Interfacial Barrier Free Organic-Inorganic Hybrid Electrolytes for All Solid-State Batteries

Organic-inorganic hybrid solid electrolytes (HSEs) are expected to overcome the inherent limitations of rigid fragile inorganic electrolytes for solid state batteries. Li-ion conductive ller such as garnet Li 7 La 3 Zr 2 O 12 (LLZO) is proposed for the high performance of HSEs, unfortunately, which suffers from native surface layer resistance to Li-ion transport. Here we present highly conductive polyvinylidene uoride (PVDF)-based HSEs incorporating LLZO llers, whose resistive barriers are eliminated by dry etching. Our optimal composition of etched LLZO llers (30 wt%) leads to ionic conductivity of 4.05 x 10 -4 S cm -1 , about two-fold improvement from non-etched counterpart. Li symmetric cells with etched llers exhibit low interfacial resistance of 110 Ω cm 2 and minimal overpotential of 46 mV. Moreover, high capacity of 79 mA h g -1 is highlighted at 4C, comparable or superior to liquid electrolyte or sulde-based electrolyte devices. Interfacial environment in HSEs ideally modied for Li-ion transport is identied by 7 Li NMR measurements.


Introduction
Since the rst commercialization in 1991, Li-ion battery (LIB) has ushered in a new era of mobile devices, including mobile phones and electric vehicles (EVs).1,2 Unfortunately, continued accidents relevant to the ignition or explosion of LIBs have raised safety concerns for more reliable LIB systems.3,4 Despite the accumulated research efforts for device design and material advancement, however, potential danger still remains, particularly involved with the ammable organic liquid electrolytes (LEs). Solid state batteries (SSBs) have been introduced in this regard, as well as to exploit pure Li metal anode with higher capacity.
SSBs also offer a wide range of operational voltages, excellent thermal stability and high voltage design by stacking bipolar electrodes.5-8 The ultimate performance of SSB is strongly governed by solid electrolytes (SEs). In the early days of SE research, high ionic conductivity was the top priority, exploiting the inorganic oxides to sul des.9-18 Recently, long-term stability under harsh electrochemical environments have re-emerged as the next critical issue for a practical utilization.
Li-stuffed garnet oxides with the nominal composition of Li7La3Zr2O12 (LLZO) are promising SEs, rst introduced by Murugan et al. in 2007. 19 Garnet LLZO not only shows genuine high ionic conductivity (10-4~10-3 S cm-1) relying on the 3D connected Li-ion network20 but also mechanical robustness to suppress surface dendrite formation and thermal stability for a wide operating temperature window. 21,22 Along with chemical modi cation and processing methodologies for the further improvement of Li-ion conductivity,23-27 moreover, garnet LLZO is found to possess the lowest reduction potential against Li metal among all previously reported SEs. 28 Presently, several challenges are remaining for the practical utilization of LLZO in SSBs. First, direct device integration of rigid and fragile ceramic LLZO is di cult. 29,30 Surface defects in the LLZO may cause the internal short circuit through Li metal pentration. 31 Second, scale up of large-size LLZO membrane for practical cells demands formidable cost and sophisticated handling process. Current studies mostly rely on pellet-type LLZO, which is inadequate for eventual scaling up. 24,29,30 Finally, LLZO frequently suffers from instability under an ambient atmosphere, due to the formation of highly resistive surface layers consisting of Li2CO3. [32][33][34][35] Motivated from the synergistic potential from organic and inorganic components, hybrid polymer electrolytes (HSEs) has emerged as an ideal composition for SSB. While a polymer matrix offers safe, lightweight and soft intimate contact/strong interfacial adhesion, highly ionic conductive inorganic SE llers can be effectively integrated without burden for brittle ceramic processing. So far, there have been a few reports for the successful utilization of HSEs with LLZO llers, particularly exploiting poly(vinylidene uoride) (PVDF) matrix with a high dielectric constant.36-38 Nevertheless, little attention has been paid to the resistive surface layer formation on LLZO llers. In this work, we introduce a straightforward approach for highly conductive PVDF-LLZO HSEs by eliminating the resistive surface layers of garnet LLZO llers. To the best of our knowledge, this is the rst demonstration for the effective removal of inactive surface layers at powder-type LLZO llers by reliable dry etching process. Along with the integration of PVDF-based HSEs lled with etched LLZO as the electrolytes, SSBs based on LiNi0.6Mn0.2Co0.2/LLZO-PVDF HSEs/Li noticeably improved the rate capability and cycling performance. Systematic monitoring of Li-ion pathway veri es the augmented Li-ion movements along the activated local environment around LLZO llers.

Results
LLZO llers with cubic phase were obtained by a solid-state synthesis employing double-doping strategy for high ionic conductivity and short reaction time. 27 Resistive Li2CO3 layers were formed at the LLZO surface upon unavoidable exposure to ambient air. Transmission electron microscopy (TEM) visualizes the Li2CO3 layer of ~15 nm thickness (Fig. 1a). 39 The poorly crystallized Li2CO3 layers are hardly detectable by X-ray diffraction (XRD) (Fig. S1). Fig. 1b schematically describes our well-re ned dry etching process for the resistive layers. Reactive ion etching (RIE) is a reliable dry etching by means of ion bombardment, widely used in the microfabrication. 40 Herein, RIE strips off the thin Li2CO3 layers at the powder state LLZO llers, which is di cult with other methods, such as mechanical polishing.39,41,42 In a RIE process, the degree of anisotropy for etching rate, generally expressed as A = 1-RL/RV (RL : lateral etch rate, RV : vertical etch rate), is expected to be 1 for an ideal vertical etching. 43 Practical RIE cannot feature such an ideal case and sidewall passivation is commonly used to prevent the undercut beneath the etching mask in microelectromechanical system (MEMS) fabrication. 44 In addition, the typical etching rate with uorocarbon gases such as CF4 is signi cantly higher for the amorphous surface layer rather than crystalline core.45 Such a non-ideal directionality of RIE along with high etching selectivity could lead to a conformal etching of Li2CO3 layers. Previously, simple thermal treatment has been exploited to remove the Li2CO3 layers. 46 Nonetheless, our etching approach can be generally extended to the unwanted surface layers of other materials even with high thermal vulnerability. We used CF4 RIE for the selective etching of amorphous surface layers. The crystalline structure of LLZO core was wellmaintained after surface etching (Fig. S1). X-ray photoelectron spectroscopy (XPS) was carried out to con rm the surface chemistry of LLZO llers.
Two distinctive peaks were detected in C 1s scan around 285 and 290 eV for adventitious carbon and carbonate, respectively (Fig. 1c).32,41 The carbonate peak diminished with the removal of Li2CO3 layers.
In O 1s spectra, the carbonate peak at ~532.1 eV decreased with etching, while the peak for crystalline LLZO increased at ~528.5 eV.34 Similarly, Li 1s spectra consist of the two principal peaks at ~55.4 and 54.8 eV for Li2CO3 and Li-O, respectively (Fig. 1e). Along with etching time, the intensity of Li2CO3 peak decreases and the overall Li 1s peak shifts from ~55.1 to ~54.8 eV, supporting the removal of Li2CO3. There is no signi cant difference in the 15 and 20 min etching, implying that 15 min is su cient to remove the majority of Li2CO3 layers (Fig. S2). Fig. 1f displays the prominent change of Zr 3d signal at LLZO upon surface etching. A weaker Zr 3d signal was observed at non-etched LLZO, which is quite plausible considering the presence of Li2CO3 layer and the detection limit of ~10 nm for XPS. After etching, the intensity of Zr 3d signal strengthened. Quantitative analysis for the etching pro le is implemented by plotting the elemental ratios of C/La and C/Zr as a function of etching time (Fig. 1g). Both ratios reduce rapidly within 10 min and saturates around 15 min (Those ratios are higher than zero throughout the etching process due to adventitious carbon).39,41 No signi cant difference in the carbon content was detected between 15 min etched sample and others treated with different conditions (Fig. S3). It is also critical to keep the etched samples in an inert atmosphere to avoid the re-growth of surface layer (Fig. S4). Energy dispersive X-ray spectroscopy (EDS) mapping could directly visualize the surface selective etching of Li2CO3 layers ( Fig.   1f and Fig S5). While the C signal became blurred after 15 min etching, no signi cant compositional change was detected in the crystalline LLZO core, as also supported by Raman spectroscopy (Fig. S6).
The typical properties of HSEs strongly depend on the size, morphology and content of inorganic llers. 47 Here, planetary ball milling was used for the size reduction of LLZO llers, considering its mechanical simplicity and potential large-scale production.48 Fig. 2a present SEM images of the assynthesized (referred to as asLLZO) and ball-milled (referred to as bmLLZO) LLZO, respectively. Compared with asLLZO (12.97 ± 4 um), bmLLZO shows a considerably decreased size (1.2 ± 0.5 um) ( Fig. 2b). XRD patterns of both samples showed identical pro les, implying the well-maintained crystalline structures during the size reduction process (Fig. S7). Next, free-standing lms of PVDF-LLZO HSE and the PVDF based solid polymer electrolyte (SPE) were prepared by solution-casting (Fig. S8). XRD patterns of LLZO, pure PVDF, PVDF-SPE and PVDF-LLZO HSEs taken at room temperature are displayed in Fig. 2c and Fig. S9. The diffraction peaks for LiClO4 salt is absent in the PVDF-SPE, indicating a successful ionic complexation.37 Upon hybridization, the crystalline phase of LLZO does not change.
Interestingly, the weak crystallinity of PVDF-SPE becomes slightly less apparent in the HSEs loaded with bmLLZO or etched bmLLZO compared to those with asLLZO. The larger PVDF/LLZO interface area with the smaller-size ball-milled LLZO llers should contribute to the further broadening of PVDF-SPE peaks, also evidenced by the planar elemental mapping (Fig. S10). The surface etching of Li2CO3 layers caused no noticeable difference again in the XRD patterns of HSE-bmLLZO30 (30 wt% bmLLZO in HSE) and HSE-etched bmLLZO30 (Fig. 2c).
While PVDF-SPE lm is transparent, HSE lms show a dramatic variation in the color, depending on the dimension, composition and surface chemistry of loaded llers ( Fig. 2d and Fig. S11). The interaction between N,N-Dimethylformamide (DMF) solvent and LLZO is known to provide an alkaline-like environment.37,49 This can induce the dehydro uorination of PVDF and thus HSE-bmLLZO30 turns dark brown. By contrast, HSE-etched bmLLZO30 exhibits light brown since a less alkaline-like condition with etched LLZO might alleviate the degradation of PVDF. Furthermore, the good mechanical exibility is illustrated by the easy deformation of HSE-etched bmLLZO30 lm.
For the optimal choice of LLZO llers, HSE lms were prepared with various ller contents of 0-50 wt%.
Electrochemical impedance spectroscopy (EIS) of all lms presents a similar trend that the ionic conductivity maximizes at 30 wt% content and declines above (Fig. 2e).37,47 The HSE-etched bmLLZO30 displays the maximum ionic conductivity of 4.05 x 10-4 S cm-1 approximately two times higher than HSE-bmLLZO30 (2.12 x10-4 S cm-1) at room temperature. Temperature-dependent conductivities and activation energies are compared where HSE-etched bmLLZO30 shows the lowest activation energy of 0.31 eV ( Fig. 2f and Table S1).
Speci c effect from eliminating Li2CO3 layers on the electrochemical performance has been investigated in a Li symmetric cell structure and SSBs with the HSEs including 30 wt% LLZO llers. All the cells were tested at room temperature. Voltage-time pro les of the cells with PVDF-SPE and HSE-asLLZO30 present stable cycling at a low current density of 0.05 mA cm-2 (Fig. 3a). However, together with the increase of current density, the overpotential abruptly increased and eventually the cell failure occurred by exceeding a pre-set voltage limit (after 168 h for PVDF-SPE, after 244 h for HSE-asLLZO30), which can be ascribed to the large internal resistance. In HSE-asLLZO30, the relatively small PVDF/asLLZO interfacial area causes a high overall cell resistance, possibly by the dendritic grow of Li metal during cycling. Fig. 3b and c presents the cycling results for HSE-bmLLZO30 and HSE-etched bmLLZO30 cells from 0.05 to 0.15 mA cm-2. Both cells displayed similar trends up to 0.10 mA cm-2 in terms of the shape of voltage plateau and cycling stability (Fig. S12). HSE-etched bmLLZO30 maintained the stable cycling with a low overpotential (0.045 mV) and at voltage plateau at 0.15 mA cm-2, whereas HSE-bmLLZO30 suffered from the increase of overpotential from 0.072 to 0.135 mV along with the irregular shape of the voltage hysteresis. The interfacial resistance of HSE-bmLLZO30 and HSE-etched bmLLZO30 was characterized by EIS upon Li stripping/plating test. (Fig. S13). As summarized in Fig. 3d, the interfacial areal speci c resistance (ASR) shows much milder increase in HSE-etched bmLLZO30 (from 72.6 to 110 Ω cm2) compared to HSE-bmLLZO30 (78 to 232.5 Ω cm2). This further supports that removal of the Li2CO3 layers can boost up Li-ion transport and ultimately suppress the increase of cell resistance.
For the demonstration of practical cell performance, SSBs based on the con guration of LiNi0.6Mn0.2Co0.2/LLZO-PVDF HSEs/Li were assembled and cycled at room temperature. The rate performances of LiNi0.6Mn0.2Co0.2/HSE-bmLLZO30/Li and LiNi0.6Mn0.2Co0.2/HSE-etched bmLLZO30/Li are compared in Fig. 3e. There is a similar level of discharge capacity for both samples at low current density (≤ 0.5 C, 1 C = 170 mA g-1). However, a higher current density reveals the superior cycling pro le of the SSB with HSE-etched bmLLZO30. Noticeably, SSB with HSE-etched bmLLZO30 exhibits a high discharge capacity of ~ 79 mA h g-1 at 4 C, which is three times higher than the one with HSE-bmLLZO30 (~26 mA h g-1) as well as comparable to or even higher than the previous conventional batteries based on LiNi0.6Mn0.2Co0.2 cathode with liquid electrolyte 50 or sul de-based electrolyte. 51 This outstanding rate performance is the highest value among HSE-based SSBs employing the high capacity cathode materials such as LiNi0.6Mn0.2Co0.2 or LiFePO4 reported to date (Fig. 3f and Table   S2). The selected charge/discharge curves for each rate step are compared for clari cation (Fig. S14). The voltage-capacity pro les during the long cycling are displayed in the inset in Fig. 3e. SSB with HSEetched bmLLZO30 exhibits a more stable cycling stability than HSE-bmLLZO30. In addition, cycling test was also conducted at a low temperature of 10 ℃ (Fig. S15). The SSB with HSE-etched bmLLZO30 still exhibited the high capacity of 167 and 163 mA h g-1 at 0.1 and 0.2 C, respectively, while the SSB with non-etched LLZO delivered relatively inferior capacity (155 and 133 mA h g-1 at 0.1 and 0.2 C).
Underlying mechanism for the promoted performance of the HSEs with etched LLZO llers could be characterized with solid-state magic angle spinning (MAS) 7Li nuclear magnetic resonance (NMR). The plausible route for Li-ion migration within HSE volume consists of PVDF polymer matrix, LLZO llers and/or the interface between them. From 7Li NMR pro les, the resonance peaks for LLZO llers and PVDF-SPE were observed at 0.36 and -0.63 ppm, respectively (Fig. 4a). The bmLLZO shows a slightly broad and weak pro le compared to the etched counterpart, indicating the in uence from surface barrier layer, resistant to mobile Li-ion. 52 Narrow and intense peak for PVDF-SPE originates from LiClO4 salt, homogeneously dissolved and complexed within the polymeric matrix. Non-blocking Li symmetric cells of 6Li/HSE-bmLLZO30 or HSE-etched bmLLZO30/6Li were assembled and cycled repeatedly at 10 μA cm-2 for every 5 min. 6Li from an electrode passes through HSE and reaches the opposite electrode while partly replacing 7Li in the preferable pathway. Monitoring of the residual 7Li in HSE can trace the possible Li-ion pathways. Noticeably, dissimilar NMR signals were detected in those samples after cycling, suggesting distinct Li-ion migration behaviors. (Fig. S16) Further, the NMR spectra can be deconvoluted to reveal a new peak at -0.28 ppm in addition to the two main signals, which is attributed to locally modi ed environment at LLZO and PVDF-SPE interface ( Fig. 4b and Fig. S17).53,54 Preferable Li-ion pathway was evaluated from the integral area of deconvoluted NMR spectra, as plotted in Fig. 4c. Signi cant reductions in 7Li were observed in the PVDF-SPE (from 30.6 to ~ 16.5 %) and interface (from 24.5 to 7.9 %) particularly for HSE-etched bmLLZO30. These results suggest the preferable Li-ion pathways through the polymer matrix and interface. Obviously, the interface seems more dominant as follows. First, LLZO llers are known to modify the local surrounding environment by increasing mobile Li-ions.55-57 Second, uniform dispersion of LLZO llers increases the total interfacial area in HSE, establishing abundant routes for ion transport, hardly achieved with a lower (< 20 wt%) or higher (> 40 wt%) ller loading (see Fig. 2e). Noteworthy that severe agglomeration of LLZO llers could occur at the higher loading and result in a less generation of active interfaces.37,47

Discussion
We have demonstrated a straightforward yet highly effective interfacial design of HSEs, consisting of PVDF, Li salt and the garnet LLZO, without resistive interfacial barrier. For the rst time, our study shows that the resistive surface layers can be removed in the powder state of inorganic SE llers by reliable dry etching process, which in turn promotes the Li-ion conductive pathways along the modi ed interfaces. Moreover, dry etching approach, which is generically compatible with pre-existing industrial infrastructure, can be utilized in a large-scale production of novel SEs with cost-effectiveness. Mechanically exible and solution-processable HSE lms were readily attained and utilized for SSBs with the record-high rate performance.
Signi cantly, our MAS 7Li NMR analysis clari es the locally modi ed polymer-inorganic interfaces and neighboring environments as the principal pathways for Li-ion transport. As supported by EIS measurements, effective increase of mobile ions in the major pathways is the enhancement mechanism of ionic conductivity in the heterogeneous system. This unprecedented clear picture among the previous controversial mechanisms for HSE enables a novel design concept and principle for the optimal interface, widely useful for the materials and devices relying on the ionic transport through practical heterogeneous structures. Our systematic research unveils a rarely explored opportunity for the high performance interface engineered HSE, generally demanded for the safe and e cient energy storage systems in deformable electronics, microbatteries, electrochemical sensors and so on.

Material Characterization
Morphology of LLZO, PVDF-SPE and PVDF-LLZO HSEs was examined by FESEM (Hitachi, S-4800) and TEM (Talos F200X). EDS-mapping images were obtained using EX-250 (Horiba). XRD patterns were obtained by a X-ray diffractometer (X'pert Pro, Philips, Cu Kα : 1.54056 Å). XPS measurements (Thermo VG Scienti c, K-alpha) provided high-resolution core scans of C 1s, O 1s, Li 1s, Zr 3d and La 3d before and after dry etching process. Raman spectra (Horiba, LabRAM HR Evolution) were measured to con rm the presence of Li2CO3 layers on the LLZO surface. The ionic conductivities and activation energies were characterized by impedance analyzer using a frequency response (Solartron HF 1225, 10-1~105 Hz) in the temperature range of 0 to 100 ℃. Solid state MAS 7Li NMR spectra were obtained using a Bruker AVANCE III 400 spectrometer with 1.9 mm HX-MAS probe at the spinning rate of 25 kHz. As a reference for 7Li chemical shift, solid LiCl at 0 ppm was used.

Electrochemical Characterization
All the cells for electrochemical characterization were assembled into a CR2032 coin-type cell in dry room. EIS measurements were performed by constructing the sandwich-type cell of stainless steel (SS)/SPE or HSEs/SS over the frequency range of 10-1 to 105 Hz at room temperature. To obtain the activation energy, the impedance spectra were examined in the temperature range of 0 to 100 ℃.

Supplementary Files
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