SolidâLiquid Electrolyte as a Nanoion Modulator for Dendrite-Free

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A Solid-Liquid Electrolyte as Nano-IonModulator for Dendrites Free Lithium Anodes Kaihua Wen, Yanlei Wang, Shimou Chen, Xi Wang, Suojiang Zhang, and Lynden A. Archer ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.8b03391 • Publication Date (Web): 01 Jun 2018 Downloaded from http://pubs.acs.org on June 1, 2018

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ACS Applied Materials & Interfaces

A Solid-Liquid Electrolyte as Nano-Ion-Modulator for Dendrites Free Lithium Anodes Kaihua Wen,†, ‡ Yanlei Wang,† Shimou Chen,*,† Xi Wang,§ Suojiang Zhang,* ,† and Lynden A. Archer‖ †

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, NY

14850, USA. *Corresponding Author E-mail: [email protected]; [email protected].

Keywords: solid-liquid electrolyte; nanochannel confinement; nano-ion-mudulator; stable electrodeposit; lithium metal batteries

ABSRACT: 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, channel phase and non-channel phase, respectively. The channel phase can achieve homogeneous Li+ flux distribution by multiple

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mechanisms, including its uniform array of nanochannels and ability to suppress lateral dendrite growth by it high modulus. In the non-channel 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 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.

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 2,567 Wh kg-1 and 3,505 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 the 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 Li anode,


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consequently resulting in irregular large volume changes, which drive the morphological instability termed dendrite growth.8 Thickening of 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 polyethylene 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 advantages of beneficial aspects of SPEs and SIEs, but which utilize inorganic fillers and plasticizers, can increase ionic conductivity, but even in state-of-the art materials it is difficult to achieve the combination of high bulk and interfacial conductivities that define 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 nano-diamonds33 and novel Li salt34 have been reported to facilitate improved ion transport and stability of solid electrolyte interface (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 micro-compartmented anode arrays41 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 nano-ion-modulator (NIM) that allows facile regulation of Li-ion transport for stable Li deposition. Nanoporous anodic aluminium 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 Li surface. 43-46 It was predicted that a transition from stable to unstable deposition when Li nuclei size exceed a value around 200 nm. 9, 44

In our work, we focus on the AAO membranes with well-aligned nanochannels

with pore sizes in the range 20-400 nm. Besides, PEG chains in the non-channel 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 indicating the great possibility of using this SLE with a mechanism of nano-ion-modulator for Li metal based energy storage system.

RESULTS AND DISCUSSION Figure 1a summarizes the synthesis procedure for the SLE films. After a series of polarization and grafting, a uniform PEG-modified AAO film was prepared.47 Figure


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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, 74.24 eV corresponding to C1s, N1s, O1s, Al2p, respectively. Table S1, S2 shows surface chemical composition of unmodified AAO membrane and PDA-coated AAO before and after PEG immobilization. The data was obtained from peak intensities of Figure 1d, S2. The N1s content of AAO membranes before and after PDA coated 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 C1s is cooperated with O1s to form 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 remain 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 Figure 1e, 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 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 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 nano-layer can efficiently reduce the interfacial impendence against Li anode, owning to the desolvation of PEG segments.48, 49

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.


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(e) Contact angle image of A@PP-5 film.

We analyzed the transport of Li+ in the two phases using molecular dynamics (MD) simulations (Figure 2a, b). In channel phase (CP), Li+ flux can be uniformly distributed in well-aligned nanochannels, facilitating homogenous Li deposition and stable SEI layers. In the simulations, we compare the diffusion coefficients in

different pore sizes, from 1 nm to 15 nm as shown in Figure 2c. We can see that the diffusion coefficient increases with the pore diameter, and always smaller than that in the bulk case. The diffusive coefficient for Li+ in bulk case is D = 0.966×10-11 m2/s, which is on the same order of LiTFSI in other solvent 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 is observed to increase. Based on the above simulation, an analytical relationship between the diffusion coefficients and pore sizes can be proposed (Figure 2c). Results reported in Figure S5 shows that the ionic conductivity follows a similar relation. Conversely, when the size of channel increases, the growth of Li dendrites could not be suppressed. Taken together, these results suggested that channels, with a pore size between 145 to 250 nm, could be considerable candidates (Figure S6). To investigated the interfacial structure of lithium salt and PEG molecules, we also constructed a model consist of 20 chains of PEG (20 monomers per chain), 43



Li+, 43 TFSI-, 510 PC molecules. In non-channel 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.

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


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the theoretical prediction based on the MD results. (d) The interaction distance between PEG and solvents, which shows that PC have a close contact with PEG compared 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.

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 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 towards tips, which finally results in large dendrites. In 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 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 Li surface and enhance ion transport on electrolyte-electrode interface, a PEG modified AAO is introduced. In our nano-ion-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 molecular and eliminate space charge layer. This configuration facilitates the application of LMBs with stable Li deposition and long lifetime performance.



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Figure 3. (a) Temperature dependent ionic conductivity of A@PP-5/LE, comparing with PE/LE, ranging from

0 °C 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 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 °C to 40 °C. Inset shows impedance spectra of A@PP-5/LE against temperatures ranging from 50 °C to 100 °C. (e) resultant Rb and Rint of 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.

Figure 3a illustrates the conductivity of A@PP-5/LE in comparison to PE/LE at temperatures

ranging

from

0~100

°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 one 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


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interfacial resistance changed from 105.1 Ω to 118.5 Ω after polarization process. The Li-ion transference number can be calculated by Eq. 4 to be t Li+=0.83, which is much higher than that measured in a PE separator (Figure S9, t than typical values (t

Li+

conclude that the high t

Li+ =

0.39). It is also higher

< 0.5) reported for PEO-based polymer electrolytes.57 We Li+

of A@PP-5/LE makes it possible to be applied in 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, 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 still larger than 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 PEG layers. The Rint for A@PP-2/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 were 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



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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 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.

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 hours. (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.

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 hours, respectively. Figure 4a shows the appearance of large, tree-like dendrites in PE/LE system. In 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.


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These nuclei are much larger than the pores of nanchannels (~200 nm). To study long-time Li deposition, EIS were 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 cm-2) and 88 Ω cm-2 (1 mA cm-2) after 20 cycles, indicating that AAO/LE can make homogenous 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), then the values of Rint drops 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 glove box. 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 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 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.

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 hours.


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To further evaluate the electrochemical compatibility of A@PP-5/LE against 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 over-potential 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 limit as 22 mV for more than 6000 hours (Figure 5c).

Figure 6. Galvanostatic cycling of Li/LTO batteries with A@PP-5/LE (red), A200/LE (blue) and PE/LE (black) at



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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 Figure 6a,6b and 6c, respectively.

To test the practical performance of SLEs, we studied Li|Li4Ti5O12 (LTO) cells at 60 °C. Figure 6a, b present 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 C 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 C and 5 C are 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, shuc as liquid-like ionic conductivity, low Rint, and long-life cell operations. In addition, the long-term cycling and rate behaviors at room temperature are shown in Figure S16, S17. This system can also be applied in Li|LiFePO4 battery (Figure S18). These results again show the potential importance of SLEs for practical


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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 Li anode at both room and elevated temperatures. The materials were assembled as Li/LTO half cells to investigate their longtime cycling performance. These measurments show 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 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.

EXPERIMENTAL SECTION Molecular dynamics simulations All the molecular dynamics (MD) simulations in this work were completed using large-scale atomic/molecular massively parallel simulator (LAMMPS).61 The timestep for integrating Newtonian equations of motion was 0.5 fs. To simulate the LiTFSI+PC liquids in the nanochannels, we built a model as shown in Figure S4, where the wall ACS Paragon Plus Environment


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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 examples, 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, 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 adopted along 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 truncate at 1.2 nm. The later one, long-range Columbic interactions, was computed using the particle-particle-particel-mesh (PPPM) algorithm.64 The system was firstly relaxed at a temperature of 300 K and 1.0 atmosphere using 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 analyse 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 the atomic trajectories from MD simulations.


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To investigated the interfacial structure of lithium salt and PEG molecules, we also constructed a model consist of 20 chains of PEG (20 monomers per chain), 43 Li+, 43 TFSI-, 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 two parts: van der Waals interaction and electrostatic force. The system was first equilibrate in NPT ensemble at a temperature of 300 K and 1.0 atmosphere using Berendsen thermostat for 5 ns. Then another 1 ns simulations were performed to capture the interfacial structure of lithium salt-PEG. Preparation of the hybrid electrolyte The polymerization and grafting reactions reacted in a buffur solution. 0.2422 g Tris base (supplied by NOVON) and 60 µL hydrochloric acid (HCl, supplied by Beijing Chemical Works) was put into 200 mL 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 aluminium 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 hours 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-hour-saking, the films were taken out, washed with distilled water and absolute ethyl alcohol for 3 times, respectively, then dried at 80 °C under vacuum for 12 h. Finally, the dried pieces were transformed into a glove box 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 were 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 HES/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 were 2.61 mg cm-2 and 3.12 mg cm-2, respectively. Characterization and performance evaluation The surface morphology of the SLE were examined by using an Ultra-high Resolution Scanning Electron Microscope (SEM, Haitich SU8020, Japan) with an acceleration voltage of 5 kV before immersing in liquid 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 were 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


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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 Eq. 1: EU(%) = [(W2–W1)/W1] × 100%, 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 were tested with contact angle machine. The thermal property 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 were 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 °C to 100 °C. The samples for the measurements were prepared by sandwiching the SLEs between two stainless steel discs electrodes (Ф=10 mm). The ionic conductivity was calculated by the following Eq. 2: σ = (1/R) × (d/S), 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, Eq. 3: σ = AT-1/2exp[-E0/R(T-T0)], where A is a constant proportional to the number carrier ions, E0 is the pseudoactivation energy barrier, T is the measurement temperatures and T0 is the ideal glass transition temperature. The Li-ion transference number (t

Li+)

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 Eq. 4: t Li+

= Is(∆V–R0I0)/[I0(∆V–RsIs)], where I0 and Is are the initial and steady-state DC

current respectively, R0 and Rs are the interfacial resistances of initial and steady-state, and ∆V 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 works 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 °C 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 mins, and then discharged at the same current density for 180 mins, following with charging 180 mins 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 C and 5 C) between 1.0 V to 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, 20 C and then 0.5 C, 1 C between 1.0 V to 3.0 V for 10 cycles for each rate.


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ASSOCIATED CONTENT Supporting Information The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami.XXXXXXX. SEM image of cross-section 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 non-channel 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, cycling and rate performances at room temperature (PDF)

AUTHOR INFORMATION Corresponding Author *Shimou Chen: E-mail: [email protected]. *Suojiang Zhang: E-mail: [email protected]. ORCID Yanlei Wang: 0000-0002-2214-8781. Shimou Chen: 0000-0002-2533-4010. Suojiang Zhang: 0000-0002-9397-954X. Lynden A. Archer: 0000-0001-9032-2772.



Notes The authors declare no competing financial interest.

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-A-08), Beijing Municipal Science and Technology Project (D171100005617001) and Henan province science and technology cooperation project (172106000061)

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Figure 1 199x164mm (300 x 300 DPI)



Figure 2 199x165mm (300 x 300 DPI)


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Figure 3 199x104mm (300 x 300 DPI)



Figure 4 199x128mm (300 x 300 DPI)


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Figure 5 199x234mm (300 x 300 DPI)



Figure 6 199x136mm (300 x 300 DPI)


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Table of Contents Graphic 82x44mm (300 x 300 DPI)


SolidâLiquid Electrolyte as a Nanoion Modulator for Dendrite-Free

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