Formation mechanism of Li7P3S11 solid electrolytes through liquid

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Formation mechanism of Li7P3S11 solid electrolytes through liquid phase synthesis Yuxing Wang, Dongping Lu, Mark Bowden, Patrick Z. El Khoury, Kee Sung Han, Zhiqun Daniel Deng, Jie Xiao, Ji-Guang Zhang, and Jun Liu Chem. Mater., Just Accepted Manuscript • Publication Date (Web): 03 Jan 2018 Downloaded from http://pubs.acs.org on January 3, 2018

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Chemistry of Materials

Formation mechanism of Li7P3S11 solid electrolytes through liquid phase synthesis Yuxing Wang a, Dongping Lu a, *, Mark Bowden b, Patrick Z El Khoury b, Kee Sung Han b, Zhiqun Daniel Deng a , Jie Xiao a, Ji-Guang Zhang a, Jun Liu a, * a

Energy and Environment Directorate, Pacific Northwest National Laboratory, Richland, WA 99352, USA

b

Environmental Molecular Sciences Laboratory, Pacific Northwest National Laboratory, Richland, WA 99354, USA Abstract Crystalline Li7P3S11 is a promising solid electrolyte for all solid state lithium/lithium ion batteries. A controllable liquid phase synthesis of Li7P3S11 is more desirable compared to conventional mechanochemical synthesis, but recent attempts suffer from reduced ionic conductivities. Here we elucidate the formation mechanism of crystalline Li7P3S11 synthesized in the liquid phase (acetonitrile, or ACN). We conclude that the crystalline Li7P3S11 forms through a two-step reaction: 1) formation of solid Li3PS4∙ACN and amorphous ‘Li2S∙P2S5’ phases in the liquid phase; 2) solid-state conversion of the two phases. The implication of this two-step reaction mechanism to the morphology control and the transport properties of liquid phase synthesized Li7P3S11 is identified and discussed. Introduction Energy storage devices are pivotal to vehicle electrification and renewable energy storage, which are key steps in order to meet decarbonisation targets around the world. Lithium/lithium ion batteries (LIBs) are the primary candidates especially in applications such as electric vehicles (EVs) and portable electronic devices where high energy densities are required. However, the high cost and inadequate energy density of state-of-the-art LIBs hinder further market penetration of EVs. It is widely accepted that further enhancement of energy density requires fundamental changes in battery chemistries.1-6 These changes may arise from electrode materials such as silicon anode,7 lithium metal anode,8, 9 Li-S,10 Li-air,11 etc., or from electrolytes.12 Solid-state electrolytes (SSE), in particular, inorganic lithium ion conductors, are generally considered safer and more thermally, chemically/electrochemically stable, although this assessment needs to be made on an individual basis.5, 13 Given the well-recognized issues with Li-metal,14 Li-S15 and Li-air batteries in liquid systems, all-solid-state design may be the solution to enable these systems.16 Other under-appreciated advantages of solid-state electrolyte systems include: 1) better thermal stability allows battery operation at elevated temperatures and simplification of thermal management, which is translated to enhanced pack-level energy/power densities;17 2) dimensional stability allows use of multiple electrolyte systems with distinct functionalities;18 3) singleion conduction nature and high carrier-density improves power performance at low temperature and ultra-high current densities.17 Despite many advantages, most solid-state electrolyte systems suffer from issues such as relatively low ionic conductivity, difficulty in forming and retaining intimate contact with electrode materials and high 1 ACS Paragon Plus Environment

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processing cost, etc..19, 20 Compared with oxides and phosphate systems, sulfide-based solid electrolytes in general have higher ionic conductivities and lower elastic modulus (softer), making them a more practical substitution for liquid electrolytes without much deviation from current battery manufacturing process.5, 20 It has been reported that some sulfide systems even have ionic conductivity exceeding that of liquid electrolytes.17, 21-23 Most notable examples are Li10GeP2S12 (12 mS cm-1)21 and Li7P3S11 (17 mS cm-1)22, 23 glass-ceramics. Li7P3S11 have been shown to form more kinetically stable interfaces with electrode materials than Li10GeP2S12.24, 25 Conventionally, the glass ceramics Li7P3S11 is synthesized by crystallization of 70Li2S∙30P2S5 glass,26 which may be prepared by mechanochemical synthesis27 or melt quenching method28, the former being preferred method. The mechanochemical synthesis involves planetary ball-milling of precursor powders under dry or wet conditions (no chemical interaction between liquid medium and the powder). Although effective, there are questions whether the process is scalable. Also, milled powders tend to aggregate into micron-size particles and an additional pulverization treatment may be necessary for use in composite electrodes.29 Liquid phase synthesis has been proven an effective method for synthesizing nanoparticles with controllable size and morphology.30 The richness of solution chemistry will endow great flexibility and tunability to the material preparation process. For instance, composites of SE and active materials or SE and conductive carbon can be obtained directly from the liquid phase.18, 31 From the manufacturing point of view, the liquid phase process is readily scalable and more compatible with conventional electrode preparation process such as composite cathode mixing and slurry coating. The liquid-phase synthesis of sulfide solid electrolytes can be further divided into two categories. In the first method, precursor powders are completely dissolved in organic solvents (methanol, ethanol, Nmethylformamide, Hydrazine, etc.)32-35 to form a homogeneous solution; in the second method, reaction between precursor powders are mediated by polar, aprotic organic solvents (tetrahydrofuran (THF), acetonitrile (ACN), 1,2-dimethoxyethane (DME), ethyl propionate, etc.) to generate precipitates.36-38 In both methods, final solid electrolyte products are obtained by drying off the solvent and subsequent heat-treatment for crystallization. While homogenous solutions are suitable for coating electrode particles, the precipitation method is conducive to producing small, uniform particles. Unfortunately, reported ionic conductivities of solid electrolytes using either method are lower than those obtained by mechanochemical or solid-state methods.18, 37, 39 Recently, various sulfide-based solid electrolytes (β-Li3PS4, Li7P3S11, Li7P2S8I, etc.) have been synthesized in the liquid phase.18, 36, 37, 40 Li7P3S11 is of particular interest due to its extraordinarily high conductivity in its crystalline form; crystalline Li7P3S11 has been synthesized in THF, acetonitrile and DME with a large variation in reported transport properties.18, 37, 39 It is unclear exactly how the Li7P3S11 crystalline phase forms from the deposited electrolyte precursor. In this research, we acquire further understanding of the formation mechanism of Li7P3S11 phase in acetonitrile by tracking the phase change of not only the precipitates, but also the dissolved phase in the supernatant liquids. We found that the deposited electrolyte precursor is actually a mixture of crystalline Li3PS4∙ACN and amorphous ‘Li2S∙P2S5∙ACN’, which then convert into crystalline Li7P3S11 through solid state reaction. Implications of this formation mechanism for the transport property and morphology of liquid phase synthesized Li7P3S11 solid electrolytes are discussed. 2 ACS Paragon Plus Environment

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Experimental Section Due to the extreme sensitivity of the sulfide compounds to moisture, all operations were carried out in an Ar-filled glovebox unless otherwise noted. The Li7P3S11 powder was synthesized from Li2S (Alfa Aesar, 99.9%) and P2S5 (Sigma Aldrich 99%) in acetonitrile (Selectilyte BASF, battery grade). All the raw materials were used without further treatment. Stoichiometric amount of Li2S and P2S5 were ground first in an agate mortar and poured into the acetonitrile solvent. The powder to solvent ratio is 1:20 g/ml. The mixture was heated at 50 °C on a hotplate and stirred for three days. A slightly greenish solution containing white precipitate was obtained. The solvent was then allowed to evaporate on a hotplate at 150 °C. The deposited precursor is denoted as DP. Alternatively, the solution mixture was sealed in a tube and centrifuged at 4400 rpm for 6 minutes. The powder precipitate and the supernatant were separated after decanting. Both were then dried under vacuum. A clear gel-like substance was obtained after solvent evaporation from the supernatant. The dried powder precipitate and the gel from the supernatant are denoted as PP and SP, respectively. The DP, PP and SP samples were then heat-treated at 200 °C or 260 °C in a sealed PTFE container filled with argon gas for 1 h. The sample denotation is summarized in Table 1. Table 1. Sample denotation and treatment conditions.

Sample ID DP

Treatment Deposited precursor from the precipitate and solution mixture after acetonitrile evaporation PP Powder precipitate After centrifuge and decanting SP Dried solution phase from the supernatant after acetonitrile evaporation DP-200 DP annealed at 200 °C for 1 h DP-260 DP annealed at 260 °C for 1 h * PP-200, PP260, SP-200, SP-260 are defined in the same manner

Thermogravimetric analysis (TGA) and differential thermal analysis (DTA) were performed on sample DP, PP and SP with Netzsch STA 449F1. Sample loading and weighing were carried out inside glovebox. The sample was transferred quickly into the TG instrument in a sealed container and the chamber was flushed promptly. The conditions were 5 °C heating from 25 °C to 300 °C under argon. The morphology of the samples was observed with a dual-focus ion beam (FIB) scanning electron microscope (SEM, Environmental, FEI Helios) at 5kV. Powder x-ray diffraction (PXRD) was used for phase characterization. Samples were sealed in thinwalled glass capillary tubes (500 µm diameter, 10 µm wall thickness, Charles Supper Co., MA) under argon. A Rigaku D/Max Rapid II micro-diffraction system with a rotating Cr target (λ = 2.2910 Å) operated at 35 kV and 25 mA was used to collect the diffraction patterns. A parallel X-ray beam collimated to 300 µm diameter was directed onto the specimen and the diffracted intensities were recorded on a large 2D image plate during a 10 min exposure. Alternatively, a desktop diffractometer (Rigaku MiniFlex II) was employed with a scan speed of 2° min-1 and a step size of 0.05°. The sample was covered by an 8-μm Kapton film during the measurement. 3 ACS Paragon Plus Environment


The Raman spectra were collected using a Raman spectrometer (Horiba LabRAM HR) coupled with an inverted optical microscope (Nikon Ti-E). The incident CW laser light source (633 nm) was attenuated using a variable neutral density filter wheel (to ~5 µW/µm2), reflected off a dichroic beamsplitter, and focused onto the sample using a 10X microscope objective. The backscattered light was collected through the same objective, transmitted though the beamsplitter cube, and dispersed through a 600 g/mm grating onto a CCD detector. Spectra were acquired as time series (10 sequentially recorded spectra, each of which was time-integrated for 5s) to ensure the integrity of our sample. As such, the final spectra shown are time-averaged and otherwise not subjected to further data analysis procedures. 31

P solid state magic angle spinning (MAS) nuclear magnetic resonance (NMR) spectra were obtained at the spinning speed of 20 kHz and at 295 K with a 3.2 mm HXY probe on a 600 MHz NMR spectrometer (Bruker, Germany). The spectra were obtained by the Fourier transformation of free induction decay after a single pulse excitation with a 90 degree pulse length of 4 µs and the repetition delay of 100 s. The 31P chemical shift (δ) was calibrated using 0 ppm of 85% H3PO4 as an external reference. Spinning sidebands occur at multiples of the spinning speed were determined by comparison between the spectra obtained with various spinning speed; 15, 20 and 23 kHz. Electrochemical impedance spectroscopy (EIS) was employed to characterize the transport properties of the samples. 100 mg of the sample powder was pelletized by cold pressing in a 10 mm diameter pressing die at 380 MPa. Indium electrodes were formed by pressing In foils onto both sides of the pellet at 130 MPa. The pellet was sandwiched by two stainless steel rods inside a Swagelok setup for impedance measurement. The measured temperature range was -40 °C to 100 °C. The low temperature testing was performed inside an environmental chamber using an electrochemical interface (Solartron 1287, Solartron Analytical) and a frequency response analyzer (Solartron 1260, Solartron Analytical); the high temperature testing was performed inside a heating oven using a Biologic potentiostat (VMP3). The frequency range was 1 MHz to 1 Hz. Results and Discussion

Figure 1. (a) Images of the deposited precursor (DP), the powder precipitate (PP) and the supernatant containing the solution phase. (b) TGA and DTA data of sample PP, DP and SP. 4 ACS Paragon Plus Environment

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White precipitates appear immediately after the mixture of Li2S and P2S5 (originally yellowish) is added into acetonitrile. The colorless solution turns bluish within a minute. The bluish color retains after the mixture is stirred at 50 °C for 3 days. The obtained DP sample (without centrifugation) appears slightly yellowish, whereas the PP sample (with centrifugation) appears entirely white (Figure 1a). Interestingly, the color of the supernatant changes from bluish to yellowish slowly after decanting. It is consistently found that the yield of DP and PP are about 0.7 g of 1.15 g, respectively; clear gel-like substance remains after the evaporation of acetonitrile. These evidences clearly suggest that part of Li2S and P2S5 precursors remain in the supernatant and the DP is actually a combination of PP and SP. According to TGA and DTA (Figure 1b), both sample PP and DP exhibit one major thermal event at around 200 °C. Close inspection reveals that the onset and peak temperatures of the DP sample are 10 °C higher than those of the PP sample (200 °C and 220 °C vs. 190 °C and 210 °C). Both samples have large weight loss during the event. Rangasamy et al. reported that Li2S and P2S5 at 3:1 molar ratio combine with acetonitrile to form an unknown crystalline phase, which the authors assigned as Li3PS4∙2ACN.40 For sample PP, the weight loss below 100 °C can be attributed to absorbed acetonitrile; minimal weight loss was observed between 100 °C and 190 °C and above 230 °C, suggesting that the complex phase is stable below 190 °C but decomposes and releases all acetonitrile above 190 °C in an endothermic reaction. This is in stark contrast to Li3PS4∙3THF (tetrahydrofuran) which decomposes at about 100 °C.36 However, based on the weight loss of the PP sample at 200 °C, we believe the formula of the acetonitrile complex should be Li3PS4∙ACN rather than Li3PS4∙2ACN. The TG and DTA curves of sample SP show a rather smooth continuous trend throughout the temperature range. Most features are too small to explain. Noticeably, there appears to be a small exothermic peak in the DTA curve, and a slight change of slope in the TG curve at 260 °C. Similar features can be seen in the TG and DTA curves of sample DP. The origin of these peaks will be discussed later. In addition, sample DP continued to lose weight after the major thermal event, similar to sample SP but to a lesser degree. The TGA and DTA results suggest sample DP has characteristics of PP and SP.

Figure 2. SEM images of sample PP-260 (a), DP-260 (b) and SP-260 (c). The morphology of PP-260, DP-260 and SP-260 were characterized by SEM (Figure 2). PP-260 and DP260 samples were in the powder form; the SP-260 was obtained by casting the supernatant onto an aluminum substrate followed by annealing. Sample PP-260 consists of submicron primary particles 5 ACS Paragon Plus Environment


(Figure 2a). Determination of the exact particle size is complicated by slight agglomeration of primary particles. Nevertheless, these particles are much more uniform and smaller compared to those synthesized via mechanochemical synthesis.29, 41 Sample SP-260 before annealing (Figure S1d) shows a featureless, complete coverage of an amorphous phase on the substrate; after annealing, the amorphous feature remains in most of the regions while cracks and particles appear (Figure 2c), probably due to shrinkage induced by acetonitrile evaporation. The morphology of DP-260 is similar to PP-260, but there seems to be more agglomeration or secondary amorphous phase binding the primary particles (Figure 2b). From the low magnification images (Figure S1), it can be seen that some primary particles agglomerate into very large secondary particles in the extreme case. In contrast, particles in sample PP-260 seems more dispersible. These results indicate the amorphous solution phase has a critical effect on the morphology of Li7P3S11. The PXRD patterns of sample PP (Figure S3) before annealing matches well with the crystalline phase ‘Li3PS4∙2ACN’;40 after annealing at 200 °C or 260 °C, the patterns (Figure 3) match well with the reported β-Li3PS4.42 Peaks corresponding to Li2S were observed in both patterns, indicating small amount of Li2S were present in both samples. The crystallinity of sample PP is higher with higher annealing temperature; small unknown peaks also disappear in sample PP-260. We can unambiguously conclude from the XRD and TGA/DTA data that the main phase of sample PP is Li3PS4∙ACN, which undergoes decomposition and phase transform into β-Li3PS4 at T > 190 °C. For sample DP, the XRD pattern before annealing is similar to PP (Figure S3), suggesting that the SP coverage on PP is indeed amorphous. After annealing, the pattern of DP-260 matches well with that of the high-conductivity Li7P3S11 phase whereas DP-200 is a combination of Li7P3S11, β-Li3PS4 and some unknown phase. It should be noted that Bragg peaks only reflect crystalline phases and amorphous phases could also be present in the sample. The completely different phases of DP-260 and PP-260 indicate the critical role of the solution phase (SP) in converting the powder precipitate (PP) into the desired Li7P3S11 phase. The conversion occurs at a slightly higher temperature (about 10 °C) than the conversion of sample PP into β-Li3PS4, as seen from TG and DTA, which explains the presence of both βLi3PS4 and Li7P3S11 phases in sample DP-200. As DP is basically PP particles coated by SP, two processes need to occur during the conversion: 1) decomposition of PP and formation of β-Li3PS4 while releasing ACN; 2) solid state reaction of SP with PP or PP derivatives. It is inconclusive whether the two processes occur concurrently or sequentially; the presence of β-Li3PS4 suggests the latter mechanism is more likely. It can also be inferred that due to the need of solid state diffusion and reaction, the formed crystalline Li7P3S11 may suffer from local inhomogeneity and non-stoichiometry. According to the SEM image (Figure 2b), the conversion did not consume the SP coating completely, as some amorphous coverage was still observed in sample DP-260.

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Figure 3. PXRD patterns of sample DP, PP and SP after heat treatment at 200 °C or 260 °C. The raw patterns and detailed pattern fitting can be found in Figure S2.



Figure 4. Raman (left) and 31P MAS-NMR (right) spectra of sample DP, PP and SP annealed at different temperatures. Asterisk (*) denotes a spinning sideband. PXRD is useful in the identification of crystalline phases, but amorphous components of the sample cannot be characterized. Raman and 31P MAS-NMR spectroscopy were employed as complementary tools to further understand the local structures of the samples. The Raman spectra of PP-200 and PP260 are similar (Figure 4). A single peak at 426 cm-1 can be assigned to the local structural unit of PS43tetrahedra, confirming that only Li3PS4 phase is present.43 The spectrum of sample DP-260 shows a major peak at 410 cm-1 corresponding to the local structural unit of P2S74- and a shoulder peak at 426 cm1 corresponding to PS43-, consistent with the crystalline structure of Li7P3S11 phase.27 The 426 cm-1 peak is stronger in sample DP-200. This is expected as residual β-Li3PS4 phase is present in the XRD pattern of DP-200. The Raman spectrum of SP-260 is not presented due to large fluorescent background, which hinders the determination of its local structure. The 31P MAS-NMR spectra of sample PP-260 and DP-260 (Figure 4) match well with the reported spectra of 75Li2S∙25P2S5 and 70Li2S∙30P2S5 glass ceramics.44 The two strong peaks at 91 ppm and 87 ppm can be assigned to the structural unit of P2S74- and PS43- in the crystalline phases, respectively.45 Seino et al. analyzed the degree of crystallization in the 70Li2S∙30P2S5 glass ceramics using 31P MAS-NMR.46 They showed that broad peaks related to P2S74- and PS43- unit in the amorphous component convolute with sharp peaks corresponding to P2S74- and PS43- in the crystalline component in poorly-crystalline samples. Minimal broadening at the base of sharp crystalline peaks was observed in the spectra of sample PP-260 8 ACS Paragon Plus Environment

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and DP-260, suggesting that both samples are mostly crystalline. Two small peaks at 109 ppm and 105 ppm in the spectrum of DP-260 can be assigned to the structural unit of P2S64- in the crystalline Li4P2S6 phase. The Li4P2S6 phase is known to arise from decomposition of glassy Li4P2S7 with release of elemental S.45 It is also reported that prolonged heating of Li7P3S11 will lead to formation of Li4P2S6 phase.47 It can be safely concluded from XRD, Raman and 31P NMR analysis that crystalline Li7P3S11 phase is formed by conversion of Li3PS4 or Li3PS4∙ACN and the solution phase through solid state reaction. The question remains what the solution phase is. To answer the question, we conducted similar synthesis but varying the molar ratio of Li2S : P2S5. Interestingly, the quantity of powder precipitates decreases with increasing P2S5 amount, e.g., at a ratio of 55:45, only 0.2 g of PP was obtained from a total of 1 g of precursors. The XRD pattern of PP at a ratio of 55:45 is similar to that at a ratio of 75:25 (Figure S3), also matching dried precursor to Li7P3S11 in Yao et al.’s report18 and the ‘Li3PS4∙2ACN’ phase in Rangasamy et al.’s report 40, indicating that all the powder precipitates are essentially the same compound. At a ratio of 1:1, Li2S and P2S5 dissolve quickly to form a clear yellowish solution. Indeed, Liang et al. has shown that the 1:1 Li2S:P2S5 solution can be used to coat Li2S particles.48, 49 Therefore, the solution phase is some amorphous form of ‘Li2S∙P2S5∙xACN’. The term ‘Li2S∙P2S5’ is used in a vague sense to refer to the phase with Li2S/P2S5 molar ratio of 1:1 that either dissolves in or precipitate out of the solution, since the solvated form of the solution phase or the structure of the amorphous solids after solvent removal were not directly determined. It should be noted that neither Li2S nor P2S5 is soluble in acetonitrile, so the dissolution of Li2S and P2S5 (1:1) is probably due to the formation of a new structural unit P2S62- which is soluble in acetonitrile. We believe that Li2S and P2S5 (in the molar ratio of x : 100-x, 50 < x < 75) participate in the reaction in acetonitrile as follows : 100 → 2 50 ∙ 275 ∙ When x < 50, soluble ‘Li2S∙P2S5’ forms and some P2S5 remains undissolved; when x > 75, Li3PS4∙ACN precipitate forms and some Li2S remains undissolved. In the case of 70Li2S∙30P2S5 under equilibrium conditions, 80 mol.% of Li3PS4∙ACN (PP) precipitates out while 20 mol.% of ‘Li2S∙P2S5’ dissolves in the solution.

Figure 5. Schematic illustration of the formation mechanism of Li7P3S11 synthesized in acetonitrile. After acetonitrile evaporation, the Li3PS4∙ACN precipitates are coated by the amorphous ‘Li2S∙P2S5’ phase and then converted to the high conductivity Li7P3S11 phase after annealing. The proposed formation 9 ACS Paragon Plus Environment


mechanism of Li7P3S11 is illustrated in Figure 5. This precipitation-solid state conversion reaction mechanism may be universal to the liquid phase synthesis of thiophosphate compounds and their derivatives, as the Li3PS4∙ACN phase may be the only stable precipitate. For instance, the Li7P2S8I phase has been synthesized by solid state reaction of liquid-phase-synthesized ‘Li3PS4∙2ACN’ and LiI.40 It follows that the morphology of these liquid-synthesized thiophosphate compounds is largely determined by that of the Li3PS4∙ACN precipitate. Therefore, the formation mechanism of the Li3PS4∙ACN precipitate is of critical importance, which has not been elucidated. Given the insoluble nature of Li2S and high solubility of ‘Li2S∙P2S5’, the formation of Li3PS4∙ACN most likely occurs through conversion of Li2S particles from outside to inside by the ‘Li2S∙P2S5’ solution. Under such assumption, the morphology of Li2S particles will determine the morphology of Li3PS4∙ACN precipitate. When the Li2S precursor particles are large, the conversion may never complete due to the passivation of the outer Li3PS4∙ACN, which can explain the presence Li2S in sample PP-200 and PP-260. With annealing of the precipitates and solution phase mixture, majority of the solution phase will be ‘assimilated’ during the solid state reaction but small amount may remain on the surface of Li7P3S11 particles. Therefore, what the solution phase transforms into at the annealing temperature will influence the transport property of DP-260. The PXRD patterns of SP-200 and SP-260 have large amorphous background, consistent with the SEM observation that the sample remains mostly amorphous. Bragg peaks in SP-200 can be attributed to P2S5 and peaks in SP-260 can be attributed to Li4P2S6 and Li2P2S6 phases.50, 51 The small exothermic peak in DTA data of sample SP can then be attributed to crystallization process of the amorphous phase. The 31P NMR spectrum of SP-260 shows much complexity. Most peaks are broad, confirming the amorphous nature. The peak at 109 ppm is associated with P2S64- in the Li4P2S6 phase; the peak at 83 ppm is due to glassy Li2S-P2S5.45 Two peaks of small chemical shift are unidentified. Nevertheless, none of the crystalline Li4P2S6, Li2P2S6, glassy Li2SP2S5 has good ionic conductivity,50-52 so the presence of the residual solution phase is expected to lower overall ionic conductivity of sample DP-260.


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Figure 6. Arrhenius plot for the total and intergrain ion conductivity of Li7P3S11 (260 °C, DP-260). The ionic conductivity of sample DP-260 was characterized by electrochemical impedance spectroscopy over the temperature range of -40 °C to 100 °C. The impedance plots are shown in Figure S4 and detailed fitting information can be found in the supplemental materials. At room temperature (22 °C) or above, the impedance plots show a straight line intercepting or trending towards the x-axis at low and intermediate frequencies whereas the bending at high frequencies is due to stray capacitance and stray inductance. The straight line represents contribution from the electrodes and the intercepts give the impedance of the materials. It cannot be determined from these data what components contribute to the impedance therefore we denote it as total impedance (Rt). At low temperatures (0 °C or below), a semicircle appeared whose left intercept did not go to zero. Apparently, there are at least two processes that significantly contribute to overall impedance. Literature references of low-temperature impedance measurement of Li7P3S11 solid electrolytes are scarce so it is difficult to determine the origin of the low frequency semicircle. We assume that the process at high frequency corresponds to intragrain (bulk) transport and the process at lower frequency corresponds to intergrain transport. The intergrain impedance may be due to secondary phases or simply grain boundary impedance. The intragrain resistance and the intergrain resistance were obtained by fitting and the ionic conductivities were calculated from the resistance and the sample geometry. The room temperature conductivity of sample DP-260 is 8.7 x 10-4 S cm-1. The activation energy was calculated based on the Arrhenius equation, σ = σ0/T exp(-Ea/kT). It can be seen (Figure 6) that the total ionic conductivity shows a good Arrhenius behavior in the entire 11 ACS Paragon Plus Environment


temperature range with an activation energy of 0.37 eV (36 kJ mol-1). This value is close to 70Li2S∙30P2S5 glass obtained by mechanical milling and much lower than 70Li2S∙30P2S5 glass-ceramics (0.18 eV). Seino et al. has shown that the activation energy of 70Li2S∙30P2S5 glass-ceramics from mechanical milling is roughly negatively correlated with the degree of crystallization. However, the observed high activation energy cannot be explained by this correlation since the XRD and NMR data indicated that sample DP260 is mostly crystalline. This observation highlights the key differences between Li7P3S11 solid electrolytes obtained via solid state route and liquid phase synthesis. Glass-ceramics obtained by solid state route arise from crystallization of a homogeneous single-phase glass, whereas liquid-phase synthesized samples arise from reaction of at least two phases. Notably, the reported activation energies of Li7P3S11 via liquid phase synthesis, despite all having the superionic crystal phase, have large variation, ranging from 23 kJ mol-1 to 38 kJ mol-1.18, 37, 39 This is not surprising considering the complexity of the formation mechanism in the liquid phase as elucidated by this study and hence how sensitive the transport properties can be to experimental conditions. Identifying the two-step reaction mechanism provides crucial clues in further improving the quality of Li7P3S11 from the liquid phase synthesis. Measures that could potentially improve transport property of liquid phase synthesized Li7P3S11 should focus on facilitating complete conversion of the two solid phases, for instance, 1) optimizing the solvent extraction process so the two phases are more uniformly mixed; 2) optimizing annealing process for better solid state conversion; 3) reducing the Li3PS4∙ACN particle size to shorten diffusion length, etc.. It is evident that the intergrain process has a much higher activation energy than that of the intragrain process. Although indeterminable from the impedance spectrum, the intergrain resistance at 22 °C and 50 °C can be calculated by extrapolation to be 10 Ω and 1 Ω, respectively; the intergrain resistance is negligible at high temperatures. Conclusions The formation mechanism of crystalline Li7P3S11 solid electrolytes synthesized in acetonitrile was revealed by analyzing the intermediate products during the reaction. It is found that the precursors (Li2S: P2S5 = 70 : 30 by mole) in acetonitrile form Li3PS4∙ACN (Li2S: P2S5 = 75 : 25) precipitates and soluble ‘Li2S∙P2S5’ (Li2S: P2S5 = 50 : 50). Unlike the β-Li3PS4 crystalline phase which forms directly from the decomposition of Li3PS4∙ACN precipitates, crystalline Li7P3S11 forms through the solid state reaction of the Li3PS4∙ACN precipitates and the amorphous ‘Li2S∙P2S5’ phase from the supernatant. The soluble species ‘Li2S∙P2S5’ appears as amorphous coverage on the precipitate particles after acetonitrile evaporation. The liquid phase synthesized Li7P3S11 has a total conductivity of 0.87 mS cm-1 at room temperature. The understanding of the formation mechanism provides clues to further optimize the transport properties and morphologies of liquid phase synthesized Li7P3S11 solid electrolytes. Author Information Corresponding Author *D.Lu. E-mail: [email protected]. *J.Liu. E-mail: [email protected]. 12 ACS Paragon Plus Environment

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ORCID Yuxing Wang: 0000-0002-7828-9399 Notes The authors declare no competing financial interest. Acknowledgment This work was supported by 1) the Energy Efficiency and Renewable Energy (EERE) Office of Vehicle Technologies of the U.S. Department of Energy (DOE) under Contract No. DEAC02-05CH11231 and DEAC02-98CH10886 for the Advanced Battery Materials Research (BMR) Program; 2) the U.S. DOE EERE Water Power Technologies Office and the U.S. Army Corps of Engineers Portland District. The SEM, solid-state NMR and Raman characterization were conducted in the William R. Wiley Environmental Molecular Sciences Laboratory (EMSL). PNNL is operated by Battelle for the DOE under Contract DEAC05-76RLO1830. Supporting Information Additional SEM images and XRD data; Impedance spectra and detailed fitting information References (1) Scrosati, B.; Garche, J., Lithium batteries: Status, prospects and future. J Power Sources 2010, 195, 2419-2430. (2) Armand, M.; Tarascon, J. M., Building better batteries. Nature 2008, 451, 652-657. (3) Goodenough, J. B.; Kim, Y., Challenges for Rechargeable Li Batteries. Chem Mater 2010, 22, 587-603. (4) Goodenough, J. B.; Kim, Y., Challenges for rechargeable batteries. J Power Sources 2011, 196, 66886694. (5) Jung, Y. S.; Oh, D. Y.; Nam, Y. J.; Park, K. H., Issues and Challenges for Bulk-Type All-Solid-State Rechargeable Lithium Batteries using Sulfide Solid Electrolytes. Isr J Chem 2015, 55, 472-485. (6) Scrosati, B.; Hassoun, J.; Sun, Y. K., Lithium-ion batteries. A look into the future. Energ Environ Sci 2011, 4, 3287-3295. (7) Ashuri, M.; He, Q. R.; Shaw, L. L., Silicon as a potential anode material for Li-ion batteries: where size, geometry and structure matter. Nanoscale 2016, 8, 74-103. (8) Xu, W.; Wang, J. L.; Ding, F.; Chen, X. L.; Nasybutin, E.; Zhang, Y. H.; Zhang, J. G., Lithium metal anodes for rechargeable batteries. Energ Environ Sci 2014, 7, 513-537. (9) Cheng, X. B.; Zhang, R.; Zhao, C. Z.; Zhang, Q., Toward Safe Lithium Metal Anode in Rechargeable Batteries: A Review. Chem Rev 2017, 117, 10403-10473. (10) Manthiram, A.; Fu, Y. Z.; Chung, S. H.; Zu, C. X.; Su, Y. S., Rechargeable Lithium-Sulfur Batteries. Chem Rev 2014, 114, 11751-11787. (11) Grande, L.; Paillard, E.; Hassoun, J.; Park, J. B.; Lee, Y. J.; Sun, Y. K.; Passerini, S.; Scrosati, B., The Lithium/Air Battery: Still an Emerging System or a Practical Reality? Adv Mater 2015, 27, 784-800. (12) Xu, K., Electrolytes and Interphases in Li-Ion Batteries and Beyond. Chem Rev 2014, 114, 1150311618.



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Formation mechanism of Li7P3S11 solid electrolytes through liquid

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