Quantifying the local Li-ion diffusion over the grain boundaries of a protective coating, revealing the impact on the macroscopic Li-ion transport in an all-solid-state battery


 The key challenge for solid-state-batteries is to design electrode-electrolyte interfaces that combine (electro)chemical and mechanical stability with facile Li-ion transport. Typically, this presents conflicting demands, the solid electrolyte-electrode interface-area should be maximized to facilitate high currents, while it should be minimized to reduce the parasitic interface reactions and enhance stability. Addressing these issues would greatly benefit from establishing the impact of interface coatings on local Li-ion transport over the grain boundaries. Here, the three-phase Li-ion transport, between solid electrolyte, coating and electrode is revealed using exchange-NMR, disentangling the detailed quantitative impact of the coating on the Li-ion transport in solid state batteries. A Li2S cathode is coated by LiI, providing a ductile and conductive interface with the argyrodite-sulfur electrolyte, where the exchange-NMR demonstrates that this enhances the interface transport to such an extent that the commonly applied nanosizing of the cathodic mixture can be abandoned. This leads to facile sulfur activation, preventing solid-electrolyte decomposition, in micron-sized cathodic mixtures, that can be related to the role of coatings on the Li-ion transport, providing perspectives for sulfur based solid-state-batteries.


Introduction
All-solid-state lithium(Li)-ion batteries are promising candidates for next-generation high energy density and safe energy storage technology [1][2] . As a liquid free system, it generally does not suffer from leakage and gas generation and the risk of a thermal runaway, inherent to liquid electrolytes used in conventional Li-ion batteries [3][4][5][6][7] . As a result, research has intensi ed towards solid electrolytes that display conductivities approaching, or even exceeding that of liquid electrolytes, including structural families such as LISICON, argyrodites, garnets, and NASICON-type structures 6,[8][9][10] .
A major obstacle for solid state batteries is the high internal resistance for Li + transfer over the solid-solid electrode-electrolyte interface [11][12] which may be due to poorly conducting electrolyte decomposition products 4,13 , contact loss due to volumetric changes 4,14 and space-charges [15][16] . The commonly applied strategy to lower the resistance for Li + charge transport, is to enlarge the electrolyte-electrode interface area, often expressed as "ionic contact area", by nanosizing the electrode and electrolyte particles [17][18][19] . However, with the interface area between electrolyte and electrode, also scale the detrimental chemical and electrochemical reactions at the interfaces between the electrolyte and electrode, raising the internal resistance for Li + transport. Thus, from the perspective of stability, larger electrolyte and electrolyte particles in the micron range are preferred, having the additional advantage of being more suitable for practical production 20 . However, this puts even higher demands on improving the Li + transport over the relatively small ionic contact area between micron sized solid electrolyte-electrode particles.
To address both stability and Li + charge transport at the interfaces has lead through the development of interphases, often realized through coating processes 6,[18][19] . The demands on these interphases are challenging, including (electro)chemical stability towards both electrode and electrolyte, poor electronic conductivity and at least reasonable ion conductivity, and a low grain boundary resistance with both solid electrode and electrolyte. This is typically achieved by good wettability (low interface energy) towards electrode and electrolyte and by soft interphase materials (low Young's modules and/or yield point). One of the challenges is to assess the impact of interphases on the local Li + charge transport, which can provide valuable insights for interphase design and preparation methodologies. Where the macroscopic charge transfer resistance is most often estimated by electrochemical impedance spectroscopy 21-22 , disentangling the three-phase diffusion between the electrode, interphase/coating and electrolyte is yet to be accomplished.
One of the solid electrolyte-electrode combinations where interphase strategies are intensively investigated is sulfur based solid electrolytes in combination with sulfur cathodes, providing a high energy density in combination with cheap raw materials and synthesis. Sul de-based electrolytes (such as Li 2 S-P 2 S 5 and Li 6 PS 5 X (X=Cl, Br and I)) are especially promising due to their high ionic conductivity and relatively low grain boundary resistance [23][24][25][26] . These advantages, unfortunately, come along with a major drawback, which is the very small electrochemical stability window 13,[27][28] . This is one of the main reasons for the extremely sluggish Li + transport over the solid-solid Li 2 S-solid electrolyte interface during battery operation, especially when micron sized solid electrolyte particles are used [3][4][29][30][31] . An additional challenge is the activation of solid sulfur cathodes due to the very low Li + diffusivity [32][33][34] , which also demands a relatively large ionic contact area 35 . To improve the Li + interfacial transport two applied strategies are bilayer solid electrolytes design (porous layer and dense layer) 30,[36][37] and mixing in binary lithium halide salts (such as LiBr, LiI) additives [29][30][31][38][39][40] . Because of the large cost and sophisticated process associated with the former, halide salt addition appears especially promising for practical application. The lower internal resistance of these three-phase mixtures has argued to be the result of the small Young's modulus (softness) of the halide salts, which effectively act as a solid wetting agent for the electrode/electrolyte interface [29][30][31][38][39][40] . However, how the halogen salt affects the local Li + transport over the grain boundaries (electrolyte-halogen salt, halogen salt-electrode and electrolyte-electrode) is di cult to establish, and especially important to develop fundamental understanding and guide the design of the optimal design and preparation of these interphases.
Making use of the ability of solid state nuclear magnetic resonance (NMR) to characterize the spontaneous Li + exchange between solid electrolyte and electrode phases in solid state cathodes [41][42] , the present work explores the three-phase Li + transport between the Li 2 S electrode, LiI coating and argyrodite Li 6 PS 5 Cl solid electrolyte, aiming to gain insight in the role of the coating in Li-ion transport in solid-state batteries. The difference in NMR chemical shift of the three phases allows to unravel the diffusivity between each phase via 6 Li 2D-EXSY exchange experiments. The sluggish Li + diffusion over the Li 2 Sargyrodite interface is drastically improved by the presence of the LiI coating, where the Li + diffusion is shown to proceed through the LiI coating. The activation energy between the three phases is equal to the bulk activation energy for Li + diffusion in LiI, demonstrating the extremely low grain boundary resistance achieved by the facile LiI coating strategy. The practical consequence during battery operation is that this prevents large overpotentials during battery cycling, even at a relatively low ionic contact area between the cathode and microscopic solid electrolyte particles. This is fortuitous, as it makes the working potential fall within the electrochemical stability of the argyrodite solid electrolyte. The result is a solidstate battery that can be activated at a very low applied voltage, and can cycle with high reversibility under a modest 2 MPa pressure for over 200 cycles. The ability to monitor the local Li + transport over the grain boundaries in this three-phase system provides valuable insight in the role of coatings in achieving low interphase resistances using micron sized solid electrolyte particles, guiding the design of stable high-performance interphases, which are crucial aspects for future solid-state batteries.
The solid-state electrolyte Li 6 PS 5 Cl (denoted as LPSC) was prepared by a simple solid-state reaction. The stoichiometric raw materials LiCl (Sigma-Aldrich), P 2 S 5 (Sigma-Aldrich), and Li 2 S (Sigma-Aldrich) were used as the starting materials and were ball milled at 110 rpm, 2hrs with the ZrO 2 coated jars using 18 ZrO 2 balls. After the ball milling, the precursor was sealed in a quartz tube with Ar and then annealed at 550 o C for 15 hours to obtain the LPSC solid electrolyte. These were subsequently crushed with an agate mortar-pestle before using the samples for further measurements. For preparation of the Li 2 S-LiI mixtures, proportional 10 mmol Li 2 S and 3.33 mmol LiI were dissolved into 1 ml ethanol and stirred for 10 mins then heated at 300 o C until totally dry.

Material characterization.
Powder X-ray diffraction (XRD) patterns were collected over a two-theta range of 10-80° to identify the crystalline phases of the prepared materials using CuKα X-rays (1.5406 Å at 45 kV and 40 mA) on an X'Pert Pro X-ray diffractometer (PANalytical). To prevent reaction with moisture and oxygen, the powder materials were sealed in an airtight XRD sample holder in an argon-lled glove box. For the TEM and energy dispersive X-ray (STEM-EDX) investigations, a suspension in dry ethanol was prepared, which was drop casted onto a standard gold grid with a holy carbon lm under a high temperature, inside an argonlled glove box. To prevent any contact with air TEM grids with the sample were loaded into a custommade vacuum transfer TEM holder. TEM measurements were carried out in a FEI-Tecnai operating at 200 kV. FE-SEM (JEOL JSM-6010LA) images were taken under dry Argon conditions, and taken using an accelerating voltage of 10 kV. The depth-pro ling sputtering was conducted by 2 min of sputtering in ve cycles (2 kV, 2 mm × 2 mm), the narrow spectra of particular elements were recorded after each cycle of sputtering. The pass energy used for the hemispheric analyzer was 58.7 eV, and the base pressure of the system was ∼1 × 10 −7 Pa. The estimated sputtering rates are 5 nm/min.

Solid-state lithium battery preparation and electrochemical performance.
Solid state cells were assembled in an argon-lled glove box in the following steps: 88 mg of LPSC was pressed relative tightly under hand pressure as the electrolyte layer, and subsequently 12.5 mg of cathode mixture and In-Li foil were brought into contact to the top and bottom faces of the solid electrolyte pallet.
After that, 2 MPa pressure was applied to press these three layers together for 30 s. The assembled cells were charged and discharged under speci c current densities to evaluate their electrochemical performances.
Solid state NMR measurements were performed on a Bruker Ascend 500 spectrometer (B 0 = 11.7 T) with a NEO console operating at a 6 Li resonance frequency of 73.578 MHz. The π/2 pulse length was determined to be 3.9 μs with an RF eld strength of 64 kHz. Chemical shifts were referenced with respect to a 1 M LiCl solution. A Bruker two channel MAS WVT 4 mm probe was used for all measurements. One pulse experiment was performed under MAS at a spinning speed of 10 kHz. A recycle delay of three times of T 1 was used each time. Two-dimensional exchange spectroscopy (2D-EXSY) measurements were performed for the mixture at a spinning speed of 10 kHz and at various mixing times ranging from 10 ms up to 10 s and at temperatures from 20 to 100 C. All 2D spectra consist of 16 scans for each of the 32 transients, each transient incremented by 8000 ms with a recycle delay of up to 5 s. Exchange between the Li 2 S and LPSC phases was quanti ed by tting the evolution of the cross-peak intensity as a function of T mix to a 3D diffusion model derived from Fick's law, which has been described previously 4, 24 .

Stability of solid electrolytes, nanosized versus micronsized
The electrochemical stability of solid electrolytes was recently shown to be signi cantly lower than initially thought, especially for sul de based solid electrolytes, where consequential decomposition reactions have been shown to have large detrimental impact on the all-solid-state battery performance 10,[43][44] . To demonstrate the redox activity of the sulfur solid electrolyte material and the impact of particle size, several mixtures of sulfur solid electrolytes with carbon are galvanostatically cycled. As shown in Figure 1a and b, when nanosized Li 3 PS 4 (nLPS) and Li 6 PS 5 Cl (nLPSC) are mixed with conductive carbon (referred to as nLPS-C and nLPSC-C respectively), they are readily oxidized at low oxidation potentials, <2 V vs. In (<2.62 V vs. Li/Li + ), at a relatively low current density of 0.033 mA/cm 2 , in agreement with recent ndings 10,25 . This illustrates that the decomposition of nanosized solid electrolytes in cathodic mixtures is the prime reason for the short cycle life of current all-solid-state batteries utilizing sul de solid electrolytes 25 . The combination of a Li 2 S cathode and the LPSC electrolyte has been intensively studied [3][4][29][30] , where the harsh activation process of Li 2 S, caused by the low bulk conductivity of Li 2 S, and sluggish Li + transport between the electrode and electrolyte represent another critical issue. A straightforward approach to achieve easier activation of Li 2 S and to improve the Li + transport between the electrode and electrolyte is to reduce the particle size of both the Li 2 S and the LPSC (here referred to as nLi 2 S-nLPSC-C electrode). Accompanied by a large overpotential of ~800 mV, this results in a substantial capacity during rst charge and discharge and upon subsequent cycling, as shown in Figure  1f and Figure S1, which is however hard to distinguish from the capacity of the electrolyte itself (because oxidation of LPSC is carried by the S 2-/S redox).
The simplest strategy to reduce the contribution of the solid electrolyte to the capacity is to lower the ionic contact area through the use of micron sized solid electrolyte particles. To verify the smaller redox activity, micron sized LPSC (average diameter of 50 mm, as shown in Figure S2) was mixed with carbon, and charged (oxidized) to 3 V vs. In-Li (3.62 vs. Li/Li + ) in the solid-state battery con guration, referred to as the mLPSC-C electrode. The result, shown in Figure 1c, demonstrates nearly no contribution of the LPSC solid electrolyte to the capacity, re ecting the smaller amount of decomposition reactions under the same current density. However, as expected, the small ionic contact area of micron sized LPSC compromises ion transport over the interface, leading to a very small capacity of the sulfur active material, even when nanosized Li 2 S is employed as shown in Figure 1d. Although nanosized Li 2 S displays a slightly larger capacity, the rapid voltage increases to the 3.5 V vs. In-Li (4.12 vs. Li/Li + ) cut-off demonstrates that in both cases the sulfur cathode material is marginally activated. At the same time, decreasing the ionic contact area increases the internal resistance and thus the overpotentials experienced at the solid electrolyte, which will induce decomposition reactions and thus driving a selfamplifying resistance growth towards battery failure.
The mechanical mixing of halogen salts like LiI with Li 2 S and sul de solid electrolytes is an often-applied strategy, to improve conductivity, although the exact mechanism has not been clari ed [29][30][31][38][39] . At present we take a different and more controlled approach to study the detailed impact of LiI on the Li-ion transport over the grain boundaries by NMR exchange experiments, and introduce LiI at the Li 2 S-LPSC interfaces. Rather than using the conventional ball-milling route, this is achieved by introducing LiI via solution, making use of the much better solubility of LiI in ethanol compared to Li 2 S (see supporting information Figure S3). The solution was then evaporated at 300 o C ( Figure S3) to obtain a LiI-Li 2 S (1:3 molar ratio) composite, where LiI precipitates on the surface of Li 2 S as discussed below. This cathode was subsequently hand mixed with LPSC and C (referred to mLi 2 S(LiI)-mLPSC-C) to prepare the cathodic mixture and an all-solid-state mLi 2 S(LiI)-mLPSC-C|mLPSC|In-Li battery was assembled under 2 MPa pressure. This has a large impact on the electrochemical charging, as shown in Figure 1e, demonstrating that the introduction of LiI results in a very low sulfur redox activation plateau at 1.69 V vs. In (2.31 V vs. Li/Li + ) of the micron sized Li 2 S combined with micron sized LPSC. The plateau is followed by a rapid increase in potential which re ects oxidation of the solid electrolyte and/or of LiI (to LiI 3 which is known to occur at ~2.3 V vs. In) [45][46] . To identify the contribution of the LiI and/or LPSC oxidation, this measurement was repeated in a battery without Li 2 S, shown in Figure 1e, leading to charging at a higher voltage marking the oxidation of LPSC and/or LiI. In conclusion, deposition of LiI on Li 2 S via solution, and hand mixing this cathode with LPSC results in an extremely low activation (oxidation) potential for Li 2 S, suggesting that facile Li + transport between the electrode and electrolyte is achieved even for a relatively small ionic contact area between the micronsized solid electrolyte and the electrode particles.

Electrode and LiI coating synthesis and characterization
To understand the role of LiI in the activation of Li 2 S, a detailed structural investigation was performed.
Both LiI and Li 2 S have a cubic structure indexed to the Fd-3m space group. Three Li 2 S-LiI composites were prepared via dissolution and precipitation where Li 2 S:LiI molar ratios of 9:1, 3:1, and 1:1 were added to ethanol, followed by evaporation of the solution at 300 °C. Li 2 S and LiI were also individually dissolved and precipitated from ethanol via evaporation for comparison. X-ray patterns of pristine Li 2 S and LiI precipitated Li 2 S and LiI and the three Li 2 S-LiI composites are provided in Figure 2a and S4. From a cursory inspection it can be observed that the peaks corresponding to the precipitated Li 2 S are much broader than those of the parent Li 2 S, while the peak width of precipitated LiI is comparable to the parent LiI. This indicates that on precipitation smaller primary crystallites of Li 2 S are obtained. In the three composite mixtures, both the Li 2 S and LiI phases could be indexed, albeit with shifts in peak positions of the Li 2 S component indicating changes in lattice parameters of this phase. Rietveld re nement was further performed of all the patterns depicted in Figure S5, and the lattice parameters obtained from the re nement are given in Figure S4. It can be seen that the lattice parameter of LiI (6.025 Å) remains unchanged from that of the pristine material. On the other hand, with increasing amounts of LiI in the composite the lattice parameter of Li 2 S keeps increasing from 5.701 Å (pristine) to 5.750 Å (1Li 2 S:1LiI), which could be due to the much smaller average crystallite size of 9.18 nm (1Li 2 S:1LiI) compared to 162.92 nm (pristine) as shown in Figure S6 or by incorporation of the I in the Li 2 S lattice.
Additional SEM and TEM measurements are performed to study the morphology of the pristine Li 2 S and the Li 2 S-LiI mixture. As shown in Figure 2c, the prepared mixture consists of a microstructure comprising of micron sized secondary particles with a relatively uniform particle size of around 5 mm similar to pristine Li 2 S in Figure 2b. TEM is used to study the morphology at smaller length scales (100 nm). As seen from the TEM image and energy spectrum (Figure 2d), the EDS mapping of the particle surface shows uniform S and I distribution, indicating a mixture on the nanoscale was obtained with this precipitation method, and the LiI was uniformly distributed over the surface structure of Li 2 S particles. To further verify the structure of the Li 2 S-LiI material, XPS depth pro ling was performed as shown in Figure   2e and f. The S 2p XPS signal is relatively low until a depth of ∼100 nm, and vice versa the I 3d is relatively high to approximately the same depth (selected window diameter is as small as 14 mm to locate only a few particles). Therefore, the present precipitation method results in micron sized secondary cathode particles, referred to as mLi 2 S(LiI), that exist of agglomerates of LiI coated nanosized primary Li 2 S particles, where the micron sized agglomerates are coated by a relatively thick LiI layer at some positions accumulating to large domains of LiI (as observed with XRD).

Li-ion conductivity and role of LiI in the Li-ion transport mechanism
To investigate the impact of the LiI coating on the conductivity, impedance spectroscopy and 6 Li solid NMR spectroscopy experiments are conducted. The temperature dependence of the ionic conductivities for the pristine Li 2 S and LiI as well as pellets of the 3:1 Li 2 S-LiI composite are presented in Figure 3a. The conductivity of all the materials follow an Arrhenius law, resulting in activation energies of 0.235, 0.107 and 0.212 eV for Li 2 S, LiI and the Li 2 S-LiI mixture, respectively. The room-temperature conductivity of the Li 2 S-LiI mixture (6.72 × 10 −9 S/cm at 25 °C) is between that of Li 2 S (7.51 × 10 −11 S/cm at 25 °C) and LiI  Figure 3c. The intensity of these cross-peaks re ects the amount of Li + exchange, which is expected to increase when the diffusion time (T mix ) and temperature are increased 4 . The absence of off-diagonal intensity, even for the maximum T mix and temperature (T mix = 20s, 373K) indicates that the Li + exchange ( ux) over the solid-solid interface between LPSC and Li 2 S (without LiI coating), is too small to be observed, re ecting sluggish Li + mobility across the interface with the solid electrolyte. This rationalizes the observation in Figure 1d, that these mixtures do not facilitate activation of Li 2 S. As expected, the addition of LiI to the cathodic mixture, mLi 2 S(LiI)-mLPSC, results in the appearance of the Li resonance at -4.56 ppm associated with LiI, in the 1D 6 Li NMR spectrum (Figure 3d). The impact of the LiI on the spontaneous Li + charge transfer, between Li 2 S and LPSC is dramatic, as can be observed in Figure   3e-i. At short mixing time, T mix = 10 ms, no appreciable cross-peak intensity is observed in the 2D EXSY spectrum (Figure 3g). However, increasing the mixing time, T mix , to 10 s, and raising the temperature to 373 K, results in a strong cross-peak intensity (Figure 3h and i), which is a measure of the Li + exchange between Li 2 S and LPSC. The evolution of the normalized cross-peak intensity as a function of T mix measured at a range of temperatures and a T mix range of 10 ms−10 s is provided in Figure 3e. Exchange between the Li 2 S and LPSC phases was quanti ed by tting the evolution of the cross-peak intensity as a function of T mix to a diffusion model derived from Fick's law, described elsewhere 24 . From the t, the diffusion coe cient (D) as a function of temperature can be obtained, which in this case pertains to Liion transport across the Li 2 S-LPSC interface. The diffusion coe cients as a function of temperature obtained from the t are given in Figure 3f. The data for Li 2 S-LPSC diffusion can be t to an Arrhenius law, yielding an activation energy of 0.107 eV for the charge transfer. This activation energy equals that of LiI as measured by impedance of the pure phase, suggesting that LiI is responsible for lowering the interfacial barrier between electrode and electrolyte. The value for the activation energy and observed Li + exchange is comparable to that obtained for interfacial diffusion between nanosized Li 2 S and nanosized (average particle size ~ 100 nm) LPSC argyrodite (0.10-0.13 eV) [3][4] . The remarkable conclusion is that, is despite the small ionic contact area of the present micron sized LPSC (average particle size ~ 50 μm) in the cathodic mixtures, the interface transport is improved to such an extent that it matches that of nanostructured mixtures having a much large ionic contact area. The two orders of magnitude difference in diameter between the LPSC in the cathodic mixtures, suggests that the LiI improves the Li + diffusion over the interface with 4 orders of magnitude.
To understand the role of the LiI in the diffusion, Figure 4a focuses on the exchange of Li in Li 2 S and LPSC with LiI, hence the three phase Li-ion exchange in the 2D EXSY measurements shown in Figure 3. Figure 4a displays clear exchange of Li + between LiI and both Li 2 S and LPSC, re ecting the equilibrium exchange of Li + between the three phases in the cathodic mixture. This represents a unique view into the charge transport between the coating and the electrode and electrolyte phases in a solid state battery. By measuring and tting the exchange intensities as a function mixing time and temperature, Figure 4c-f, similar to the evaluation of the direct exchange between Li 2 S and LPSC, the diffusion coe cients and activation energy over both the LPSC-LiI and LiI-Li 2 S interfaces is quanti ed. To the best of our knowledge, this is the rst quanti cation of the local ion diffusion between a coating and its facing solid phases, providing insight in the impact of a coating on the Li-ion transport. The high diffusivity and very low activation energies for charge transfer from Li 2 S to LiI and LPSC to LiI, 0.142 eV and 0.117 eV respectively, are similar to the overall charge transfer between Li 2 S and LPSC. This indicates that LiI facilitates the charge transfer and thus functions as the bridge between electrode and electrolyte as summarized schematically in Figure 4b. Apparently, the ductile LiI 40 creates grain boundaries between both the electrolyte and electrode material that do not pose an additional barrier for Li-ion diffusion, and the diffusivity of LiI itself dictates the diffusivity between electrode and electrolyte.

Electrochemical performance
To test the e cacy of the Li 2 S-LiI cathode in combination with the micron sized LPSC, mLi 2 S(LiI)-mLPSC-C|mLPSC|In-Li, all-solid-state batteries were assembled, during which only a mild pressure, 2MPa, was applied (see methods section for assembly details). The battery performance is shown in Figure 5.
Although the battery can be activated during charge to over 900 mAh/g capacity as shown in Figure 1e, discharge leads to large overpotentials. The increase of the oxidation potential towards 900 mAh/g exceeds the oxidation potential of LiI (2.3 V vs. In-Li) and that of the LPSC electrolyte (2.1 V vs. In-Li), which will result in poorly conducting species near the interfaces that increase the impedance. To prevent this, the battery is charged to speci c capacities ( xed charge capacity) as shown in Figure 5a, followed by discharging to a xed potential (0.8 V vs. In-Li) to achieve complete discharge. After the battery is initially charged to 600 mAh/g directly, the battery was cycled up to 50 cycles at different currents with an average Coulombic e ciency higher than 97.8% (Figure 5b). The 1 st , 25 th and 50 th charge and discharge curves of the mLi 2 S(LiI)-mLPSC-C|mLPSC|In-Li battery cycled at 0.132 mA/cm 2 , shown in Figure 5c, demonstrate an ultra-low activation potential of 1.7 V vs. In-Li, amounting an overpotential of only ~100 mV. This is the lowest overpotential reported (Figure 5d and Table S1) to date and the consequence of the facile Li-ion transport induced by the LiI coating as observed by the exchange NMR experiments. Postmortem XRD analysis after different states of charge, and NMR analysis after 50 cycles (ending in the charged state) of the cycled mLi 2 S(LiI)-mLPSC-C active materials can be found in Figure S7. After the rst charge to 600 mAh/g and even after the 50 th charge to 600 mAh/g, only the Li 2 S is oxidized, towards an amorphous structure, and the micron sized LPSC solid electrolyte remains intact as no decomposition products are observed 47 . The long cycling stability of the battery with a higher mass loading, 6.4 mg/cm 2 , is demonstrated in Figure 5e. After 200 cycles, the Coulombic e ciency maintains values exceeding 99.9% resulting in an average Coulombic e ciency of 99.6%. Most notable is that this is achieved in combination with micron sized electrolyte particles, a small ionic contact area, and that further optimization can be expected upon variation of the amount of LiI, the applied pressure, charge capacity and the cathodic mixture.

Conclusion
In summary, the impact of a ductile coating on the diffusion over the grain boundary between the electrode and solid electrolyte is investigated. 2D exchange NMR reveals and quanti es the spontaneous equilibrium Li + between Li 2 S electrode, LiI coating and argyrodite Li 6 PS 5 Cl solid electrolyte, thus uniquely disentangling the impact of the coating on the Li + diffusion. Effectively, the ductile LiI lowers the barrier for grain boundary diffusion towards both the Li 2 S (electrode) and Li 6 PS 5 Cl (electrolyte) phases to such an extent that the conductivity of the thin LiI coating dominates. This improves the charge transfer between electrode and electrolyte with several orders of magnitude and enables to move from nanostructured solid state cathode mixtures to micron-sized solid state cathode mixtures, the latter having the practical advantages of high stability and facile material and electrode production. The impact of the improved charge transfer is demonstrated by a sul de based solid state battery which combines easy activation of the sulfur electrode at very low overpotentials with stable cycling. This work demonstrates the ability of exchange NMR unambiguously quantify and disentangle the Li + transport over the interfaces between electrode, coating and solid electrolyte (three phase exchange) in solid state batteries. As one of the key challenges towards solid state batteries is the development of interphases to establish stability and facile Li + transport, the present approach and insights provides valuable insights to guide future understanding and material design. Figure 1 Electrochemical activity of nanosized and micronsized solid sulfur electrolytes and solid-state sulfur electrodes. (a-c) voltage pro les of (a) nanosized nLPS-C, (b) nLPSC-C and (c) micron sized mLPSC-C cathodes; In this con guration the solid electrolyte acts as a cathode electrode material due to the conductive carbon mixed in, which allows direct evaluation of the electrochemical activity as these would occur is a solid-state battery, excluding speci c (electro)chemical reactions with a cathode material. (d) Activation voltage pro les of mLi2S-mLPSC-C, nLi2S-mLPSC-C cathodes; (e, f) Voltage pro les of mLi2S(LiI)-mLPSC-C, LiI-mLPSC-C and nLi2S-nLPSC-C cathodes; For all voltage pro les the current density is 0.033 mA/cm2.  Arrhenius law, yielding an activation energy (Ea) of 0.107 eV. (g-i) Two-dimensional 6Li-6Li exchange spectra of the mixture of micron sized LPSC and Li2S-LiI (3:1) cathodes measured at a spinning speed of 10 kHz at (g,h) 293K with mixing times of 0.01s and 10 s and at (i) 373K with a mixing time of 10 s. The spectra consist of 16 scans for each of the 32 slices, each slice incremented by 8 ms with a recycle delay of 5 s. The cross-peak at the off-diagonal positions in the dashed boxes represent diffusion of Li-ions between solid electrolyte and electrode.    Table S1 with the present work; (e) Discharge capacity of the mLi2S(LiI)-mLPSC-C|mLPSC|In-Li battery cycled to 600 mAh/g at 0.264 mA/cm2 current density up to 200 cycles.