an average size of 350 nm are plated on the Li anode. These nuclei are much larger than the pores of nanchannels (∼200 nm). To study long-time Li deposition, EIS was carried out to characterize the interfacial stability in different cells before and after cycling (Figure 4i, 4j). Before cycling, the Li|PE/LE|Li cells show large interfacial resistances of about 179 Ω cm−2 (0.5 mA cm−2) and 171 Ω cm−2 (1 mA cm−2), which could be attributed to the formation of thick and random passivation films on Li anodes. The values of Rint drop to 59 Ω cm−2 (0.5 mA cm−2) and 92 Ω cm−2 (1 mA cm−2) after 20 cycles. The Li|AAO/LE|Li cells show interfacial resistances of about 179 Ω cm−2 (0.5 mA cm−2) and 168 Ω cm−2 (1 mA cm−2) and dropped to 65 Ω cm−2 (0.5 mA 20416

DOI: 10.1021/acsami.8b03391 ACS Appl. Mater. Interfaces 2018, 10, 20412−20421

Research Article

ACS Applied Materials & Interfaces cm−2) and 88 Ω cm−2 (1 mA cm−2) after 20 cycles, indicating that AAO/LE can make homogeneous Li deposition and SEI layers by its well-aligned nanochannels. In comparison, the Li| A@PP-5/LE|Li cells show relatively lower interfacial resistances of about 132 Ω cm−2 (0.5 mA cm−2) and 138 Ω cm−2 (1 mA cm−2), and then the values of Rint drop to 54 Ω cm−2 (0.5 mA cm−2) and 50 Ω cm−2 (1 mA cm−2) after 20 cycles. The results clearly reveal that cells with NIM can reduce space charge layer on the electrode and accelerate the interfacial transfer of Li+. After running the symmetric cells for 20 cycles (120 h), we disassembled some of the cells in a glovebox. As shown in Figure 4g and Figure S14a, the surface and side section of pristine Li electrodes are uniform in general except for small defects. For Li|PE/LE|Li cells, after 20 cycles, the surface becomes rough with large cracks and thick SEI films on the Li surface (Figure 4d and Figure S14b). For Li|AAO/LE|Li cells, after 20 cycles, the surface shows similar features (Figure 4e and Figure S14c). By contrast, the anodes using A@PP-5/LE present a smooth and flat surface with thin SEI corresponding to efficiency Li plating and striping (Figure 4f and Figure S14d). As a result, Li dendrites are inhibited with the structure of NIM, and the interphase with reduced space charge layer is provided by PEG. To further evaluate the electrochemical compatibility of A@ PP-5/LE against the Li anode, Li stripping/plating experiments were performed in symmetric cells. Figure 5a compares the voltage profiles obtained in Li|A@PP-5/LE|Li and Li|PE/LE|Li cells at a fixed current density of 0.5 mA cm−2. It shows that the Li|PE/LE|Li cells exhibit a sudden drop after approximately 400

h, which is attributed to short circuits induced by dendrites. On the contrary, the cycling performance of Li|A@PP-5/LE|Li reaches a low and steady-state overpotential for over 700 h under the same current density. At 1 mA cm−2 (Figure 5b), stable cycling can be achieved by A@PP-5/LE, while PE/LE becomes short-circuited at around 200 h. Most notably, Li|A@ PP-5/LE|Li can achieve long-time plating/stripping at 0.05 mA cm−2 with polarization as a limit at 22 mV for more than 6000 h (Figure 5c). To test the practical performance of SLEs, we studied Li| Li4Ti5O12 (LTO) cells at 60 °C. Figure 6a,b presents the capacity as a function of cycle number at current densities of 0.175 A g−1 (1 C) and 0.875 A g−1 (5 C), respectively. It is obvious that the cells with NIM exhibit stable, high efficiency cycling over 500 and 1000 charge−discharge cycles at 1 and 5 C, with only minimal capacity fading over the first few cycles. The specific capacity of Li|A@PP-5/LE|LTO at 1 and 5 C is about 165 mAh g−1 (94% of theoretical capacity) (Figure 6d) and 135 mAh g−1 (77% of theoretical capacity) (Figure 6e), respectively. However, the Li|PE/LE|LTO cells exhibit significant capacity fading in the process of charging and discharging at 5 C, which is attributed to weight loss of PE/LE at high temperature (Figure S15). To further emphasize the significance of the SLEs, we evaluate the rate behaviors of the Li/LTO batteries at different rates at 60 °C (Figure 6c). The cell of Li|A@PP-5/LE|LTO shows enhanced discharge capacity, especially at high rates. It displays a capacity of 131 mAh g−1 at 10 C and 103 mAh g−1 at 20 C, whereas the capacities for cells with PE/LE cannot stand such high currents. A breif comparison of electrochemical performances were performed in Table S4 with various systems of Li metal batteries.58−60 Overall, our SLEs exhibited superior electrochemical performances, such as liquid-like ionic conductivity, low Rint, and longlife cell operations. In addition, the long-term cycling and rate behaviors at room temperature are shown in Figures S16 and S17. This system can also be applied in Li|LiFePO4 battery (Figure S18). These results again show the potential importance of SLEs for practical application.

■

CONCLUSIONS In summary, we propose a SLE as NIM for stable LMBs. Our SLE shows high ionic conductivity, low Rint, high mechanical modulus and good stability against the Li anode at both room and elevated temperatures. The materials were assembled as Li/ LTO cells to investigate their longtime cycling performance. These measurments show that more than 1000 charge/ discharge cycles can be achieved with no evidence of dendritic deposition. The Li|A@PP-5/LE|Li cells exhibit more than 6000 h of stable cycling at current densities ranging from 0.05 to 1 mA cm−2. The results indicated that the model of NIM can make homogeneous Li + distribution and rapid transportation in its two special nanophases. This configuration facilitates the application of LMBs with stable Li deposition and long lifetime performance. We believe that the concept of NIM is simple and easy to be used in next-generation energy storage systems for advanced stable battery technologies.

■

Figure 5. Galvanostatic cycling performance of Li|A@PP-5/LE|Li (red) and Li|PE/LE|Li (black) cells at a fixed current density of (a) 0.5 mA cm−2 and (b) 1 mA cm−2 at room temperature. (c) Galvanostatic cycling performance of Li|A@PP-5/LE|Li cells at a fixed current density of 0.05 mA cm−2 (red) and 0.1 mA cm−2 (blue) at room temperature. Each charge and discharge time is set as 3 h.

EXPERIMENTAL SECTION

Molecular Dynamics Simulations. All the molecular dynamics (MD) simulations in this work were completed using a large-scale atomic/molecular massively parallel simulator (LAMMPS).61 The time step for integrating Newtonian equations of motion was 0.5 fs. To simulate the LiTFSI + PC liquids in the nanochannels, we built a 20417

DOI: 10.1021/acsami.8b03391 ACS Appl. Mater. Interfaces 2018, 10, 20412−20421

Research Article

ACS Applied Materials & Interfaces

Figure 6. Galvanostatic cycling of Li/LTO batteries with A@PP-5/LE (red), A200/LE (blue), and PE/LE (black) at 1 C (a) and 5 C (b). (c) Rate performances of Li/LTO batteries with A@PP-5/LE (red) and PE/LE (black) at 60 °C. The voltage profiles of Li/LTO battery with A@PP-5/LE at 1 C (d), 5 C (e), and different C rates (f) at 60 °C corresponding to (a), (b), and (c), respectively. Preparation of the Hybrid Electrolyte. The polymerization and grafting reactions reacted in a buffur solution. Amounts of 0.2422 g of Tris base (supplied by NOVON) and 60 μL of hydrochloric acid (HCl, supplied by Beijing Chemical Works) were put into 200 mL of distilled water in a beaker under constant stirring to get a uniform buffer with pH = 8.5. Dopamine hydrochloride (DA·HCl, supplied by Aladdin) was dissolved in the above buffer with a concentration of 2 mg mL−1. Nanoporous anodic aluminum oxide membranes (AAO, Whatman Anodisc 25 with pore sizes of 200 nm and a thickness of 60 μm) were immersed into the DA solutions, and the beakers were continuously shaken for 24 h on a shaking table at room temperature. After that, the PDA-modified AAOs (A@P, 0.500 mg cm−2 PDA on AAO) were rinsed by distilled water and dried at 80 °C. The prepared A@P films were immersed into methoxypolyethylene glycol amine (mPEG-NH2, M.W. 5000, supplied by Aladdin) solutions with certain concentrations (2, 5, 10, and 20 mg mL−1) in the initial buffer solution to get PDA-g-PEG-modified AAOs with different thickness of the PEG layers, hereafter referred to as A@PP-2, A@PP-5, A@PP-10, and A@ PP-20. After 24 h saking, the films were taken out, washed with distilled water and absolute ethyl alcohol for 3 times, respectively, and then dried at 80 °C under vacuum for 12 h. Finally, the dried pieces were transformed into a glovebox and immersed into 1 M LiTFSI/PC (supplied by Shanghai Xiaoyuan Energy Technology Co., Ltd.) for more than 24 h to obtain the SLE/LE for further use. Electrode Preparation and Cell Assembly. The slurry consisting of 80 wt % Li4Ti5O12, 10 wt % Super P carbon, and 10 wt % PVdF was casted on copper foil and dried in a vacuum oven at 80 °C for 24 h. The foil was punched into slices used as cathode, and Li metal foils were used as anode. A half-cell was assembled by sandwiching the SLE/LEs between a Li4Ti5O12 cathode and a Li anode in a CR2025 coin cell. A Li/LiFePO4 half cell was assembled the same way as mentioned above. The areal loading of dried Li4Ti5O12 and LiFePO4 was 2.61 mg cm−2 and 3.12 mg cm−2, respectively. Characterization and Performance Evaluation. The surface morphology of the SLE was examined by using an Ultrahigh Resolution Scanning Electron Microscope (SEM, Haitich SU8020, Japan) with an acceleration voltage of 5 kV before immersing in liquid

model as shown in Figure S4, where the wall is Al2O3 representing the AAO nanochannels in the experiments. The cross-section of the system was 4.2 × 3.9 nm2, and a series distance (d) of nanochannel was considered, for example, w = 0.5, 1, 2.5, 5, 7.5, 10, and 15 nm. In all the simulations, the LiTFSI solution with a concentration of 1.0 M lithium salt was considered, and for the case w = 10 nm, the system consisted of 1092 PC molecules, 92 Li+, and 92 TFSI−. The periodic boundary conditions (PBCs) were used in the y and z directions, while an open boundary condition was adopted along the x direction. The BKS potential was used to describe the atomic potential of Al2O3, and the OPLS-AA potential was employed for the LiTFSI and PC molecules.62,63 The interactions between Al2O3, Li+, TFSI−, and PC include two parts: van der Waals interactions and electrostatic terms. The former one was described using the Lennard-Jones potential function, and the Lorentz−Berthelot mixing rules were used to model the parameters, which are truncated at 1.2 nm. The later one, longrange Columbic interactions, was computed using the particle− particle−particel-mesh (PPPM) algorithm.64 The system was first relaxed at a temperature of 300 K and 1.0 atm using a Berendsen thermostat for 5 ns. Then the system is relaxed in the NVE ensemble for 2 ns to up to a equilibrated state. After the equilibrium was archived, the additional 1 ns was run to analyze the structure and property of lithium salt. The self-diffusion coefficient D of Li+ can be computed from the mean square displacement (MSD) of atomic trajectories based on the Einstein relation.65 The mass density and radial distribution function (RDF) were also calculated based on the atomic trajectories from MD simulations. To investigate the interfacial structure of lithium salt and PEG molecules, we also constructed a model consisting of 20 chains of PEG (20 monomers per chain), 43 Li+, 43 TFSI−, and 510 PC molecules. The CVFF force field was used to describe the atomic structure of PEG.66 The interactions between PEG and lithium salt mainly consisted of two parts: van der Waals interaction and electrostatic force. The system was first equilibrated in NPT ensemble at a temperature of 300 K and 1.0 atm using a Berendsen thermostat for 5 ns. Then more 1 ns simulations were performed to capture the interfacial structure of lithium salt−PEG. 20418

DOI: 10.1021/acsami.8b03391 ACS Appl. Mater. Interfaces 2018, 10, 20412−20421

ACS Applied Materials & Interfaces

■

electrolyte. To facilitate realistic testing, the pieces were sputter coated with gold. To further verify the chemical composition of the SLE films, an X-ray photoelectron spectroscopy (XPS) study was carried out. The surface of Li metal anode after cycling was also studied by SEM, with a special vacuum transfer box to avoid Li oxidation. The films were immersed into liquid electrolyte for 24 h. The electrolyte uptake (EU) was determined by the weight change before and after immersion and can be calculated as the following

EU (%) = [(W2 − W1)/W1] × 100%

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.8b03391. SEM images of cross-sections of modified films and Li anodes, XPS results of AAO, A@P, and A@PP-x films, contact angles of PE and AAOs, theoretical model for channel and nonchannel phases, relative short-circuit time and relative ionic conductivity, ionic conductivities of AAO/LEs with different pore sizes, tLi+ of PE/LE, EIS equivalent circuit model, Rb and Rint for different systems, TGA results of PE/LE and A@PP-5/LE, and cycling and rate performances at room temperature (PDF)

(1)

where W1 and W2 are the weights of the dry films, such as AAO, A@P, A@PP-x (x stands for the concentration of mPEG-NH2), and PE, before and after immersion in liquid electrolyte. Extra liquid electrolytes were wiped by kimwipes. The wettability of SLE films was tested with a contact angle machine. The thermal properties of PE/LE and A@PP-x/LE were examined by a thermogravimetric analyzer (Setaram Labsys) with a heating rate of 5 °C min−1. The ionic conductivity of the SLEs was determined by measuring electrochemical impedance spectra (EIS) on the electrochemical station Metrohm Autolab (PGSTAT302N) in the frequency range of 1 Hz to 1.0 MHz with 10 mV of AC amplitude in the temperature range from 0 to 100 °C. The samples for the measurements were prepared by sandwiching the SLEs between two stainless steel disc electrodes (Φ = 10 mm). The ionic conductivity was calculated by the following

σ = (1/R ) × (d /S)

Research Article

■

AUTHOR INFORMATION

Corresponding Authors

*Shimou Chen. E-mail: [email protected]. *Suojiang Zhang. E-mail: [email protected]. ORCID

Shimou Chen: 0000-0002-2533-4010 Xi Wang: 0000-0003-3910-9575 Suojiang Zhang: 0000-0002-9397-954X Lynden A. Archer: 0000-0001-9032-2772

(2)

Notes

The authors declare no competing financial interest.

where R is bulk resistance of SLEs obtained from AC impedance spectrum; d is the thickness of electrolytes; and S is the effective area between electrolyte and electrode. The effective activation energy barrier of the electrolyte was fitted by experimental results according to the VTF equation σ = AT −1/2 exp[− E0 /R(T − T0)]

■

ACKNOWLEDGMENTS This work was finacially supported by National Key Projects for Fundamental Research and Development of China (No. 2016YFB0100104), National Natural Science Foundation of China (Nos. 91534109 and 91434203), the Fund of State Key Laboratory of Multiphase Complex Systems (MPCS-2017-A08), Beijing Municipal Science and Technology Project (D171100005617001), and Henan province science and technology cooperation project (172106000061).

(3)

where A is a constant proportional to the number carrier ions; E0 is the pseudoactivation energy barrier; T is the measurement temperature; and T0 is the ideal glass transition temperature. The Li-ion transference number (tLi+) was obtained using a symmetric cell of Li|SLE/LE|Li by the direct current (DC) polarization combined with EIS method, and it can be calculated according to

t Li + = Is(ΔV − R 0I0)/[I0(ΔV − R sIs)]

■

REFERENCES

(1) Xu, K. Electrolytes and Interphases in Li-Ion Batteries and Beyond. Chem. Rev. 2014, 114, 11503−11618. (2) Sun, C.; Liu, J.; Gong, Y.; Wilkinson, D. P.; Zhang, J. Recent Advances in All-Solid-State Rechargeable Lithium Batteries. Nano Energy 2017, 33, 363−386. (3) Li, Q.; Chen, J.; Fan, L.; Kong, X.; Lu, Y. Progress in Electrolytes for Rechargeable Li-Based Batteries and Beyond. Green Energy & Environment 2016, 1, 18−42. (4) Bruce, P. G.; Freunberger, S. A.; Hardwick, L. J.; Tarascon, J. M. Li−O2 and Li−S Batteries with High Energy Storage. Nat. Mater. 2012, 11, 19−29. (5) Zhang, S.; Ueno, K.; Dokko, K.; Watanabe, M. Recent Advances in Electrolytes for Lithium−Sulfur Batteries. Adv. Energy Mater. 2015, 5, 1500117. (6) Grande, L.; Paillard, E.; Hassoun, J.; Park, J.-B.; Lee, Y.-J.; Sun, Y.K.; Passerini, S.; Scrosati, B. The Lithium/Air Battery: Still an Emerging System or a Practical Reality? Adv. Mater. 2015, 27, 784− 800. (7) Goodenough, J. B.; Park, K. S. The Li-Ion Rechargeable Battery: A Perspective. J. Am. Chem. Soc. 2013, 135, 1167−1176. (8) Cheng, X. B.; Zhang, R.; Zhao, C. Z.; Zhang, Q. Toward Safe Lithium Metal Anode in Rechargeable Batteries: A Review. Chem. Rev. 2017, 117, 10403−10473. (9) Tikekar, M. D.; Choudhury, S.; Tu, Z.; Archer, L. A. Design Principles for Electrolytes and Interfaces for Stable Lithium-Metal Batteries. Nat. Energy 2016, 1, 16114.

(4)

where I0 and Is are the initial and steady-state DC current, respectively; R0 and Rs are the interfacial resistances of the initial and steady state; and ΔV is the applied potential. Electrochemical stability of SLE was estimated by linear sweep voltammetry (LSV) and cyclic voltammetry (CV) using a three-electrode cell Li|SLE/LE|SS at a scan rate of 1 mV s−1 at room temperature on Autolab. Stainless-steel disc acts as the working electrode and Li metal as the reference electrode and counter. The cells were swept in potential ranges from 2 to 5 V and −1 to 2 V, respectively. The compatibility of SLEs/LE with Li anode was measured on Autolab by EIS using a symmetrical cell Li|SLEs/LE|Li with 10 mV of AC amplitude in the frequency range of 0.1 Hz to 100 kHz. This measurement was studied both at different temperatures (0 to 100 °C) and at different storage time (1 day to 30 days). The Li planting/striping experiment was performed under different current densities (0.1, 0.5, 1 mA cm−2). The cell was initially charged at a fixed current density for 90 min and then discharged at the same current density for 180 min, followed with charging 180 min to continue the cycling. For the galvanostatic cycling experiment, the assembled half cells were performed on a Neware CT-3008 battery tester under different charging/discharging rates (1 and 5 C) between 1.0 and 3.0 V for 500 or 1000 times. The C-rate capability was conducted at the rates of 0.2 C, 0.5 C, 1 C, 2 C, 5 C, 10 C, and 20 C and then 0.5 C and 1 C between 1.0 and 3.0 V for 10 cycles for each rate. 20419

DOI: 10.1021/acsami.8b03391 ACS Appl. Mater. Interfaces 2018, 10, 20412−20421

Research Article

ACS Applied Materials & Interfaces

Lithium Anodes for Fast Interfacial Ion Transport. Angew. Chem., Int. Ed. 2017, 56, 13070−13077. (30) Li, N.-W.; Yin, Y.-X.; Li, J.-Y.; Zhang, C.-H.; Guo, Y.-G. Passivation of Lithium Metal Anode via Hybrid Ionic Liquid Electrolyte toward Stable Li Plating/Stripping. Adv.Sci. 2017, 4, 1600400. (31) Lu, Y.; Korf, K.; Kambe, Y.; Tu, Z.; Archer, L. A. Ionic-Liquid− Nanoparticle Hybrid Electrolytes: Applications in Lithium Metal Batteries. Angew. Chem., Int. Ed. 2014, 53, 488−492. (32) Nugent, J. L.; Moganty, S. S.; Archer, L. A. Nanoscale Organic Hybrid Electrolytes. Adv. Mater. 2010, 22, 3677−3680. (33) Cheng, X.-B.; Zhao, M.-Q.; Chen, C.; Pentecost, A.; Maleski, K.; Mathis, T.; Zhang, X.-Q.; Zhang, Q.; Jiang, J.; Gogotsi, Y. Nanodiamonds Suppress the Growth of Lithium Dendrites. Nat. Commun. 2017, 8, 336. (34) Wu, F.; Qian, J.; Chen, R.; Lu, J.; Li, L.; Wu, H.; Chen, J.; Zhao, T.; Ye, Y.; Amine, K. An Effective Approach To Protect Lithium Anode and Improve Cycle Performance for Li−S Batteries. ACS Appl. Mater. Interfaces 2014, 6, 15542−15549. (35) Li, Q.; Zhu, S.; Lu, Y. 3D Porous Cu Current Collector/LiMetal Composite Anode for Stable Lithium-Metal Batteries. Adv. Funct. Mater. 2017, 27, 1606422. (36) Wang, S.-H.; Yin, Y.-X.; Zuo, T.-T.; Dong, W.; Li, J.-Y.; Shi, J.L.; Zhang, C.-H.; Li, N.-W.; Li, C.-J.; Guo, Y.-G. Stable Li Metal Anodes via Regulating Lithium Plating/Stripping in Vertically Aligned Microchannels. Adv. Mater. 2017, 29, 1703729. (37) Liu, W.; Lin, D.; Pei, A.; Cui, Y. Stabilizing Lithium Metal Anodes by Uniform Li-Ion Flux Distribution in Nanochannel Confinement. J. Am. Chem. Soc. 2016, 138, 15443. (38) Cheng, X.-B.; Hou, T.-Z.; Zhang, R.; Peng, H.-J.; Zhao, C.-Z.; Huang, J.-Q.; Zhang, Q. Dendrite-Free Lithium Deposition Induced by Uniformly Distributed Lithium Ions for Efficient Lithium Metal Batteries. Adv. Mater. 2016, 28, 2888−2895. (39) Wu, B.; Lochala, J.; Taverne, T.; Xiao, J. The Interplay Between Solid Electrolyte Interface (SEI) and Dendritic Lithium Growth. Nano Energy 2017, 40, 34−41. (40) Zhang, R.; Chen, X.-R.; Chen, X.; Cheng, X.-B.; Zhang, X.-Q.; Yan, C.; Zhang, Q. Lithiophilic Sites in Doped Graphene Guide Uniform Lithium Nucleation for Dendrite-Free Lithium Metal Anodes. Angew. Chem., Int. Ed. 2017, 56, 7764−7768. (41) Zou, P.; Wang, Y.; Chiang, S.-W.; Wang, X.; Kang, F.; Yang, C. Directing Lateral Growth of Lithium Dendrites in Micro-Compartmented Anode Arrays for Safe Lithium Metal Batteries. Nat. Commun. 2018, 9, 464. (42) Pei, A.; Zheng, G.; Shi, F.; Li, Y.; Cui, Y. Nanoscale Nucleation and Growth of Electrodeposited Lithium Metal. Nano Lett. 2017, 17, 1132−1139. (43) Tu, Z.; Kambe, Y.; Lu, Y.; Archer, L. A. Nanoporous PolymerCeramic Composite Electrolytes for Lithium Metal Batteries. Adv. Energy Mater. 2014, 4, 1300654. (44) Tu, Z.; Zachman, M. J.; Choudhury, S.; Wei, S.; Ma, L.; Yang, Y.; Kourkoutis, L. F.; Archer, L. A. Nanoporous Hybrid Electrolytes for High-Energy Batteries Based on Reactive Metal Anodes. Adv. Energy Mater. 2017, 7, 1602367. (45) Han, J.-H.; Khoo, E.; Bai, P.; Bazant, M. Z. Over-limiting Current and Control of Dendritic Growth by Surface Conduction in Nanopores. Sci. Rep. 2015, 4, 7056. (46) Kang, S. J.; Mori, T.; Suk, J.; Kim, D. W.; Kang, Y.; Wilcke, W.; Kim, H.-C. Improved Cycle Efficiency of Lithium Metal Electrodes in Li−O2 Batteries by a Two-Dimensionally Ordered Nanoporous Separator. J. Mater. Chem. A 2014, 2, 9970−9974. (47) Lee, H.; Dellatore, S. M.; Miller, W. M.; Messersmith, P. B. Mussel-Inspired Surface Chemistry for Multifunctional Coatings. Science 2007, 318, 426−430. (48) Schaefer, J. L.; Moganty, S. S.; Yanga, D. A.; Archer, L. A. Nanoporous Hybrid Electrolytes. J. Mater. Chem. 2011, 21, 10094− 10101.

(10) Janek, J.; Zeier, W. G. A Solid Future for Battery Development. Nat. Energy 2016, 1, 16141. (11) Manthiram, A.; Yu, X.; Wang, S. Lithium Battery Chemistries Enabled by Solid-State Electrolytes. Nat. Rev. Mater. 2017, 2, 16103. (12) Zhou, G.; Li, F.; Cheng, H.-M. Progress in Flexible Lithium Batteries and Future Prospects. Energy Environ. Sci. 2014, 1307−1338. (13) Zhu, Y.; He, X.; Mo, Y. Origin of Outstanding Stability in the Lithium Solid Electrolyte Materials: Insights from Thermodynamic Analyses Based on First-Principles Calculations. ACS Appl. Mater. Interfaces 2015, 7, 23685−23693. (14) Wen, K.; Liu, L.; Chen, S.; Zhang, S. A Bidirectional Growth Mechanism for a Stable Lithium Anode by a Platinum Nanolayer Sputtered on a Polypropylene Separator. RSC Adv. 2018, 8, 13034− 13039. (15) Zhang, J.; Zhao, J.; Yue, L.; Wang, Q.; Chai, J.; Liu, Z.; Zhou, X.; Li, H.; Guo, Y.; Cui, G.; Chen, L. Safety-Reinforced Poly(Propylene Carbonate)-Based All-Solid-State Polymer Electrolyte for AmbientTemperature Solid Polymer Lithium Batteries. Adv. Energy Mater. 2015, 5, 1501082. (16) Pan, Q.; Smith, D. M.; Qi, H.; Wang, S.; Li, C. Y. Hybrid Electrolytes with Controlled Network Structures for Lithium Metal Batteries. Adv. Mater. 2015, 27, 5995−6001. (17) Zeng, X. X.; Yin, Y. X.; Li, N. W.; Du, W. C.; Guo, Y. G.; Wan, L. J. Reshaping Lithium Plating/Stripping Behavior via Bifunctional Polymer Electrolyte for Room-Temperature Solid Li Metal Batteries. J. Am. Chem. Soc. 2016, 138, 15825−15828. (18) Tung, S.-O.; Ho, S.; Yang, M.; Zhang, R.; Kotov, N. A. A Dendrite-Suppressing Composite Ion Conductor from Aramid Nanofibres. Nat. Commun. 2015, 6, 6152. (19) Chinnam, P. R.; Wunder, S. L. Engineered Interfaces in Hybrid Ceramic−Polymer Electrolytes for Use in All-Solid-State Li Batteries. ACS Energy Lett. 2017, 2, 134−138. (20) Liu, Y.; Lin, D.; Jin, Y.; Liu, K.; Tao, X.; Zhang, Q.; Zhang, X.; Cui, Y. Transforming from Planar to Three-Dimensional Lithium with Flowable Interphase for Solid Lithium Metal Batteries. Sci. Adv. 2017, 3, eaao0713. (21) Yang, L.; Wang, Z.; Feng, Y.; Tan, R.; Zuo, Y.; Gao, R.; Zhao, Y.; Han, L.; Wang, Z.; Pan, F. Flexible Composite Solid Electrolyte Facilitating Highly Stable “Soft Contacting” Li−Electrolyte Interface for Solid State Lithium-Ion Batteries. Adv. Energy Mater. 2017, 7, 1701437. (22) Xue, Z.; He, D.; Xie, X. Poly(ethylene oxide)-Based Electrolytes for Lithium-Ion Batteries. J. Mater. Chem. A 2015, 3, 19218−19253. (23) Fu, X.; Yu, D.; Zhou, J.; Li, S.; Gao, X.; Han, Y.; Qi, P.; Feng, X.; Wang, B. Inorganic and Organic Hybrid Solid Electrolytes for LithiumIon Batteries. CrystEngComm 2016, 18, 4236−4258. (24) Zhou, W.; Wang, S.; Li, Y.; Xin, S.; Manthiram, A.; Goodenough, J. B. Plating a Dendrite-Free Lithium Anode with a Polymer/Ceramic/Polymer Sandwich Electrolyte. J. Am. Chem. Soc. 2016, 138, 9385−9388. (25) Li, Y.; Xu, B.; Xu, H.; Duan, H.; Lv, X.; Xin, S.; Zhou, W.; Xue, L.; Fu, G.; Manthiram, A.; Goodenough, J. B. Hybrid Polymer/Garnet Electrolyte with a Small Interfacial Resistance for Lithium-Ion Batteries. Angew. Chem., Int. Ed. 2017, 56, 753−756. (26) Fu, K.; Gong, Y.; Dai, J.; Gong, A.; Han, X.; Yao, Y.; Wang, C.; Wang, Y.; Chen, Y.; Yan, C.; Li, Y.; Waachsman, E. D.; Hu, L. Flexible, Solid-State, Ion-Conducting Membrane with 3D Garnet Nanofiber Network for Lithium Batteries. Proc. Natl. Acad. Sci. U. S. A. 2016, 113, 7094−7099. (27) Lu, Y.; Tu, Z.; Archer, L. A. Stable Lithium Electrodeposition in Liquid and Nanoporous Solid Electrolytes. Nat. Mater. 2014, 13, 961− 969. (28) Ding, F.; Xu, W.; Graff, G. L.; Zhang, J.; Sushko, M. L.; Chen, X.; Shao, Y.; Engelhard, M. H.; Nie, Z.; Xiao, J.; Liu, X.; Sushko, P. V.; Liu, J.; Zhang, J. G. Dendrite-Free Lithium Deposition via Self-Healing Electrostatic Shield Mechanism. J. Am. Chem. Soc. 2013, 135, 4450− 4456. (29) Choudhury, S.; Tu, Z.; Stalin, S.; Vu, D.; Fawole, K.; Gunceler, D.; Sundararaman, R.; Archer, L. A. Electroless Formation of Hybrid 20420

DOI: 10.1021/acsami.8b03391 ACS Appl. Mater. Interfaces 2018, 10, 20412−20421

Research Article

ACS Applied Materials & Interfaces (49) Tikekar, M. D.; Archer, L. A.; Koch, D. L. Stability Analysis of Electrodeposition across a Structured Electrolyte with Immobilized Anions. J. Electrochem. Soc. 2014, 161, A847−A855. (50) Yamaguchi, T.; Nagao, A.; Matsuoka, T.; Koda, S. A Theoretical Study on the Anomalous Pressure Dependence of the Transport Properties of Ionic Liquids: Comparison Among Lithium Bromide, Silica, and Water. J. Chem. Phys. 2003, 119, 11306−11317. (51) Borodin, O.; Smith, G. D.; Henderson, W. Li + Cation Environment, Transport, and Mechanical Properties of the LiTFSI Doped N-Methyl-N-alkylpyrrolidinium + TFSI− Ionic Liquids. J. Phys. Chem. B 2006, 110, 16879−16886. (52) Borodin, O.; Smith, G. D. Structure and Dynamics of N-MethylN-propylpyrrolidinium Bis(trifluoromethanesulfonyl)imide Ionic Liquid from Molecular Dynamics Simulations. J. Phys. Chem. B 2006, 110, 11481−11490. (53) Deshpande, A.; Kariyawasam, L.; Dutta, P.; Banerjee, S. Enhancement of Lithium Ion Mobility in Ionic Liquid Electrolytes in Presence of Additives. J. Phys. Chem. C 2013, 117, 25343−25351. (54) Bayley, P. M.; Best, A. S.; MacFarlane, D. R.; Forsyth, M. The Effect of Coordinating and Non-Coordinating Additives on the Transport Properties in Ionic Liquid Electrolytes for Lithium Batteries. Phys. Chem. Chem. Phys. 2011, 13, 4632−4640. (55) Lin, D.; Liu, Y.; Cui, Y. Reviving the Lithium Metal Anode for High-Energy Batteries. Nat. Nanotechnol. 2017, 12, 194−206. (56) Suo, L.; Hu, Y.-S.; Li, H.; Armand, M.; Chen, L. A New Class of Solvent-in-Salt Electrolyte for High-Energy Rechargeable Metallic Lithium Batteries. Nat. Commun. 2013, 4, 1481. (57) Mauger, A.; Armand, M.; Julien, C. M.; Zaghib, K. Challenges and Issues Facing Lithium Metal for Solid-State Rechargeable Batteries. J. Power Sources 2017, 353, 333−342. (58) Zhao, C.-Z.; Zhang, X.-Q.; Cheng, X.-B.; Zhang, R.; Xu, R.; Chen, P.-Y.; Peng, H.-J.; Huang, J.-Q.; Zhang, Q. An AnionImmobilized Composite Electrolyte for Dendrite-Free Lithium Metal Anodes. Proc. Natl. Acad. Sci. U. S. A. 2017, 114, 11069−11074. (59) Xu, R.; Zhang, X.-Q.; Cheng, X.-B.; Peng, H.-J.; Zhao, C.-Z.; Yan, C.; Huang, J.-Q. Artifcial Soft−Rigid Protective Layer for Dendrite-Free Lithium Metal Anode. Adv. Funct. Mater. 2018, 28, 1705838. (60) Chen, L.; Li, Y.; Li, S.-P.; Fan, L.-Z.; Nan, C.-W.; Goodenough, J. B. PEO/Garnet Composite Electrolytes for Solid-State Lithium Batteries: From “Ceramic-in-Polymer” to “Polymer-in-Ceramic. Nano Energy 2018, 46, 176−184. (61) Plimpton, S. Fast Parallel Algorithms for Short-Range Molecular Dynamics. J. Comput. Phys. 1995, 117, 1−19. (62) Dushanov, E.; Kholmurodov, Kh.; Yasuoka, K. Molecular Dynamics Studies of the Interaction Between Water and Oxide Surfaces. Physics of Particles and Nuclei Letters 2012, 9, 541−551. (63) Wu, X.; Liu, Z.; Huang, S.; Wang, W. Molecular Dynamics Simulation of Room-Temperature Ionic Liquid Mixture of [Bmim][BF4] and Acetonitrile by a Refined Force Field. Phys. Chem. Chem. Phys. 2005, 7, 2771−2779. (64) Hockney, R. W.; Eastwood, J. W. Computer simulation using particles; Taylor & Francis, Inc.: Bristol, PA, USA. 1988. (65) Wang, Y.; Qin, Z.; Buehler, M. J.; Xu, Z. Intercalated Water Layers Promote Thermal Dissipation at Bio-Nano Interfaces. Nat. Commun. 2016, 7, 12854. (66) Sun, H. Ab initio calculations and force field development for computer simulation of polysilanes. Macromolecules 1995, 28, 701− 712.

20421

DOI: 10.1021/acsami.8b03391 ACS Appl. Mater. Interfaces 2018, 10, 20412−20421