Mechanism of Formation of Li7P3S11 Solid Electrolytes through

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Cite This: Chem. Mater. 2018, 30, 990−997

Mechanism of Formation 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*,† †

Energy and Environment Directorate, Pacific Northwest National Laboratory, Richland, Washington 99354, United States Environmental Molecular Sciences Laboratory, Pacific Northwest National Laboratory, Richland, Washington 99354, United States

‡

S Supporting Information *

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 than conventional mechanochemical synthesis, but recent attempts suffer from reduced ionic conductivities. Here we elucidate the mechanism of formation of crystalline Li7P3S11 synthesized in the liquid phase [acetonitrile (ACN)]. We conclude that crystalline Li7P3S11 forms through a two-step reaction: (1) formation of solid Li3PS4·ACN and amorphous “Li2S·P2S5” phases in the liquid phase and (2) solid-state conversion of the two phases. The implication of this two-step reaction mechanism for morphology control and the transport properties of liquid phase synthesized Li7P3S11 is identified and discussed.

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INTRODUCTION Energy storage devices are pivotal to vehicle electrification and renewable energy storage, which are key steps for meeting decarbonization 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 stateof-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 a silicon anode,7 a lithium metal anode,8,9 Li−S,10 Li−air,11 etc., or from electrolytes.12 Solid-state electrolytes (SSEs), in particular, inorganic lithium ion conductors, are generally considered safer and more thermally and 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− S,15 and Li−air batteries in liquid systems, an all-solid-state design may be the solution to enable these systems.16 Other underappreciated advantages of solid-state electrolyte systems include the following. (1) Better thermal stability allows battery operation at elevated temperatures and simplification of thermal management, which is translated to enhanced packlevel energy/power densities.17 (2) Dimensional stability allows the use of multiple electrolyte systems with distinct functionalities.18 (3) The single-ion conduction nature and high carrier density improve power performance at low temperatures and ultrahigh current densities.17 © 2018 American Chemical Society

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, high processing costs, etc.19,20 Compared with oxides and phosphate systems, sulfide-based solid electrolytes in general have higher ionic conductivities and lower elastic moduli (softer), making them a more practical substitution for liquid electrolytes without much deviation from the 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 interfaces with electrode materials more kinetically stable than those of Li10GeP2S12.24,25 Conventionally, glass ceramic Li7P3S11 is synthesized by crystallization of 70Li2S·30P2S5 glass,26 which may be prepared by mechanochemical synthesis27 or the melt quenching method,28 the former being the preferred method. The mechanochemical synthesis involves planetary ball-milling of precursor powders under dry or wet conditions (no chemical interaction between the liquid medium and the powder). Although it is effective, there are questions about whether the process is scalable. Also, milled powders tend to aggregate into micrometer-sized particles, and an Received: November 18, 2017 Revised: January 2, 2018 Published: January 3, 2018 990

DOI: 10.1021/acs.chemmater.7b04842 Chem. Mater. 2018, 30, 990−997

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Chemistry of Materials

substance was obtained after the solvent had evaporated 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 or 260 °C in a sealed PTFE container filled with argon gas for 1 h. The sample notation is summarized in Table 1.

additional pulverization treatment may be necessary for use in composite electrodes.29 Liquid phase synthesis has proven to be 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 a manufacturing point of view, the liquid phase process is readily scalable and more compatible with a 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, reactions between precursor powders are mediated by polar, aprotic organic solvents [tetrahydrofuran (THF), acetonitrile (ACN), 1,2dimethoxyethane (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 homogeneous 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 because of its extraordinarily high conductivity in its crystalline form; crystalline Li7P3S11 has been synthesized in THF, acetonitrile, and DME with a large variation in the 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 improve our understanding of the mechanism of formation of the 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.

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Table 1. Samples and Treatment Conditions samplea DP PP SP DP-200 DP-260 a

treatment deposited precursor from the precipitate and solution mixture after acetonitrile evaporation powder precipitate after centrifugation and decanting dried solution phase from the supernatant after acetonitrile evaporation DP annealed at 200 °C for 1 h DP annealed at 260 °C for 1 h

PP-200, PP-260, SP-200, and SP-260 are defined in the same manner.

Thermogravimetric analysis (TGA) and differential thermal analysis (DTA) were performed on samples DP, PP, and SP with a Netzsch STA 449F1 instrument. Sample loading and weighing were performed inside glovebox. The sample was transferred quickly into the TG instrument in a sealed container, and the chamber was flushed promptly. The samples were heated from 25 to 300 °C under argon. The morphology of the samples was observed with a dual-focus ion beam (FIB) scanning electron microscope (Environmental, FEI Helios) at 5 kV. Powder X-ray diffraction (PXRD) was used for phase characterization. Samples were sealed in thin-walled glass capillary tubes (500 μm diameter, 10 μm wall thickness, Charles Supper Co.) under argon. A Rigaku D/Max Rapid II microdiffraction 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 a 300 μm diameter was directed onto the specimen, and the diffracted intensities were recorded on a large two-dimensional 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 with an 8 μm Kapton film during the measurement. 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 beam splitter, and focused onto the sample using a 10× microscope objective. The back-scattered light was collected through the same objective, transmitted though the beam splitter 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 5 s) to ensure the integrity of our sample. As such, the final spectra shown are timeaveraged 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 a spinning speed of 20 kHz and 295 K with a 3.2 mm HXY probe on a 600 MHz NMR spectrometer (Bruker). The spectra were obtained by the Fourier transformation of free induction decay after a single-pulse excitation with a 90° pulse length of 4 μs and a repetition delay of 100 s. The 31P chemical shift (δ) was calibrated using 0 ppm of 85% H3PO4 as an external reference. Spinning sidebands at multiples of the spinning speed were determined by comparison between the spectra obtained at various spinning speeds (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

EXPERIMENTAL SECTION

Because of the extreme sensitivity of the sulfide compounds to moisture, all operations were performed 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 amounts of Li2S and P2S5 were ground first in an agate mortar and poured into the acetonitrile solvent. The powder:solvent ratio is 1:20 in grams per milliliter. The mixture was heated at 50 °C on a hot plate and stirred for 3 days. A slightly greenish solution containing white precipitate was obtained. The solvent was then allowed to evaporate on a hot plate 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 min. The powder precipitate and the supernatant were separated after decanting. Both were then dried under vacuum. A clear gel-like 991

DOI: 10.1021/acs.chemmater.7b04842 Chem. Mater. 2018, 30, 990−997

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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 PP, DP, and SP samples. impedance measurement. The measured temperature range was −40 to 100 °C. 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); high-temperature testing was performed inside a heating oven using a Biologic potentiostat (VMP3). The frequency range was 1 MHz to 1 Hz.

appears to be a small exothermic peak in the DTA curve, and a slight change in 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 below. 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 samples PP and SP. The morphologies of PP-260, DP-260, and SP-260 were characterized by scanning electron microscopy (SEM) (Figure 2). PP-260 and DP-260 samples were in the powder form; SP-

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RESULTS AND DISCUSSION White precipitates appear immediately after the mixture of Li2S and P2S5 (originally yellowish) is added to acetonitrile. The colorless solution turns bluish within 1 min. The bluish color is preserved 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 yields of DP and PP are ∼0.7 of ∼1.15 g, respectively; a clear gel-like substance remains after the evaporation of acetonitrile. These pieces of evidence clearly suggest that some of the 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 samples PP and DP exhibit one major thermal event 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 and 220 °C vs 190 and 210 °C, respectively). Both samples experience large weight losses during the event. Rangasamy et al. reported that Li2S and P2S5 at a 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 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 ∼100 °C.36 However, on the basis of 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

Figure 2. Scanning electron microscopy images of samples (a) PP-260, (b) DP-260, and (c) SP-260.

260 was obtained by casting the supernatant onto an aluminum substrate followed by annealing. Sample PP-260 consists of submicrometer primary particles (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 than 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 that of PP-260, but there seems to be more agglomeration or secondary amorphous phase binding of 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 seem to be 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 match well with those of the crystalline phase 992

DOI: 10.1021/acs.chemmater.7b04842 Chem. Mater. 2018, 30, 990−997

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Chemistry of Materials “Li3PS4·2ACN”;40 after annealing at 200 or 260 °C, the patterns (Figure 3) match well with those of the reported β-

desired Li7P3S11 phase. The conversion occurs at a temperature that is slightly higher (∼10 °C) than that of the conversion of sample PP into β-Li3PS4, as seen from TG and DTA, which explains the presence of both β-Li3PS4 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 with the release of ACN and (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 for solid-state diffusion and reaction, the formed crystalline Li7P3S11 may suffer from local inhomogeneity and nonstoichiometry. 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. 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 improve our understanding of the local structures of the samples. The Raman spectra of PP-200 and PP-260 are similar (Figure 4). A single peak at 426 cm−1 can be

Figure 4. Raman (left) and 31P MAS NMR (right) spectra of samples DP, PP, and SP annealed at different temperatures. The asterisk denotes a spinning sideband.

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

assigned to the local structural unit of PS43− tetrahedra, confirming that only the 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 cm−1 corresponding to PS43−, consistent with the crystalline structure of the Li7P3S11 phase.27 The 426 cm−1 peak is stronger in sample DP-200. This is expected as the residual β-Li3PS4 phase is present in the XRD pattern of DP200. The Raman spectrum of SP-260 is not presented because of the large fluorescence background, which hinders the determination of its local structure. The 31P MAS NMR spectra of samples 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 and 87 ppm can be assigned to the structural units 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− units in the

42

Li3PS4. Peaks corresponding to Li2S were observed in both patterns, indicating small amounts of Li2S were present in both samples. The crystallinity of sample PP is higher with a 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 transforms into β-Li3PS4 phase at >190 °C. For sample DP, the XRD pattern before annealing is similar to that of PP (Figure S3), suggesting that the SP coverage on PP is indeed amorphous. After annealing, the pattern of DP260 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 reflect only 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 993

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Figure 5. Schematic illustration of the mechanism of formation of Li7P3S11 synthesized in acetonitrile.

converted to the high-conductivity Li7P3S11 phase after annealing. The proposed mechanism of formation 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 mechanism of formation of the Li3PS4·ACN precipitate is critically important, which has not been elucidated. Given the insoluble nature of Li2S and high solubility of “Li2S·P2S5”, the formation of Li 3 PS 4 ·ACN most likely occurs through conversion of Li2S particles from outside to inside by the “Li2S·P2S5” solution. Under such an assumption, the morphology of Li2S particles will determine the morphology of the Li3PS4·ACN precipitate. When the Li2S precursor particles are large, the conversion may never reach completion because of the passivation of the outer Li3PS4·ACN, which can explain the presence of Li2S in samples PP-200 and PP-260. With annealing of the precipitates and solution phase mixture, a majority of the solution phase will be “assimilated” during the solid-state reaction but a small amount may remain on the surface of Li7P3S11 particles. Therefore, the state into which the solution phase transforms at the annealing temperature will influence the transport property of DP-260. The PXRD patterns of SP-200 and SP-260 have large amorphous backgrounds, consistent with the SEM observation that the sample remains mostly amorphous. Bragg peaks in SP200 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 the DTA data of sample SP can then be attributed to crystallization of the amorphous phase. The 31P NMR spectrum of SP-260 is quite complex. 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 shifts are unidentified. Nevertheless, none of the crystalline Li 4 P 2S 6 , Li 2 P 2 S 6 , or glassy Li 2 S−P 2 S 5 has good ionic conductivity,50−52 so the presence of the residual solution phase is expected to lower the overall ionic conductivity of sample DP-260. The ionic conductivity of sample DP-260 was characterized by electrochemical impedance spectroscopy over the temperature range of −40 to 100 °C. The impedance plots are shown in Figure S4, and detailed fitting information can be found in the Supporting Information. At or above room temperature (22 °C), the impedance plots show a straight line intercepting or trending toward the x-axis at low and intermediate frequencies whereas the bending at high frequencies is due to stray

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 samples PP-260 and DP-260, suggesting that both samples are mostly crystalline. Two small peaks at 109 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 the Li4P2S6 phase.47 We can safely conclude from XRD, Raman, and 31P NMR analysis that the crystalline Li7P3S11 phase is formed by conversion of Li3PS4 or Li3PS4·ACN and the solution phase through solid-state reaction. There remains a question about what the solution phase is. To answer the question, we conducted a similar synthesis but varied the Li2S:P2S5 molar ratio. Interestingly, the quantity of powder precipitates decreases with an increase in the amount of P2S5; e.g., at a 55:45 ratio, only 0.2 g of PP was obtained from a total of 1 g of precursors. The XRD pattern of PP at a 55:45 ratio is similar to that at a 75:25 ratio (Figure S3), also matching the dried precursor:Li7P3S11 ratio of Yao et al.18 and that of the “Li3PS4· 2ACN” phase of Rangasamy et al.,40 indicating that all the powder precipitates are essentially the same compound. At a 1:1 ratio, Li2S and P2S5 dissolve quickly to form a clear yellowish solution. Indeed, Liang et al. have 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 a Li2S:P2S5 molar ratio of 1:1 that either dissolves in or precipitates out of the solution, because the solvated form of the solution phase or the structure of the amorphous solids after solvent removal was 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 P2S62− unit that is soluble in acetonitrile. We believe that Li2S and P2S5 (in an x:100 − x molar ratio, where 50 < x < 75) participate in the reaction in acetonitrile as follows: x Li 2S + (100 − x)P2S5 + ACN → 2(x − 50)Li3PS4 ·ACN(p) + 2(75 − x)Li 2S· P2S5(s)

When x < 50, soluble “Li2S·P2S5” forms and some P2S5 remains undissolved; when x > 75, the Li3PS4·ACN precipitate forms and some Li2S remains undissolved. In the case of 70Li2S· 30P2S5 under equilibrium conditions, 80 mol % Li3PS4·ACN (PP) precipitates out while 20 mol % “Li2S·P2S5” dissolves in the solution. After acetonitrile evaporation, the Li3PS4·ACN precipitates are coated by the amorphous “Li2S·P2S5” phase and then 994

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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 for further improving the quality of Li7P3S11 from the liquid phase synthesis. Measures that could potentially improve the 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 the annealing process for better solid-state conversion, and (3) reducing the size of the Li3PS4·ACN particle to shorten diffusion length, etc. It is evident that the intergrain process has an activation energy much higher than that of the intragrain process. Although indeterminable from the impedance spectrum, the intergrain resistance at 22 and 50 °C can be calculated by extrapolation to be 10 and 1 Ω, respectively; the intergrain resistance is negligible at high temperatures.

capacitance and stray inductance. The straight line represents a 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), 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 × 10−4 S cm−1. The activation energy was calculated on the basis of the Arrhenius equation: σ = σ0/T exp(−Ea/kT). It can be seen (Figure 6) that the total ionic conductivity shows a good

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CONCLUSIONS The mechanism of formation of crystalline Li7P3S11 solid electrolytes synthesized in acetonitrile was revealed by analyzing the intermediate products during the reaction. It is found that the precursors (70:30 Li2S:P2S5 molar ratio) in acetonitrile form Li3PS4·ACN (75:25 Li2S:P2S5) precipitates and soluble “Li2S·P2S5” (50:50 Li2S:P2S5). Unlike the β-Li3PS4 crystalline phase that 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 for further optimizing the transport properties and morphologies of liquid phasesynthesized Li7P3S11 solid electrolytes.

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Figure 6. Arrhenius plot for the total and intergrain ion conductivity of Li7P3S11 (260 °C, DP-260).

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.chemmater.7b04842.

Arrhenius behavior in the entire temperature range with an activation energy of 0.37 eV (36 kJ mol−1). This value is close to that of 70Li2S·30P2S5 glass obtained by mechanical milling and much lower than that of 70Li2S·30P2S5 glass ceramics (0.18 eV). Seino et al. showed 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 because the XRD and NMR data indicated that sample DP-260 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 a 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 to 38 kJ mol−1.18,37,39 This is not surprising considering the complexity of the formation mechanism in the liquid phase as

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Additional SEM images and XRD data, impedance spectra, and detailed fitting information (PDF)

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Dongping Lu: 0000-0001-9597-8500 Patrick Z. El Khoury: 0000-0002-6032-9006 Kee Sung Han: 0000-0002-3535-1818 Ji-Guang Zhang: 0000-0001-7343-4609 Jun Liu: 0000-0001-8663-7771 Notes

The authors declare no competing financial interest. 995

DOI: 10.1021/acs.chemmater.7b04842 Chem. Mater. 2018, 30, 990−997

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Chemistry of Materials

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ACKNOWLEDGMENTS This work was supported by the Energy Efficiency and Renewable Energy (EERE) Office of Vehicle Technologies of the U.S. Department of Energy (DOE) under Contract DEAC02-05CH11231 and DEAC02-98CH10886 for the Advanced Battery Materials Research (BMR) Program and the U.S. DOE EERE Water Power Technologies Office and the U.S. Army Corps of Engineers Portland District. SEM, solidstate 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 DE-AC05-76RLO1830.

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