Understanding the Effect of Interlayers at the Thiophosphate Solid

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Understanding the Effect of Interlayers at the Thiophosphate Solid Electrolyte/Lithium Interface for All-Solid-State Li Batteries Lingzi Sang, Kimberly L. Bassett, Fernando C. Castro, Matthias J. Young, Lin Chen, Richard T. Haasch, Jeffrey W. Elam, Vinayak P. Dravid, Ralph G Nuzzo, and Andrew A. Gewirth Chem. Mater., Just Accepted Manuscript • DOI: 10.1021/acs.chemmater.8b02368 • Publication Date (Web): 30 Nov 2018 Downloaded from http://pubs.acs.org on December 1, 2018

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Understanding the Effect of Interlayers at the Thiophosphate Solid Electrolyte/Lithium Interface for All-Solid-State Li Batteries Lingzi Sang,† Kimberly L. Bassett,† Fernando C. Castro,  Matthias J. Young,§ Lin Chen,§ Richard T. Haasch,# Jeffrey W. Elam,§ Vinayak P. Dravid,



Ralph G. Nuzzo,†* and

Andrew A. Gewirth†* †Department

of Chemistry and #Frederick Seitz Materials Research Laboratory, University

of Illinois at Urbana−Champaign, Urbana, Illinois 61801, United States Department

of Materials Science and Engineering, Northwestern University, Evanston,

Illinois 60208, United States ‡Chemical

Sciences and Engineering and §Applied Materials Division, Argonne National

Laboratory, Lemont, Illinois 60439, United States Surface

and Corrosion Science, School of Engineering Sciences in Chemistry,

Biotechnology and Health, KTH Royal Institute of Technology, Drottning Kristinasväg 51, 100 44 Stockholm, Sweden TOC figure

Li+ Li+

SiAu

Li

Li0

LPS

LiAlO

Li0

e-

E0 ~ ~ E’SiAu e-

Li+

LPS

E0 < E’LiAlO

Li

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Chemistry of Materials

E0 = 0V (vs. Li/Li+) Li2S Abstract All-solid-state Li-ion batteries afford possibilities to enhance battery safety while improving their energy and power densities. Current challenges to achieving high performance all-solid-state batteries with long cycle life include shorting resulting predominantly from Li dendrite formation and infiltration through the solid electrolyte (SE), and increases in cell impedance induced by SE decomposition at the SE/electrode interface. In this work, we evaluate the electrochemical properties of two interlayer

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Chemistry of Materials 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

materials, Si and LiXAl(2-x/3)O3 (LiAlO), at the Li7P3S11 (LPS)/Li interface. Compared to the Li/LPS/Li symmetric cells in absence of interlayers, the presence of Si and LiAlO both significantly enhance the cycle number and total charge passing through the interface before failures resulting from cell shorting. In both cases, the noted improvements were accompanied by cell impedances that had increased substantially. The data reveal that both interlayers prevent the direct exposure of LPS to the metallic Li, and therefore eliminate the intrinsic LPS decomposition that occurs at Li surfaces before electrochemical cycling. After cycling, a reduction of LPS to Li2S occurs at the interface when a Si interlayer is present; LiAlO, which functions to drop the potential between Li and LPS, suppresses LPS decomposition processes. The relative propensities towards SE decomposition follows from the electrochemical potentials at the interface which are dictated by the identities of the interlayer materials. This work provides new insights into the phase dynamics associated with specific choices for SE/electrode interlayer materials and the requirements they impose for realizing high efficiency, long lasting all-solid-sate batteries. Introduction All-solid state Li-ion batteries are receiving substantial attention due to their potential advantages in safety and reduced balance-of-plant requirements.1-2 While extensive work has been performed to improve the Li+ conductivity of solid-state electrolyte (SE) materials, 3-10 challenges still exist for using them in full batteries, among which the development of battery shorts and slow Li+ kinetics are two major hurdles. Battery shorting is most frequently initiated by Li dendrite formation and metallic Li infiltration through the electrolyte. Modeling suggest that compared to liquid electrolytes, SEs exhibit a high shear modulus sufficient to suppress dendrite growth. 11-12 More recent reports argue that Li preferentially infiltrates through pre-existing defects at the SE/Li interface and this infiltration drives crack propagation.13-14 A current density threshold, defined as the onset current density prior to short circuit development, was found critical in driving Li dendrite growth and crack propagation. 13, 15 More generally, the properties of the SE/electrode interface such as physical contact, interface morphology, and resistivity also contribute to define the critical current density of a specific cell. 5 The slow kinetics (i.e. low charge and discharge currents) resulting from poor interfacial contact, and undesired (electro-) chemical decomposition at the SE/electrode interface, is another area of focus for all-solid-state batteries. In particular, thiophosphate-


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based SEs, which exhibit Li+ conductivities comparable to liquid electrolytes, are predicted to exhibit narrow electrochemical windows of stability16 and more notably found to decompose upon contact with metallic Li.17-19 For example, Li7P3S11 (LPS) and Li10GeP2S12 (LGPS) reduce to Li2S and Li3P when directly exposed to metallic Li.17-18 The -Li3PS4 material, another thiophosphate-based solid Li+ conductor, also reduces to Li4P2S6 and Li2S at negative potentials (-0.1 V vs. Li/Li+).19 These interfacial decomposition products are poor Li+ conductors. The ionic conductivity of Li2S, for example, is ~10-12 mS/cm and for Li4P2S6, ~10-7 mS/cm. These values are 7-12 orders of magnitude lower than that of the thiophosphate SE. The formation of these poor Li+ conductors at the SE/electrode interface likely results in a cell impedance increase which probably underpins the slow battery kinetics seen in test devices. In addition, these Li+-blocking materials are intrinsically quite heterogeneous in form and likely non-uniformly distributed at the interface where they are formed, leading to local current spikes exceeding critical current densities, and in consequence battery shorting. 13, 15 Introducing a Li+ conducting layer can address the challenges associated with the SE/electrode interface.2, 20-23 Ideally, this interlayer could protect the SE from exposure to low potentials, modify the SE surface morphology, and tailor the physical contact between the SE and electrode materials. Metal coatings such as Mg, Al, and Si have been applied to the Li7La3Zr2O12 (LLZO)/Li electrode interface previously.20,

23

These metal coatings

alloy with Li and enhance Li wettability at the LLZO surface, leading to reduced cell impedance. For the more Li+-conductive yet less stable thiophosphate-based SEs, a few interfacial materials have been reported, including Li-metal alloys,3 liquid wetting layers,2425

and Au.26 While these interlayers improve battery cycle life to some extent, the impacts

of these interlayer materials on critical factors associated with battery performance — such as interfacial chemistry, morphology, current density, Li dendrite formation, and battery shorting — remain problematic and the chemical speciations associated with them at the anode surface are incompletely understood. In this work, we address the essential properties an ideal interlayer should exhibit by evaluating and comparing the effect of two interlayer materials, Si and Al2O3, on the SE/Li interface via fundamental modes of physical and spectroscopic characterizations. LPS is chosen as an exemplar of the thiophosphate SE for its higher Li+ conductivity compared to that of -Li3PS4, and better structural stability compared to LGPS. We show that while the bare LPS/Li interface is unstable, the presence of either a Si or Al2O3



interlayer engenders additional stability, albeit at the cost of additional interfacial impedance. Degradation products are still found following cycling with the Si interlayer, while fewer such products are found with the Al2O3. We show that the origin of the different behaviors relates to the identity of lithiated components present at the interface and the potentials experienced by the SE during cycling. In this work, we also provide new insights into how the progression of the interfacial conversion reaction demonstrate important details of the potential profiles present at the interface. Experimental Section Materials and All-Solid Cell Preparation. Li7P3S11 (99.5%) (LPS) was purchased from MSE Supplies (Tucson, AZ, USA) and used without further purification. LPS SE pellets (0.6 mm thick and 13 mm in diameter) were obtained by pressing 0.15 g of LPS powder under 3000 psi in an Ar atmosphere. Electrochemical measurements were performed on three types of symmetric cells. Li/LPS/Li cells were assembled by placing the LPS pellet between two pieces of Li foil. For the Li/AuSi/LPS/SiAu/Li cell, 20 nm of Si and a 30 nm of Au were sputtered (AJA Orion 3) sequentially on both sides of the LPS pellet. In order to prevent post-sputtering air exposure-induced Si oxidation, we encapsulated the Si layer with an additional 30 nm of sputtered Au (this Au film does not block Li transport as it forms a Au-Li alloy during Li deposition). Au functions as a Litransport permissive hermetic layer that also provides electrical contact. The AuSi-coated LPS pellet was placed between two pieces of Li foil. The Al2O3 film was coated onto Li foil via atomic layer deposition (ALD) as previously described.27 Al2O3, exhibiting low Li+ conductivity ( = 5.0 10-7 mS/cm) deposited onto Li likely converts to a more Li+ conductive material such as LiXAl(2-x/3)O3 ( = 6.0 10-5 mS/cm).27-28 An Al2O3 thickness of 10 nm was chosen because 5 nm films were found to be insufficiently stable while 20 nm films were found to be too resistive. The film thickness was verified by X-ray fluorescence (XRF) measurements. Two pieces of LiXAl(2-x/3)O3 (LiAlO) coated Li foil were used as the electrodes for the Li/LiAlO/LPS/LiAlO/Li cell. Scheme 1. Illustrates the structure of the symmetric cells in the presence of interlayers. Li/LPS/Au (or Li/LPS/Si/Au) samples used for in-situ Raman experiments were prepared by sputtering 50 nm of Au (or 20 nm Si and 30 nm Au sequentially) on one side of the LPS pellet and placing Li on the other side. The Au film serves as the working electrode and Li serves as the counter/reference electrode. Au was chosen due to its chemical stability, excellent conductivity, and compatibility with the Raman experiment.


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Li

Si

LPS

Au*

Li

LPS

LiAlO

Li

Li

A

B

*Li permissive hermetic layer and electrical contact.

Scheme 1. The structures of the symmetric electrochemical cells in the presence of (A) SiAu interlayer and (B) LiAlO interlayer. Au is used as a Li-transport permissive hermetic layer to prevent the oxidation of Si during cell assembly and at the same time serves as an electrical contact. Electrochemistry. All electrochemical measurements were performed in a Swagelok cell. A spring was used to ensure sufficient contact between electrodes and solid-state electrolyte. The external pressure of the cell was in the range of 36 to 43 psi. Relative to a Li counter/reference electrode, the potential of the other Li electrode was swept between -0.1 V and +0.1 V. Electrochemical impedance spectroscopy (EIS) was measured along with CV cycles at frequencies from 1 MHz to 5 mHz with an amplitude of 30 mV, using a SP-150 potentiostat/galvanostat (BioLogic Science Instruments) Raman Spectroscopy. Raman spectra presented in this paper were measured using a custom-built spectro-electrochemical cell described in our previous work.19 The cell was assembled in an Ar-filled glove box prior to measurements. A 532 nm laser (B&W Tek Inc.) was used as the excitation source. Raman spectra in a range between 170 cm1

and 2300 cm-1 were collected by a Shamrock SR-303i monochromator (Andor Tech)

coupled to an iDus 420 spectroscopy CCD detector (Andor Tech). Each spectrum was a co-addition of 60 spectra, each with a 2 second integration time. X-ray Photoelectron Spectroscopy (XPS).

The XPS data were measured using a

monochromatic Al K source on Kratos Axis Ultra spectrophotometer (Kratos Analytical) operating, for acquisition of high resolution core level spectra, with an analytical energy resolution of 0.4 eV. The binding energies were calibrated relative to the C 1s peak at 284.6 eV in order to eliminate the influence of surface charging effects. The S 2p peak was fitted with a Gaussian line shape in CasaXPS software (Casa Software Ltd.) with 1.2 eV energy splitting between the S2p 1/2 and 3/2 peaks. Scanning Electron Microscopy (SEM). SEM images were collected on a JEOL, JSM-7000F (GenTech) at 15 keV. For LPS at the LPS/SiAu interface, SEM measurements were performed on the SiAu coated LPS surface following a 48 h Li contact, after which



the Li foil was peeled from the interface to expose the interlayer. Measurements were also performed on a SiAu coated LPS surface recovered from a shorted Li/AuSi/LPS/SiAu/Li cell after cycling, after which the Li foil was peeled from the interface to expose the interlayer. For LPS at the LPS/LiAlO/Li interface, SEM images were recorded at the LPS surface pressed against LiAlO coated Li, after which LPS was peeled from the interface to expose the interlayer. Measurements were also performed on a LPS surface recovered from a cycled and shorted Li/LiAlO/LPS/LiAlO/Li cell, after which LPS was peeled from the interface to expose the interlayer. Transmission Electron Microscopy (TEM). A FEI Helios Nanolab 600 (ThermoFisher) was used to prepare cross-sectioned samples via standard focused ion beam (FIB) techniques. The cross-sections were studied with TEM and Scanning TEM (STEM) in order to observe sample morphologies and elemental distributions before and after battery cycling. A JEOL 2100 TEM/STEM with 200 kV operating voltage was used for imaging, energy dispersive X-ray spectroscopy (EDS), and electron energy loss spectroscopy (EELS). Additional EDS characterizations were carried out using a JEOLARM300F Grand ARM TEM/STEM operating at 300 kV. Detailed EDS analyses and mapping were processed using Oxford AZtec EDS software. In particular, spectral overlap between the Si K and Au M EDS peaks was deconvoluted using the AZtec TrueMap functionality to produce EDS maps of Si and Au that more accurately represent the elemental distribution present within the sample. Results and Discussion

B

A

300 0 -300 -600 -0.10

Cycle 1 Cycle 2 Cycle 3 Cycle 4

-0.05 0.00 0.05 Potential (V vs. Li/Li+)

0.10

RLPS

Rint

CPELPS

CPEint

Cycle 3

-Im (Z) / 

600

1000 Hz

50 

700 kHz

Current Density (A/cm2)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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Cycle 2 Cycle 1 Pre-cycle

100

200

300

Re (Z) / 

400

Figure 1. (A) Cyclic voltammetry (CV) of Li/LPS/Li cell obtained at 0.2 mV/sec; and (B) The Nyquist plot of the same cell before (red open circle) and after the first three CV


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cycles (red, blue, and green solid circle) along with the EIS spectrum of LPS between two aluminum blocking electrodes (grey open circle).

Fig. 1A shows the first 4 CV cycles obtained from a Li/LPS/Li symmetric cell. The potential was scanned between -0.1 V and +0.1 V at a scan rate of 0.2 mV/sec. The CV scans were initiated from 0 V (OCP) towards negative potentials. Fig. 1A shows that the CV evolves with cycle number. Initially, the CV shows the presence of deposition and stripping behavior. However, at cycles 3 and 4, “noise” features, i.e random and drastic changes of current density during potential sweeps, were observed, likely due to the formation and breakdown of Li dendrites which function as electron conductive channels (vide infra). As the CV evolves, the overall current density grows as a function of cycle number. (The first 6 CV cycles are provided in Fig. S1) At cycle 15, the current density was found to increase markedly (> 10 A/cm2) and presents as spikes in the CV; this behavior is consistent with the presence of a short across the electrolyte. These phenomena suggest Li dendrites form, grow, and eventually interconnect the two electrodes. Starting with cycle 2, and following through cycles 3 and 4, we also observed a decrease in the current density reached at the scan limits when the potential was cycled in a range lying between -0.1 V and +0.1 V. This decrease likely results from Li+ depletion layers that form near -0.1 V and +0.1 V. Such a cause for a depletion behavior directly implicates the formation and accumulation of a Li+ blocking layer near the electrode surface due to SE decomposition.16, 18-19 This Li+ blocking layer at the LPS/Li interface likely inhibits Li+ diffusion across the interface. The impedance of the same Li/LPS/Li symmetric cell was monitored before and after the CV cycles. Fig. 1B shows the Nyquist plot obtained from the Li/LPS/Li cell after the first three cycles, comparing it to that of a cell before cycling (Fig. S2A shows the Bode plot). The Nyquist plot of LPS assembled between two Al (ion blocking) electrodes is also shown in Fig. 1B (grey open circle) as a reference. Qualitatively, as soon as the LPS is placed in contact with Li, significant growth of the overall cell impedance was observed relative to that obtained between the Al electrodes. This growth indicates that Li metal facilitates the formation of a high impedance layer at the Li/LPS interface. For the first 2 CV cycles, the cell impedance increases as a function of cycle number, which is consistent with the Li+ depletion layer formation inferred during the first 2 CV cycles shown in Fig.



1A. After cycle 3, a decrease in cell impedance is observed, which is likely a consequence of Li dendrite formation. This observation is consistent with the presence of the “noise feature” in the third cycle CV results. Starting from cycle 4, as the Li dendrites grow, the Li/LPS/Li cell yields real and imaginary phase resistivity equal to zero with no frequency dependence; this behavior is consistent with a shorted cell. In the low frequency range of the EIS for the first three CV cycles, the real phase resistivity decreases as a function of frequency while the imaginary phase is close to zero. Similar low frequency artifacts in EIS spectra are seen in other measurements where electrode surfaces are polarized for an extended period of time. The origin of these artifacts are ascribed to: (1) formation of intermediate species during electrode polarization;29 (2) significant ion adsorption and slow charge (ion) relaxation;30 and (3) preferential electrodeposition on one electrode inducing geometrical and/or electrical misalignment.31-32 In any event, the low frequency behavior observed in Fig. 1B likely suggest the presence of irreversible reaction chemistry. We quantified the impedance of Li/LPS/Li cell to better define the evolution of interfacial impedance during Li deposition and stripping processes. The high frequency semi-circles of the Nyquist plots were fitted into a two-component equivalent circuit (inset in Fig. 1B, dash lines show the fitted results), representing two interfacial components existing in the Li/LPS/Li symmetric cell: (1), a bulk electrolyte resistance (RLPS) in parallel with a constant phase element (CPELPS) originating from the grain boundaries inherently existing in LPS; and (2), an “apparent interfacial component” represented by an interfacial resistance (Rint) in parallel with and an interfacial constant phase element (CPEint).18-19 The latter interfacial component reflects a combined contribution from the electrode/electrolyte interface near both the anode and the cathode. The RLPS and CPELPS were determined from a separate EIS experiment where the same LPS pellet was placed in between two Al blocking electrodes (grey open circles in Fig. 1B). The conductivity obtained from the Al/LPS/Al cell was 1.1 mS/cm, consistent with previous reports.33 The values of Rint and CPEint representing the interfacial components were obtained and are provided in Table S1. An interfacial resistance of 107  is found once LPS was assembled between two Li electrode. The Rint value grows to 231 and 295  during the first two CV cycles before dropping to 221  after cycle 3 (due to the initiation of Li dendrite formation).


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900 Hz

700 kHz

4 30

A

15 0

Cycle 1 Cycle 2 Cycle 3 Cycle 4 Cycle 20 Cycle 30

-15 -30 -0.10

-0.05

0.00

0.05

Potential (V vs. Li/Li+)

0.10

-Im (Z) / k


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


0.5 k

Page 9 of 25

B

Cycle 30 Cycle 20 Cycle 4 Cycle 3

3 2

Cycle 2 Cycle 1

1 0 0

RLPS

Rint

CPELPS

CPEint

1 2 3 Re (Z) / k

Pre-cycle

4

Figure 2. (A) Cyclic voltammetry (CV) of Li/AuSi_LPS_SiAu/Li cell obtained at 0.2 mV/sec; and (B) the Nyquist plot of the same cell before (red open circle) and after CV cycles

In order to prevent LPS material from direct contacting metallic Li, we applied a thin Si interlayer, encapsulated by Au, between LPS and the Li electrode. Fig. 2A shows the first four cycles along with the 20th and 30th CV cycles scanned between -0.1 V and +0.1 V, at a scan rate of 0.2 mV/sec. The figure shows that the I-V behavior was relatively constant as a function of cycle number. In contrast with the Li/LPS/Li cell, the Li/AuSi/LPS/SiAu/Li cell ran at least 30 scans without shorting and no obvious Li+ depletion phenomena were observed. These observations suggest that the AuSi interlayer effectively prevented both interfacial decomposition and Li dendrite formation. Fig. 2A shows, however, that the current density values near -0.1 V and +0.1V are approximately one order of magnitude lower compared to that seen in the Li/LPS/Li cell (Fig. 1A). Thus, the AuSi interlayer likely results in a cell impedance increase. The impedance of the same Li/AuSi/LPS/SiAu/Li cell was monitored after each CV cycle. Fig. 2B shows the Nyquist plots obtained following cycles 1-4 and cycles 20 and 30. (Fig. S2B shows the Bode plot.) The figure shows that no significant impedance change occurs during the CV cycles. Quantitatively, the EIS spectra were fit into the twocomponent circuit model described above (dashed lines overlaid on the Nyquist plots show the fitted results). The RSE and CPESE values associated with the bulk LPS were fixed and the interfacial components were obtained from the fit. Table S1 provides these



Rint and CPEint values. In the presence of the AuSi interlayer, the interfacial resistivity is found to be near 3000 , an order of magnitude larger than the interface absent the AuSi. The impedance increase in the presence of the AuSi interlayer suggests that the amount of charge passed per cycle in the presence of the interlayer will be smaller than that passed in its absence. In the Li/LPS/Li cell during the first 2 CV cycles, the total Li deposition charge is found to be 0.49 C, with an accompanying stripping charge of 0.38 C. With the presence of the AuSi interlayer, we were able to run at least 60 CV cycles, (Fig. S3), with the total charge reaching 0.65 C during Li deposition, and 0.64 C during Li stripping. The electrode surface area of the cell with the AuSi interlayer was 0.71 cm2, which is only 56% of the area for the cell without the interlayer (1.27 cm2). Nonetheless, the cell with the AuSi interlayer evinces no evidence of dendrite formation even though substantially more Li+ has passed through it. RLPS

Rint

CPELPS

CPEint

B

60

A

30 0

Cycle 1 Cycle2 Cycle 3 Cycle 4 Cycle 20 Cycle 30

-30 -60 -90 -0.10

Cycle 30

Cycle 20

-0.05

0.00

0.05

-Im(Z)/k

90

400 Hz

0.5 k

700 kHz


1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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Cycle 4

Cycle 3 Cycle 2

0.10 Cycle 1

+

Potential (V vs. Li/Li )

Pre-cycle

0.0 0.5 1.0 1.5

Re(Z)/k

Figure 3. (A) Cyclic voltammetry (CV) of Li/LiAlO/LPS/LiAlO/Li cell obtained at 0.2 mV/sec; and (B) the Nyquist plot of the same cell before (red open circle) and after CV cycles. Another interlayer examined comprised an LiAlO thin film formed by ALD. The data in Fig. 3A show the first 4 cycles along with the 20th and 30th CV cycles scanned between a potential of -0.1 V to +0.1 V at a rate of 0.2 mV/sec. These CV scans show the presence of well-behaved deposition and stripping behaviors, suggesting that no Li


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dendrites were formed to short the cell. A gradual decrease in current density was observed during the first four CV cycles, a value that was found to stabilize after cycle 20. The electrochemical data presented in Fig. 3A show that near –0.06 V and +0.06 V, current depletion occurs in this case, which is likely in consequence of Li+ depletion in the (less Li+ conductive) LiAlO interlayer (for example LiAlO2 = 6.0 10-5 mS/cm, LPS = 1.1 mS/cm). This suggests that the Li deposition (stripping) current is likely limited by Li+ diffusion within the LiAlO interlayer. During the first 30 CV cycles, the total Li deposition charge was -1.25 C, and stripping charge was 1.15 C, which is substantially larger than that found with an interlayerfree cell. The average Columbic efficiency of the Li/LiAlO/LPS/LiAlO/Li cell was 91.5%, slightly smaller than the Li/AuSi/LPS/SiAu/Li cell (98.43%), which is likely a result of the different interfacial chemistry. In order to compare the total amount of charge passed through LiAlO and AuSi interlayers before cell shorting, CV scans at a slower scan rate (0.02 mV/sec) were carried out using symmetric cells assembled with each of the two types of materials. For the cell with SiAu interlayer, a dramatic current density increase (to ~100 mA/cm2) was observed at cycle 2, indicating the presence of a shorted cell (Fig. S4). For the cell with the LiAlO interlayer, current spikes were observed at cycle 19 followed by a current density increase to ~19 mA/cm2 at cycle 20, again suggesting the presence of a shorted cell (Fig. S5). The total Li+ deposition and stripping charge for the Li/LiAlO/LPS/LiAlO/Li cell during the first 18 CV cycles was 2.93 C and 2.75 C respectively, while only 0.22 C of Li+ deposition charge and 0.29 C of Li+ stripping charge were observed prior to cell shorting in the Li/AuSi/LPS/SiAu/Li

cell.

These

data

demonstrate

persuasively

that

the

Li/LiAlO/LPS/LiAlO/Li cell can pass substantially more charge prior to shorting than the Li/AuSi/LPS/SiAu/Li cell (we note that the electrode area of the Li/AuSi/LPS/SiAu/Li cell is 56% of that for the Li/LiAlO/LPS/LiAlO/Li cell; nonetheless the latter can still pass substantially more charge through the same electrode surface area). From the above, we see that the LiAlO interlayer affords a more durable/stablyoperable cell, one with a current density is approximately a factor of 2 greater than that for the SiAu interlayer. This suggest an approximately doubled resistance of the SiAu layer versus that of the LiAlO layer. The Li+ conductivity of a Si layer is determined by an interfacial conductance related to the Li insertion/extraction from the Si film, and a Li+



diffusivity within the Si film.34 Considering both contributions, the Li+ conductivity of a single crystal Si layer is on the order of 10-5 – 10-4 mS/cm,34 close to that of a LiAlO2 film. (6105 mS/cm)28

Since the thickness of the LiAlO interlayer is half of that for the SiAu interlayer,

the approximately doubled resistance measured for SiAu is consistent with this analysis. In order to better understand the cell impedance changes seen during the CV cycles, we performed a series of EIS experiments. The data in Fig. 3B shows the Nyquist plots obtained from an EIS study of a Li/LiAlO/LPS/LiAlO/Li cell as a function of cycle number (Fig. S2C shows the Bode plots). The cell impedance increased after the first four CV cycles before stabilizing at cycle 20. We fit the impedance data with the same two-component circuit model described previously and, assuming constant values of RLPS and CPELPS for bulk LPS, extracted the interfacial resistance and capacitance. These Rint and CPEint values are shown in Table S1. We found the interfacial resistivity increased significantly (343  to 905 ) after the first CV followed by slower growth to 1235  after the 30th cycle. The increase in impedance with cycle number is consistent with the drop in current density seen in Fig. 3A.


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S 2p XPS

S 2p XPS

AuSi

LPS

LPS

Pristine

A

Li

Li

LPS

LPS

AuSi

48 hr press

B

F

Li

Li LPS

Li

after 10 cycles

LiAlO

Li

C

after 10 cycles

Li

Li

LPS

LPS

AuSi

Li

short cell

LiAlO

48 hr press

LPS

AuSi

E

Bare

D

168 166 164 162 160 158 BE (eV)

Li short cell

G

LiAlO

H

168 166 164 162 160 158 BE (eV)

Figure 4. S 2p XPS spectra obtained at LPS surface before and after assembled/cycled in Li/LPS/Li symmetric cell in the presence of AuSi (A - D), and LiAlO (B - H) interlayers: (A) pristine LPS, (E) bare LPS, (B) and (F) after 48 hours press against Li, (C) and (G) after 10 CV cycles, (D) and (H) after cell shorts The CV and impedance results suggest that the presence of an interlayer stabilizes the solid electrolyte and prevents cracking and other decomposition that might serve to facilitate Li dendrite formation. In order to evaluate the stabilization mechanism(s) afforded by the interlayer, we collected XPS data to characterize the interfacial speciation present following cycling. These core level spectra were measured on disassembled cells from which one of the Li electrodes was removed, as shown schematically in Fig 4. In the case of the AuSi interlayer, the XPS survey scans revealed the presence of Au on the LPS side, but none on the Li side. Alternatively, for the case of the LiAlO interlayer, the survey scans showed the presence of Al on the Li side, with none found on the LPS side. This result suggests that the LiAlO interlayer is strongly adherent to Li, while Au more strongly associates with the LPS. Fig. 4 shows the S 2p XPS spectrum of a pristine AuSi coated LPS surface (Fig. 4A) along with S 2p spectrum of this same material after different treatments that include:



48 hours contact with metallic Li (Fig. 4B); 10 CV cycles between -0.1 V and +0.1 V, at 0.2 mV/sec (Fig. 4C); and following cell shorting (Fig. 4D). The S region was chosen for XPS analysis because changes in sulfur oxidation state provide a good indicator of interfacial evolution. 18-19, 35 The spectrum of the pristine AuSi coated LPS surface (Fig. 4A) can be fit to a predominate speciation involving two sulfur components: a major component exhibiting a lower binding energy at 161.4 eV (red), attributed to the terminal S (P-S-Li) in PS43- and P2S74-; and a minor component at 162.8 eV (blue), associated with the bridge S (P-S-P) in P2S74-.

18-19, 35-37

The peak fitting parameters and the numerical peak area percentages

calculated for all the spectra are given in Table S2. Fig. 4B shows that, once LPS contacts Li, a slight change in the peak ratio between the terminal and bridging S is observed. This behavior is in contrast to that reported for LPS in direct contact with Li, where increases in reduced S species are seen.18 This observation indicates the AuSi coating prevents LPS reduction by metallic Li. Following 10 CV cycles, the XPS data (Fig. 4C) show the presence of such reduced S species (seen at BE = 160.0 eV, green) in the S2p spectrum. This S component is associated with Li2S. 18-19, 35 Fig. 4D shows that following cell shorting, the intensity of the BE=160.0 eV component increases, suggesting the presence of additional reduced sulfur. Compared to the AuSi surface before cycling, we observe much higher relative intensities in the survey spectra of the S and P core level peaks (as compared to the Au 4f emission) on the AuSi surface after cycling. This observation suggest LPS migrates through (or within defects in) the AuSi interlayer and eventually covers it. Figs. 4E-H report the S2p XPS data obtained from the bare LPS surface (Fig. 4E) and those following 48 hrs of contact with Li (Fig. 4F); 10 CV cycles between -0.1 V to +0.1 V at a scan rate of 0.2 mV/sec (Fig. 4G); and following cell shorting (Fig. 4H). Fig. 4E shows that the bare LPS surface exhibits peaks associated with the terminal and bridging S species discussed above. Following Li contact (Fig. 4F) and 10 CV cycles (Fig. 4G), only slight changes in the terminal S to bridging S ratio are seen, and there is no evidence for further reduced S species. Following cell shorting (Fig. 4H), however, new features associated with the presence of Li2S are observed (green component) as discussed above.


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The XPS results suggest both AuSi and LiAlO can isolate LPS from direct contacts with Li and inhibit the subsequent reduction of LPS to products including Li2S at the Li surface. Following 10 CV cycles, Li2S was found at the LPS/AuSi surface. This result suggests that potentials sufficient to reduce LPS were present at the Si/LPS interface. It is known that Si lithiates at potentials of ca. +0.050 V vs. Li+/0, which suggests that the lithiated Si in the interlayer is very reducing. In contrast, no Li2S was found at the LPS surface proximal to the LiAlO coating Li following 10 cycles. This observation suggests that the LiAlO functions (in part) to drop the potential between Li metal and the LPS, protecting the LPS from reduction. When the cells short, Li dendrites grow between the two electrodes. Not surprisingly, the existence of these dendrites means that LPS is exposed to the reducing potentials attendant with metallic Li (Fig. S6 and Fig S7 show examples). Consequently, Li2S forms, as is seen in the post-mortem XPS from the shorted cells. 2 m

A

LPS/SiAu/Li before cycling

2 m

B

LPS/SiAu/Li shorted cell

2 m

C

LPS/LiAlO/Li before cycling

2 m

D

LPS/LiAlO/Li shorted cell

Figure 5. SEM image of AuSi covered LPS surface (A) pressed against Li, (B) recovered from a shorted Li/AuSi/LPS/SiAu/Li cell after cycling. SEM image of LPS surface (C) pressed against LiAlO coated Li, and (D) recovered from a shorted Li/LiAlO/LPS/LiAlO/Li cell after cycling. We further compared the morphology of the LPS surfaces cycled in the symmetric cells when the two types of interlayers were present. Figs. 5A and 5B report SEM images of the AuSi-coated LPS surface either pressed against Li without cycling (Fig. 5A) or recovered from a shorted Li/AuSi/LPS/SiAu/Li cell (Fig 5B). Before cycling, a generally smooth surface texture embedding a more sparse population of small (ca. 10 nm) particles was observed, consistent with the presence of a near continuous AuSi film



on the LPS surface. After cell shorting (Fig. 5B), the AuSi coated LPS surface exhibits a dramatic change in morphology—here showing the formation of larger particles ca. 500 nm in size. This morphology change likely results from the formation of Li2S as is revealed by the XPS results. Indeed, by examining the cross-section of LPS/SiAu interface with TEM, we found a physical destruction of the AuSi interface after CV cycling; detailed EDX analysis found Si to be intermixed with the LPS and its reduced products (Fig. S8 and Fig. S9). The morphology change could also be augmented by the volume change of the SiAu interlayer during lithiation. Figs. 5C and 5D show SEM images obtained from a LPS surface which was pressed against LiAlO coated Li without cycling (Fig. 5C), and recovered from a shorted Li/LiAlO/LPS/LiAlO/Li cell (Fig. 5D). The images in these figures show the presence of relatively smooth surfaces both before and after cell shorting, in addition to some cracks. This observation suggests that no substantial conversion chemistry occurs at the LPS surface when the LiAlO interlayer is present. The observation further suggests that the Li2S seen in the shorted cell is likely located in the regions of cracks and other imperfections, rather than being uniformly distributed across the electrode. The contrast between the AuSi and LiAlO interlayers suggests that LiAlO inhibits Li2S formation, likely because LiAlO functions to drop the potential between the Li electrode and LPS.

f 4.3 V

2

2.0V

c 2 2

-0.2V -0.1 V 0V 2.4 V (OCP)

200 300 400 500 600 700 Raman Shift (cm-1)

a

Au

LPS

Si

Li

f

b

2.0 V 1.0 V

c

0.8 V

2

0.4 V

2

-0.2 V

2

0V 2.0 V (OCP)

+ E sweep

2

12 hrs @

b

d e

B

Li

Relative Intensity (a.u.)

a

Au

LPS

- E sweep

d e

A Relative Intensity (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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200 300 400 500 600 700 Raman shift (cm-1)

Figure 6. In-situ Raman spectra obtained at Au electrode in (A) Li/LPS/Au cell, and (B) Li/LPS/SiAu cell


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The XPS results show the presence of Li2S as a consequence of cycling at the Li/AuSi/LPS interface. The formation of Li2S should be accompanied by other degradation products at the LPS interface. Fig. 6A shows a series of Raman spectra measured at a LPS/Au interface at various potentials. As shown in the black trace in Fig 6A, five LPS signature peaks are observed. Major peaks d (411 cm-1) and the shoulder e (426 cm-1) are associated to the symmetric stretch of PS4 in P2S74- and PS43- respectively, assignments consistent with previous work.4, 38-39 Peak a (178 cm-1) is assigned as the lattice deformation mode, and b (272 cm-1) is attributed to the δdef(S−P−S) deformation in PS43−. Peak f (610 cm-1) originates from the asymmetric stretch of PS43-. Complete peak assignments are provided in Table 1. At 0 V, the intensity of PS43- and P2S74- associated peaks d and e decrease along with a, b, and f, suggesting the decomposition of the original thiophosphate-based counter ion (PS43- and P2S74-). When more negative potentials are applied (-0.1 V, -0.2 V), a new peak c near 380 cm-1 emerges at the expense of peaks d and e. Peak c is associated with P2S64-, the formation of which was previously observed at the -Li3PS4/Au interface at Li deposition potentials.19 As the potential stepped to positive values, peak c diminishes in intensity while peaks d and e recover, indicating conversion of P2S64- to PS43- and P2S74-. Note that even at 2.0 V, the intensities of peaks d and e do not completely recover. The intensity of the five LPS signature peaks do eventually recover after a long time positive potential hold. This observation suggests that at the LPS/Au interface, Li4P2S6 forms at 0.1 V and persists over a wide range of potentials before converting back to PS43- and P2S74- at positive potentials. Li4P2S6 exhibits low Li+ conductivity (~10-7 mS/cm), and its presence likely functions to block Li+ from depositing at the electrode surface. Table 1. Vibrational Assignments of LPS/Au, and LPS/SiAu interface during Li deposition and stripping peak label a b c d e f

frequency (cm-1) before cycle during cycle 178 272 380 411 411 426 426 610 610

Assignment Lattice deformation δdef(S−P−S) in PS43− s(PS3) and (P−P) in P2S64− s(PS4) in P2S74− s(PS43-) as(PS43−)


refs 40-42 40, 42-44 45-46 4, 38-39 40, 42, 44-47 40, 42, 47-48


We performed a similar measurement at the LPS/SiAu surface to explore the influence of the Si interlayer on speciation during Li deposition and stripping. Fig. 6B shows the in-situ Raman spectra measured at the LPS/SiAu interface as a function of potential. The Raman spectra of the pristine SiAu-coated LPS surface are similar to that of the Au coated surface. The intensity of peak c associated with P2S64- increases while the PS43-/P2S74- related peaks (a, b, and d-f) decreases in intensity at 0 V and -0.2 V. When positive potentials are applied, peaks a, b, d-f grow in intensity while that of peak c decreases. These observations suggest that a similar PS43-/P2S74- to P2S64- conversion takes place at the LPS/SiAu interface. In this case, however, the PS43-/P2S74- to P2S64conversion occurs at a slightly less negative potential (0 V), and the reverse reaction (P2S64- to PS43-/P2S74-) takes place at lower limiting positive potentials. (~0.8 V) The complete spectral recovery (which occurs at 1.0 V) does not require a long positive potential hold, a behavior that is in marked contrast to the Au/LPS case. These results indicate that the presence of Si facilitates both the forward and backward reaction of the PS43-/P2S74- to P2S64- conversion. One likely origin of this more facile conversion behavior has to do with the Si complexes available at the interface. As the Si-Li alloy delithiates, it generates the much less electropositive elemental form of Si. By way of contrast, when Au is delithiated, it would do so (under conditions of thermodynamic control) via the decomposition of a Aurich intermetallic phase.

From the data, this latter speciation sustains a strongly

electropositive character within the region of interfacial contact. Fig. 7 shows a series of cartoons illustrating the proposed potential evolution at the LPS/electrode interface during Li deposition and stripping process. As shown in Fig. 7A, when a reducing potential (E0 < 0 V vs. Li/Li+) is applied at the Au electrode surface, Au in vicinity of LPS surface lithiates and forms a Li-Au alloy. In consequence, LPS at the LPS/Li-Au interface experiences reducing potentials (E1=E0 < 0V) causing the Li3PS4/Li4P2S7 to Li4P2S7 conversion. When a positive potential is applied at the Au electrode surface, the Li-Au alloy delithiates near the LPS/Au interface forming a Li-depleted layer, before complete delithiation (Fig. 7B). As Au is a good conductor, the potential (E1) at the LPS/Au interface is expected to be close to the potential of the Li-Au alloy (E2~0V vs. Li/Li+) and consequently Li2P2S6 persists. Following a potential hold at positive values for an extended period of time, the Au electrode delithiates completely, the potential at the LPS/Au interface approaches the potential at the Au electrode surface, (E0>0V vs. Li/Li+) and Li4P2S6 converts back to Li3PS4 and Li4P2S7 at the interface (Fig. 7C).


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Fig. 8 shows a similar cartoon illustrating the Si situation. Here the Si lithiates at the LPS/Si interface when a reducing potential (E0 < 0 V vs. Li/Li+) is applied, (Fig. 8A) and at these potentials Li3PS4/Li4P2S7 converts to Li4P2S6. Similarly, as Si near the interface delithiates, a Li-depleted layer forms near the interface region, while the Li-Si alloy probably still exists in the bulk Si electrode (Fig. 8B). The Si so produced need not form ohmic interfaces with the lithiated materials they contact—a heterogeneity that could underpin physicochemical mechanisms sustaining the significant interfacial heterogeneity of potential within the barrier layer region as is evidenced by the data. In contrast to the Au, the Li-depleted Si near the interface likely functions to introduce a potential difference between the Li-Au layer (E2) and the LPS/Si interface (E1). This implicitly suggests that the heterogeneity in the potential distribution present in the Si layer provides a less effective medium through which to sustain a reducing potential at LPS/Si interface, allowing a facile reconversion of Li4P2S6 back to Li3PS4/Li4P2S7 prior to the complete delithiation of the Si interlayer. Eventually, however, the Si is completely delithiated (Fig. 8C). The Raman results (Fig. 6) further demonstrate that the identity of the interlayer plays an essential role in controlling the electrochemical potential at the SE/electrode interface. The potential that the SE experiences at the SE/electrode interface dictates the interfacial chemistry and subsequent battery performance.

Li4P2S6 E0 < 0V

C

Li4P2S6 E0 > 0V (initial)

E1=E0

Au

E1=E2~0 V E0

LPS

Li-Au Au

E0

B

LPS

E1< 0V

Li-Au Au

A LPS

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


E0 > 0V (holding)

Figure 7. Cartoon presents electrochemical potential of the LPS/Au interface at various electrode potentials (E0). E0 represents negative potential in A, and positive potential in B and C. C shows the scenario when a positive potential is held at the electrode surface for an extended period of time. All potentials are reference to Li/Li+.



Li4P2S6 E0 < 0V

E0

C

E0 > 0V (initial)

LPS

E1>E2~0 V E0

Si

E1>E2~0 V

LPS

E0

B

Li-Si Si

E1< 0 V

Li-Si Si

A LPS

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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E0 > 0V (holding)

Figure 8. Cartoon presents electrochemical potential of the LPS/Si interface at various electrode potentials (E0). E0 represents negative potential in A, and positive potential in B and C. C shows the scenario when a positive potential is held at the electrode surface for an extended period of time. All potentials are reference to Li/Li+.

Conclusion In summary, we report the effects of two interlayer materials, Si and LiAlO, on the potential dependent speciation and evolution of the SE/Li interface. Cyclic voltammetry and impedance spectroscopy were employed to characterize the cycling performances of the Li/LPS/Li symmetric cells with and without interlayers in place. Both Si and LiAlO interlayer dramatically enhance the attainable cycle numbers and total charge passable through the LPS/Li interface, although at the same time substantially increase the cell impedances as compared to Li/LPS/Li cells in the absence of these interlayers. The LPS/Li, LPS/SiAu/Li, and LPS/LiAlO/Li interfaces were further characterized by XPS, SEM and in-situ Raman spectroscopy before and after electrochemical cycling. Taken together, the results suggest that, Li+ blocking materials such as Li4P2S6 and Li2S are formed at the interface at Li+ reducing potentials. The LiAlO interlayer functions to drop the potential between LPS and Li, and thus suppress LPS decomposition. In contrast, Si interlayers lithiate and efficiently maintain the strongly reducing potentials (and Li transport) that allow LPS reduction chemistry to occur at the interface. These results illustrate the impact of the interlayer materials on the electrochemical potential at the SE/electrode interface. The latter determines the interfacial decomposition chemistry which ultimately influence the battery performances. This current work provides two important observations related to the minimal requirements that an effect Li/SE interlayer must meet. The interlayer must facilitate a high rate of Li+ transport, and simultaneously function to drop the potential between Li and SE, thus preventing SE decomposition. Theoretical prediction of the electrochemical potential window where a particular SE


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material exhibits sufficient thermodynamic stability16 could serve as an appropriate

starting point to identify new interlayer materials. . The present results suggest alloying barrier layers may have diminished capabilities of stabilizing SE/Li interfaces in consequence of this latter requirement. Even so, any effective strategy irrespective of the class of materials selected will have to harbor capabilities for lowering the interfacial impedance. Designs that can do so will be crucial requirements for the development of new interlayer materials. Supporting Information (Fig. S1) CV results of Li/LPS/Li cell, 1-6 cycles; (Fig. S2) Bode plots obtained from the Nyquist plots shown in Fig 1-3; (Fig. S3) CV and EIS results of the Li/AiSi/LPS/SiAu/Li cell, 1-4 cycles and the 50th and 60th cycles; (Fig. S4) CV results of the Li/AiSi/LPS/SiAu/Li cell (Fig. S4), and the Li/LiAlO/LPS/LiAlO/Li cell; (Fig. S5) obtained at slow scan rate (0.02 mV/sec); (Fig. S6) SEM image of LiAlO covered Li before and after cycling; (Fig. S7) SEM images of bare Li before and after cycling (Fig. S8) TEM image and EDX maps at the LPS/SiAu/Li interface before and after cycling; (Fig. S9) EDX and EELS line scan of the LPS/SiAu interface; (Table S1) Interfacial Resistivity and Capacity upon cycling; (Table S2) Peak fitting results for XPS spectra; (Fig. 4A-D) measured from the LPS/SiAu/Li interface; (Table S3) Peak fitting results for XPS spectra (Fig. 4E-H) measured from the LPS/LiAlO/Li interface. Corresponding Authors Ralph G. Nuzzo Email: [email protected] Andrew A. Gewirth Email: [email protected]

Author Contributions The manuscript was written through contributions of all authors. All authors have given approval to the final version of the manuscript. L. S. conceived the operando Raman and electrochemistry measurements; K. L. B. performed SEM experiment; F. C. C. contributed on the TEM and STEM measurements; R.T.H. performed XPS measurements; L.S. performed the XPS data analysis. M. J. Y and C. L. fabricated LiAlO coated Li electrodes. Notes The authors declare no competing financial interest. Acknowledgements



This work was supported as part of the Center for Electrochemical Energy Science, an Energy Frontier Research Center funded by the U.S. Department of Energy, Office of Science, Basic Energy Sciences. This work was carried out in part in the Frederick Seitz Materials Research Laboratory Central Facilities, University of Illinois. K.L.B. acknowledges support from the National Science Foundation Graduate Research Fellowship Program under Grant Number DGE-1144245. This work also made use of the EPIC facility of Northwestern University’s NUANCE Center, which has received support from the Soft and Hybrid Nanotechnology Experimental (SHyNE) Resource (NSF ECCS1542205); the MRSEC program (NSF DMR-1720139) at the Materials Research Center; the International Institute for Nanotechnology (IIN); the Keck Foundation; and the State of Illinois, through the IIN. References 1. Kato, Y.; Hori, S.; Saito, T.; Suzuki, K.; Hirayama, M.; Mitsui, A.; Yonemura, M.; Iba, H.; Kanno, R., High-power all-solid-state batteries using sulfidesuperionic conductors. Nature Energy 2016, 1, 1-7. 2. Zhonghui, G.; Huabin, S.; Lin, F.; Fangliang, Y.; Yi, Z.; Wei, L.; Yunhui, H., Promises, Challenges, and Recent Progress of Inorganic Solid‐State Electrolytes for All‐Solid‐State Lithium Batteries. Adv. Mater. 2018, 30 (17), 1705702. 3. Kamaya, N.; Homma, K.; Yamakawa, Y.; Hirayama, M.; Kanno, R.; Yonemura, M.; Kamiyama, T.; Kato, Y.; Hama, S.; Kawamoto, K.; Mitsui, A., A lithium superionic conductor. Nat. Mater. 2011, 10 (9), 682-686. 4. Mizuno, F.; Hayashi, A.; Tadanaga, K.; Tatsumisago, M., New, highly ionconductive crystals precipitated from Li2S-P2S5 glasses. Adv. Mater. 2005, 17 (7), 918+. 5. Seino, Y.; Ota, T.; Takada, K.; Hayashi, A.; Tatsumisago, M., A sulphide lithium super ion conductor is superior to liquid ion conductors for use in rechargeable batteries. Energy Environ. Sci. 2014, 7 (2), 627-631. 6. Sahu, G.; Lin, Z.; Li, J. C.; Liu, Z. C.; Dudney, N.; Liang, C. D., Air-stable, highconduction solid electrolytes of arsenic-substituted Li4SnS4. Energy Environ. Sci. 2014, 7 (3), 1053-1058. 7. Ito, S.; Nakakita, M.; Aihara, Y.; Uehara, T.; Machida, N., A synthesis of crystalline Li7P3S11 solid electrolyte from 1,2-dimethoxyethane solvent. J. Power Sources 2014, 271, 342-345. 8. Kozen, A. C.; Pearse, A. J.; Lin, C. F.; Noked, M.; Rubloff, G. W., Atomic Layer Deposition of the Solid Electrolyte LiPON. Chem. Mater. 2015, 27 (15), 5324-5331. 9. Liu, Z.; Fu, W.; Payzant, E. A.; Yu, X.; Wu, Z.; Dudney, N. J.; Kiggans, J.; Hong, K.; Rondinone, A. J.; Liang, C., Anomalous High Ionic Conductivity of Nanoporous betaLi3PS4. J. Am. Chem. Soc. 2013, 135 (3), 975-978. 10. Bachman, J. C.; Muy, S.; Grimaud, A.; Chang, H.-H.; Pour, N.; Lux, S. F.; Paschos, O.; Maglia, F.; Lupart, S.; Lamp, P.; Giordano, L.; Shao-Horn, Y., Inorganic Solid-State Electrolytes for Lithium Batteries: Mechanisms and Properties Governing Ion Conduction. Chem. Rev. 2016, 116 (1), 140-162. 11. Monroe, C.; Newman, J., The Effect of Interfacial Deformation on Electrodeposition Kinetics. J. Electrochem. Soc. 2004, 151 (6), A880-A886. 12. Monroe, C.; Newman, J., The Impact of Elastic Deformation on Deposition Kinetics at Lithium/Polymer Interfaces. J. Electrochem. Soc. 2005, 152 (2), A396-A404.


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Understanding the Effect of Interlayers at the Thiophosphate Solid

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