Amorphous vs. Crystalline Li3PS4: Local Structural Changes During

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C: Energy Conversion and Storage; Energy and Charge Transport 3

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Amorphous vs. Crystalline LiPS: Local Structural Changes During Synthesis and Li Ion Mobility Heike Stöffler, Tatiana Zinkevich, Murat Yavuz, Anna-Lena Hansen, Michael Knapp, Jozef Bednar#ík, Simon Randau, Felix H. Richter, Jürgen Janek, Helmut Ehrenberg, and Sylvio Indris J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.9b01425 • Publication Date (Web): 01 Apr 2019 Downloaded from http://pubs.acs.org on April 1, 2019

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The Journal of Physical Chemistry

Amorphous vs. Crystalline Li3PS4: Local Structural Changes During Synthesis and Li Ion Mobility Heike Stöffler*,a, Tatiana Zinkevicha,b, Murat Yavuza, Anna-Lena Hansena, Michael Knappa, Jozef Bednarčíkc,d, Simon Randaue, Felix H. Richtere, Jürgen Janeke,f, Helmut Ehrenberga,b, Sylvio Indrisa,b

aInstitute

for Applied Materials – Energy Storage Systems (IAM-ESS), Karlsruhe Institute of

Technology (KIT), Hermann-von-Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germany bHelmholtz

Institute Ulm, Helmholtzstraße 11, 89081 Ulm, Germany

cDeutsches

Elektronen Synchrotron DESY, Notkestrasse 85, 22607 Hamburg, Germany

dDepartment

of Condensed Matter Physics, Institute of Physics, P. J. Safarik University, Park

Angelinum 9, 041 54 Kosice, Slovakia eInstitute

of Physical Chemistry, Justus-Liebig-University Giessen, Heinrich-Buff-Ring 17, 35392

Giessen, Germany fBELLA-Batteries

and Electrochemistry Laboratory, Institute of Nanotechnology, Karlsruhe Institute of

Technology (KIT), Hermann-von-Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germany

Corresponding Author: *Heike Stöffler Institute for Applied Materials – Energy Storage Systems (IAM-ESS) Karlsruhe Institute of Technology (KIT) Hermann-von-Helmholtz-Platz 1 76344 Eggenstein-Leopoldshafen Germany Tel.: +49-721-680-28502 Email: [email protected] 1 ACS Paragon Plus Environment

The Journal of Physical Chemistry 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

Abstract Glass–ceramic solid electrolytes have been reported to exhibit high ionic conductivities. Their synthesis can be performed by crystallization of mechanically milled Li2S-P2S5 glasses. Herein, the amorphization process of Li2S–P2S5 (75:25) induced by ball milling was analyzed via X-ray diffraction (XRD), Raman spectroscopy, and 31P magic-angle spinning (MAS) nuclear magnetic resonance (NMR) spectroscopy. Several structural building blocks such as [P4S10], [P2S6]4-, [P2S7]4-, and [PS4]3- occur during this amorphization process. In addition, high-temperature XRD was used to study the crystallization process of the mechanically milled Li2S–P2S5 glass. Crystallization of phase-pure β-Li3PS4 was observed at temperatures up to 548 K. The kinetics of crystallization was analyzed by integration of the intensity of the Bragg reflections. 7Li

NMR relaxometry and pulsed field-gradient (PFG) NMR was used to investigate the short-range and

long-range Li+ dynamics in these amorphous and crystalline materials. From the diffusion coefficients obtained by PFG NMR, similar Li+ conductivities for the glassy and heat-treated sample were calculated. For the glassy sample and the glass-ceramic β-Li3PS4 (calcination at 523 K for 1 h), a Li+ bulk conductivity of σLi = 1.6 × 10-4 S/cm (298 K) was obtained, showing that for this system a wellcrystalline material is not essential to achieve fast Li-ion dynamics. Impedance measurements reveal a higher overall conductivity for the amorphous sample, suggesting that the influence of grain boundaries is small in this case.

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Introduction All-solid-state batteries (ASSB) using inorganic solid electrolytes instead of flammable organic liquid electrolytes are highly desired for a safe operation of lithium-ion batteries. In particular, lithium-ion conducting thiophosphates have attracted interest as solid electrolytes due to their high Li+ conductivity and compliant mechanical properties. Some crystalline compounds such as Li7P3S11 and Li10GeP2S12 have been classified as superionic conductors.1,2 The formation of such crystalline compounds in the Li2S-P2S5 system can be performed via heat treatment of the corresponding glasses. A large difference in the Li ion conductivities has been observed depending on the synthesized phase. For 70 mol % Li2S, a large increase from 5 × 10-5 Scm−1 to 3.2 × 10−3 Scm−1 was observed by heating the glass to 633 K. This increase is due to the formation of the superionic crystalline phase Li7P3S11.3 However, for other lithium thiophosphates such as Li4P2S6 and Li2P2S6 a very low ionic conductivity has been reported.4,5 Besides high ionic conductivities the chemical stability of the solid electrolyte plays a crucial role in solid state batteries. In the Li2S-P2S5 system, Li3PS4 is likely to be one of the most suitable compounds for use with lithium metal.6,7 β-Li3PS4 is a highly ion-conductive phase, while γ-Li3PS4 is a poor ionic conductor.8 Liu et al. have reported the stabilization of the highly conductive β-Li3PS4 phase by the formation of a nanoporous structure using a solvent-based synthesis. They also suggest that the transition from amorphous to crystalline β-Li3PS4 is important to achieve high ionic conductivity.6 An increase in the conductivity achieved by heating a mechanically milled glass (75 mol% Li2S) above the crystallization temperature was also reported by Mizuno et al.9 On the other hand, the synthesis of amorphous Li3PS4 via mechanical milling has also shown to result in high conductivities of >2 × 10-4 S/cm.10,11 In addition, Tsukasaki et al. reported that the crystallization of a 75Li2S-25P2S5 glass at temperatures above 453 K leads to a decrease of the ionic conductivity.10 It has been shown by neutron diffraction12 and by nuclear magnetic resonance (NMR) relaxometry13, that the diffusion in crystalline -Li3PS4 is two-dimensional. In this paper, we report on investigations via X-ray diffraction (XRD), Raman spectroscopy, and 31P magic-angle spinning (MAS) NMR spectroscopy of the amorphization process of Li2S–P2S5 (75:25) induced by ball milling. High-temperature XRD was used to observe the structural evolution and phase 3 ACS Paragon Plus Environment


transitions of the glassy sample during the heating process. Furthermore, the kinetics of crystallization was analyzed by integration of the intensity of the Bragg reflections. The short-range and long-range Li+ dynamics in these amorphous and crystalline materials are studied by 7Li NMR relaxometry and pulsed field-gradient (PFG) NMR. These results are compared with those from electrochemical impedance spectroscopy (EIS).

Experimental Li2S-P2S5 glassy solid electrolyte was prepared by mechanical milling of the binary sulfides Li2S and P2S5. Subsequently, the glassy sample was calcined in order to obtain the glass-ceramic β-Li3PS4. For the ball-milling process, first a mixture of 75 mol % Li2S (Alpha Aesar 99.9 %) and 25 mol % P2S5 (Honeywell-Fluka 99 %) was placed in a ZrO2 bowl of 45 ml together with 100 g ZrO2 balls with a diameter of 3 mm. A batch of 4 g of the powder mixture was ball-milled using a planetary ball mill Pulverisette 7 Premium Line (Fritsch) under argon atmosphere at 510 rpm. Up to 350 milling cycles (5 min of milling alternating with 15 min of resting for cooling) were applied to complete the amorphization process. The calcination was performed in Al2O3 crucibles sealed in quartz glass tubes under vacuum. The sample labeled as “calcined” was calcined at 523 K in the furnace for 1 h and slowly cooled down to room temperature at a natural rate. The heating rate was 4 K/min. Raman spectroscopy was performed with a LabRAM HR Evolution spectrometer (HORIBA Scientific), using a 100x magnification objective and an excitation wavelength of 632.81 nm. The spectral data were acquired with exposure times of 10 s and with 15 accumulations over the wavenumber range of 100 to 1198 cm-1. Samples were measured in sealed glass capillaries. All Raman spectra were baseline corrected and normalized to unity using the Horiba Labspec 6 software. Room temperature and some of the high-temperature X-ray diffraction patterns were collected using a STOE Stadi P powder diffractometer with Mo-Kα1 radiation (λ= 0.70932 Å). Powder samples were measured in sealed glass capillaries with a diameter of 0.5 mm. For the high-temperature measurements, sealed quartz glass capillaries were used, which were placed in a graphite heater. A water-cooling system


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was employed together with N2 as inert gas to prevent oxidation of the graphite rod. The heating/cooling rate was 323 K/min. Room temperature X-xay total scattering experiments were performed using a STOE Stadi P powder diffractometer with Ag-Kα1 radiation (λ= 0.55940 Å), equipped with a Dectris MYTHEN 1k detector. Powder samples were measured in sealed glass capillaries with a diameter of 0.5 mm. An empty capillary was measured under the same conditions and used for background subtraction. To account for instrumental resolution LaB6 (NIST660a) was measured. The corresponding Pair distribution function (PDF) was calculated using the xPDFsuite with a Qmax =17.3 Å-1. Qdamp = 0.0997Å-1 was determined by fitting LaB6 PDF using PDFgui.14,15 Nanoparticle models and their PDFs were calculated using DISCUS.16 Room temperature and high-temperature synchrotron diffraction studies were performed at the highresolution powder diffraction beamline (P02.1) at PETRA-III, DESY, using synchrotron radiation with a photon energy of 60 keV (λ = 0.2072 Å).17 The sample was sealed in a 0.5 mm glass (or quartz glass for the HT-XRD) capillary and heated in a ceramic oven from room temperature up to 973 K. The diffraction patterns were acquired using a Perkin Elmer area detector with a sample-detector distance of 1610 mm. A heating rate of 20 K/min was used and the exposure time for each diffraction pattern was 1 min. The obtained 2D images were integrated to 1D patterns by using the program Fit2D.18 The patterns were analyzed by the Rietveld method using the Full-Prof software.19 Differential Scanning Calorimetry (DSC) coupled with Thermogravimetric Analysis (TGA) was done on a STA 449 C Jupiter instrument (NETZSCH, Germany). The experiments were performed at temperatures up to 973 K in argon with constant heating rates of 5 K/min and a flow rate of 312.5 ml/min. The data treatment was performed with the NETZSCH Proteus Thermal Analysis software. 31P

MAS NMR spectra were obtained at a spinning speed of 20 kHz on a Bruker Avance 500 MHz

spectrometer at a field of 11.7 T, corresponding to a resonance frequency of 202.5 MHz. For these measurements, the sample was packed into a 2.5 mm zirconia MAS rotor in an argon filled glovebox. The spectra were acquired with a rotor-synchronized Hahn-echo pulse sequence. The 31P chemical shift was calibrated using H3PO4 (85 %, 0 ppm). 7Li NMR spin-lattice relaxation time measurements were 5 ACS Paragon Plus Environment


performed using a Bruker Avance 200 MHz spectrometer at a field of 4.7 T (77.8 MHz), on samples sealed in 10 mm glass vials. A saturation-recovery pulse sequence was used and the temperature range was 210 to 560 K. 7Li PFG-NMR measurements were performed on a Bruker Avance 300 MHz spectrometer operated at a 7Li frequency of 116.6 MHz. The spectrometer was equipped with a PFG NMR probe, which provides pulsed field gradients up to 30 T/m. A stimulated-echo pulse sequence with bipolar gradients20 was used to suppress the influence of eddy currents. The gradient duration and the diffusion time were 2.51 ms and 300 ms, respectively. Scanning electron microscopy (SEM) images were recorded on a Zeiss Merlin microscope using 5 kV acceleration voltage. Nitrogen physisorption experiments were carried out on a Quantachrome Quadrasorb evo instrument. The used measurement parameters were pressure tolerance of 0.05 Torr, equilibration time of 60 seconds, and equilibrium timeout of 120 seconds. The Brunauer-Emmett-Teller (BET) specific surface area was determined from five data points between a relative pressure (p/p0) of 0.1 and 0.25. Electrochemical impedance spectroscopy (EIS) on β-Li3PS4 (sample with pellet geometry, 60 mg, 425 µm thickness, 10 mm diameter) was performed using an EC-Lab Electrochemistry VMP-300 Biologic in the frequency range of 7 MHz to 1 Hz, applying a 10 mV signal amplitude. Impedance plots were fitted using RelaxIS software.

Results and discussion Glass formation and crystallization. Figure 1 shows the XRD patterns of the starting materials Li2S and P2S5 and the evolution of the powder mixture after different ball milling times. The pattern of the powder mixture without milling (0 h) is simply a superposition of the patterns of the initial binary compounds. Subsequently, a gradual broadening and intensity decrease of the Bragg reflections can be observed with increasing milling time. After a milling period of more than 25 h, no reflections of the starting materials are visible, indicating a complete glass formation. This is in accordance with literature, where in the Li2S-P2S5 system XRD patterns characteristic of amorphous materials were obtained for Li2S contents of up to 75 mol %.21,11,22 In addition, Dietrich et al. verified the amorphous character of the samples by employing differential scanning calorimetry to detect the glass transition temperature. 6 ACS Paragon Plus Environment

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The effect of the glass modifier Li2S was described by Dietrich et al.: The PS4 polyhedra in the precursor P2S5 (P4S10) are connected to three other polyhedra by P-S-P bridges. By incorporating Li2S, the sulfur content is increased and the P-S-P bridges are broken, resulting in terminal sulfur atoms with Li+ acting as the counter ion. However, unreacted crystalline Li2S should be avoided, since it has low ionic conductivity.

Figure 1: XRD patterns (Molybdenum Kα1 radiation) of the binary starting materials Li2S and P2S5, and of 75Li2S-25P2S5 powders after different milling times.

Figure 2a shows the characteristic Raman spectra of the starting materials Li2S and P2S5 and the spectra of the 75Li2S-25P2S5 powder mixture after different ball milling times. The Li2S Raman spectrum shows



only one vibration mode around 377 cm-1, which can be assigned to the Raman mode with F2g symmetry in Li2S (372 cm-1).23,24 The Raman spectrum of P2S5 (P4S10) is also in good agreement with literature25,26,27,28, although some additional peaks (indicated by *) evidence the presence of decomposition products. The additonal peaks can be seen in the Raman spectrum of P2S5 at 388 and 307 cm-1. They can be assigned to P4S9, and a small peak at 234 cm-1 is ascribed to P4S7. These phases might derive from a minor decomposition of P4S10 and a secondary decomposition of P4S9 via P4S8.27 The initial compound P2S5 was analyzed by XRD. The Rietveld Refinement reveals phase pure P4S10, (see Supporting Information, Figure S1). This indicates that the observed decomposition products observed in the Raman spectrum are minor and were only detected due to an inhomogeneous distribution in the small sample volume. Similar to earlier reports27, also a minor amount of sulfur was observed in the Raman spectrum at 486 cm-1, resulting from a decomposition of P4S10 by separation of exocyclic sulfur. Somer et al. classified the P4S10 Raman spectra into 4 categories: Exocyclic P-S valence vibrations (680 - 720 cm-1), P-S-P bridge vibrations (395 - 560 cm-1), endocyclic deformations (180 - 350 cm-1) and exocyclic P-S deformations (110 - 165 cm-1). A Raman spectrum of an empty capillary was also measured, showing no detectable signals and thus indicating that there is no interference with the spectra of the samples (see Supporting Information, Figure S2). The Raman spectrum of the initial powder mixture (0 h) is again the superposition of the spectra of the starting materials Li2S and P2S5. Upon amorphization the characteristic bands of the starting materials disappear and several peaks corresponding to the building blocks present in Li7P3S11, Li4P2S6 and Li3PS4 are observed. The spectra of the mixture milled for 8 h reveal the presence of isolated [PS4]3- units (418 cm-1) and [P2S7]4- (404 cm-1). The Raman band at 386 cm-1 is assigned to the [P2S6]4- units of Li4P2S6.5,29 For the amorphous sample (30 h) and the calcined sample a Raman band around 420 cm-1 is observed. This band can be attributed to the symmetric stretching vibration of the P-S bonds in the isolated [PS4]3units of Li3PS4.29 The small shoulder in the amorphous sample (30 h) at 386 cm-1 indicates the presence of the anionic building block [P2S6] 4-.4,29 Furthermore, the presence of a small amount of [P2S7]4- (404 cm-1) is also possible. 8 ACS Paragon Plus Environment

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Figure 2: (a) Raman spectra and (b) 31P MAS NMR spectra of Li2S-P2S5 after different ball milling times and after calcination (523 K for 1 h). The asterisk indicates the presence of a small amount of decomposition products. The 31P NMR spectra of these samples are shown in Figure 2b. The spectrum of the initial Li2S-P2S5 mixture (0 h) shows a broad peak between 65 ppm and 45 ppm. This can be explained by the structure of the precursor P2S5, consisting of adamantane-like 3-dimensional P4S10 cages, which are not highly symmetric but slightly distorted. Therefore, the four P positions are similar but not fully equivalent, giving rise to the multiple overlapping peaks in the region between 65 and 45 ppm. During milling, these peaks successively disappear and new peaks occur. For the sample milled for 8 h, two broad overlapping peaks can be observed between 70 - 118 ppm. These broad peaks can be assigned to isolated [PS4]3tetrahedra (87 ppm), [P2S6]4- units (105 ppm), and [P2S7]4- units (91 ppm).5 The large width of these peaks results from the strong variations of the bond angles inside the local structural building blocks which are consistent with an overall amorphous structure, in good agreement with the XRD results. The spectrum of the amorphous sample (30 h) shows a broad peak at 83 ppm corresponding to [PS4]3units and a small broad peak (105 ppm) that can be assigned to [P2S6]4-. Again, some small contribution of [P2S7]4- units (91 ppm) might be present. Overall the Raman and 31P NMR spectra of the sample ballmilled for 30 h match well with the reported spectra of 75Li2S-25P2S5 glass.11 9 ACS Paragon Plus Environment


After calcination, a strong and narrow peak at 87 ppm is observed, which is assigned to the uniform [PS4]3- units in the crystalline Li3PS4 phase.30,11,12 In addition, a tiny peak corresponding to pyrothiophosphate [P2S7]4- is observed, which is formed during crystallization by a redox process in which P4+ in the glassy [P2S6]4- is oxidized to P5+ [P2S7]4-. A similar redox process was observed by Seino et al. for the crystallization of Li7P3S11 by 31P NMR spectroscopy.1 Based on in house diffraction experiments using Ag radiation the corresponding pair distribution functions (PDF) for three different samples were calculated, a sample before ball milling (0 h), an Xray amorphous sample (30 h) and a calcined sample (calcined), respectively. To give an impression of the altered crystallinity after ball milling and calcination a general comparison of these PDFs is given in Figure 3. As already demonstrated by XRD analysis the crystallinity decreases markedly and the PDF of the corresponding sample (blue curve in Figure 3) reveals no reliable interatomic distances, i.e. structural information, above ~ 10 Å. Above this value, the PDF is dominated by continuous fluctuations, so called termination ripples, due to the restricted Qmax of 17.3 Å-1. Furthermore, a closer look on the short-range region of the PDF, corresponding to the local structure, indicates even more alterations of the structure during the calcination process (Figure 4a). Similar deviations were detected by Shiotani et al.31, though our findings suggest a different underlying cause.


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Figure 3: Comparison of the pair distribution functions for a sample before ball milling (0 h), an X-ray amorphous sample (30 h), and a calcined sample (calcined). All PDFs were calculated based on in house XRD data collected with Ag radiation. Since NMR and Raman spectroscopy already indicated the presence of a [P2S6]4- containing compound in the ball milled sample, different models of Li3PS4 and Li4P2S6 nanoparticles were calculated. The best agreement was achieved using a 25 Å model for Li3PS4 and a nanoparticle half this size for Li4P2S6. The corresponding PDFs are depicted in Figure 4a, completed by assignments of the major radial distances in the two anionic species Figure 4b. The major deviations of the PDFs (marked by grey dotted lines) could indeed be attributed to the disappearance of Li4P2S6 during calcination. Furthermore, a notable deviation of the relative intensity of the first two peaks were observed, marked by a black arrow in Figure 4a. The second peak at 2.5 Å is attributed to Li–S bonds in tetrahedral coordination. Therefore, an increasing probability of randomly distributed Li vacancies was added to the models, where the maximum Li occupancy was set to 0.65 for Li3PS4 and 1 for Li4P2S6, respectively. Not surprisingly, the intensity of the [Li-S4] peak is decreased when the probability of vacancies is increased (red to black curve in Figure 4). The same holds for the [Li-S6] peak for octahedrally coordinated Li in Li4P2S6 (blue to black curve).



(a)

(b)

Figure 4: (a) Zoom into the short range region of the PDFs depicted in Figure 3 for the ball milled (30 h) and calcined sample (calcined) compared with modeled PDFs of Li3PS4 (red) and Li4P2S6. Details on the models are given in the text. Grey dotted lines mark major deviations caused by the presence of Li4P2S6. (b) Assignments of the major radial distributions present in the two anionic species. Comparing the modeled PDFs with the measured ones indicates that a major amount of Li is randomly distributed in the sample and not located on specific crystallographic positions in the lattice. Nevertheless, keeping in mind that this data was measured on an in house Ag radiation diffractometer and has severe termination ripples due to the limited Q space, these findings remain indications and further neutron-PDF studies have to be employed to conclusively determine the local structure especially for Li.


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Figure 5: (a) High-temperature synchrotron X-ray diffraction patterns of the amorphous sample for temperatures between 298 K and 973 K, (b) phase fractions obtained from the Rietveld refinements between 473 K and 873 K, and (c) TGA/DSC scans of the amorphous sample. Figure 5a shows the patterns of the high-temperature X-ray diffraction measurements at temperatures between 298 K and 973 K. The crystallization from the amorphous state to β-Li3PS4 starts at 473 K and the crystallinity increases with increasing temperature. Between 573 K and 773 K the β-phase is transformed to α-Li3PS4, which is present until 873 K. According to Homma et al., during the cooling process the previously formed α-phase (between 724 K and 811 K) is transformed to the γ-phase (without the appearance of β-phase). The formation of γ-phase was also confirmed in this study (see Supporting Information Figure S3 and Figure S4). Since the γ-phase has been reported to exhibit a very low ionic conductivity (3  10-7 S/cm)29, careful temperature treatment of the glass is very important. Above 923 K the sample melts and no reflections are visible. From the Rietveld refinements between 473 K and 873 K, phase fractions of the β- and α-phase were extracted (Figure 5b). Very small additional reflections were seen in the patterns from 573 K to 723 K, which might be explained by presence of a



small amount of Li7P3S11 (space group P-1). Due to the low symmetry and small amount of this phase, it was not included in the refined structure model. The TGA and DSC results (Figure 5c) are in good agreement with the HT-XRD measurements. The exothermic peak at around 508 K shows the crystallization process to β-Li3PS4. The exothermic broad peak, observed upon further heating (between 548 K and 773 K), corresponds to the phase transition from the β- to the α-phase, while the endothermic peak at 933 K describes the melting process. In the TGA results a total mass loss of around 6 wt% was observed. This mass loss might result from some sulfur sublimation. In the crystalline phase the mass loss is around 3 % (not shown here). The observed X-ray diffraction pattern for β-Li3PS4 (calcined at 523 K for 1 h) with the calculated one after Rietveld refinement is shown in Figure 6. The observed pattern can be well described by the single β-Li3PS4 phase with symmetry Pnma and lattice constants a = 12.8829(1), b = 8.1404(2), c = 6.1459(2). The cell volume is V = 644.5(3) Å3.

Figure 6: X-ray diffraction pattern (Mo-Kα1) of β-Li3PS4 (calcined at 523 K for 1 h), measured at room temperature, with the calculated profiles after Rietveld refinement (298 K).


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Figure 7: (a) XRD patterns (Mo-Kα1) of one ball milled Li2S-P2S5 mixture after different resting times at 453 K and (b) integrated XRD intensities of the Bragg reflections at 11.3° and 13.6° as a function of resting time. The crystallization kinetics was studied by heating the amorphous powder to 453 K and holding this temperature for several hours. Figure 7a shows the recorded XRD patterns for the different resting times. A stepwise crystallization of the powder to β-Li3PS4 can be observed. In order to achieve this gradual evolution, the choice of the temperature is important. Higher temperatures result in a very fast crystallization, making it difficult to record the crystallization step by step. For lower temperatures no or a too slow crystallization is observed (see Supporting Information, Figure S5). For RbSn2F5, a crystallization starting from nm-sized crystallites was observed even at room temperature during storage over five months.32 In our case, such a room-temperature crystallization can be excluded from both XRD and NMR measurements. Figure 7b shows the integrated intensities of the two Bragg reflections at 13.3° and 13.6° as a function ―𝑡

of heating time. The overall behavior can be well described by an exponential function (𝐼 = 𝐼0 +𝐼1 ∙ 𝑒 𝜏 ). From the corresponding fit (dashed line) a time constant for the crystallization of τ = (4.1 ± 0.7) h can be extracted. 15 ACS Paragon Plus Environment


Figure 8 summarizes the structural changes during the synthesis of crystalline Li3PS4. The synthesis of β-Li3PS4 includes an amorphization and a crystallization step. The amorphous sample can be characterized by the presence of mainly [PS4]3- and [P2S6]4- units, while the crystalline sample is dominated by uniform [PS4]3- units.

Figure 8: Summary of the structural changes by amorphization and crystallization.

Ionic conductivity. The local transport of the Li+ ions was studied by 7Li NMR relaxometry. Figure 9 shows the T1-1 relaxation rates as a function of inverse temperature for the calcined sample. A clear maximum can be observed at about 408 K with a low- and high-temperature flank. The activation barrier extracted from the low temperature side is lower (0.16 eV) than that extracted from the high temperature side (0.27 eV). Such an asymmetric behavior around the maximum is characteristic of disordered systems33 and reveals that the motion of the lithium ions is correlated34,35, leading to a deviation from the symmetric behavior described by Bloembergen et al.36 Assuming a completely uncorrelated motion of the Li+ ions and using the activation energy from the high-temperature side of the T1-1 maximum, a hopping rate of 2.9 × 107 s-1 at 298 K was obtained. For a jump length l of 2 Å, which is the shortest Li-Li distance in the structure, a diffusion coefficient of 1.9 × 10-13 m2/s was estimated via the Einstein-Smoluchowski equation37,38. Using the Nernst-Einstein equation together with the Li+ ion concentration (1.9 × 1028 m-3) a Li+ ion conductivity of 0.22 mS/cm can be estimated. For the amorphous sample, such a maximum could not be observed due to the crystallization of the sample that starts at about 433 K. 16 ACS Paragon Plus Environment

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Figure 9: 7Li NMR relaxation rates T1-1 versus inverse temperature for the calcined and amorphous sample.

The long-range transport of the Li+ ions can be directly measured by pulsed field-gradient NMR. Figure 10 shows the echo damping vs. gradient field strength at five different temperatures between 303 K and 343 K for the amorphous and the calcined sample. The damping can be well described by a Gaussian function according to the Stejskal-Tanner equation.39 At higher temperatures, a stronger damping is observed, as expected, due to faster diffusion. The extracted diffusion coefficients versus inverse temperature are shown for the amorphous and calcined sample in Figure 11. The diffusion coefficient at 298 K was extrapolated and is about D298 K = 1.4 × 10-13 m2/s for both samples. Within the observation time of 300 ms, this corresponds to a diffusion length of about 0.5 μm. Using again the Nernst-Einstein equation, the Li+ ion conductivity at 298 K was determined to be σLi = 1.6 × 10-4 S/cm for the amorphous and the calcined sample.



Figure 10: Echo damping vs. gradient field strength g measured with 7Li PFG NMR on the (a) amorphous (b) and calcined sample.

Figure 11: Diffusion coefficients of the amorphous and calcined (523 K for 1 h) sample for temperatures between 303 and 343 K, as determined by 7Li PFG NMR. The activation energies, determined from the slope between 303 K and 343 K for the amorphous and the calcined sample were EA = 0.33 ± 0.01 eV and EA = 0.31 ± 0.03 eV, respectively. The value for the conductivity and the activation energy are in good agreement with the values estimated from the T1-1 measurements described above.


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Figure 12: Arrhenius plots from impedance measurements for the (a) amorphous and (b) calcined sample. From the impedance measurements on pelletized samples, an overall conductivity of σLi=2.98  10-4 S/cm for the amorphous sample and σLi=1.09  10-4 S/cm for the calcined sample (see Supporting Information Figure S6) was found. The value for the amorphous sample is close to the value determined by PFG NMR. The difference in the Li-ion conductivities for the crystalline sample might be explained by the stronger influence of grain boundaries on the dc conductivity probed during the impedance measurements, while PFG NMR is sensitive on time scales of about 1 s and thus to somehow shorter length scales, mainly inside the interior of the crystallites. For the amorphous sample, the influence of grain boundaries seems to be weaker. The activation energies obtained from the corresponding Arrhenius plots were EA=0.37 eV for both the amorphous and the calcined sample (Figure 12), which is in good agreement with the PFG measurements. A similar behavior was already observed earlier by impedance spectroscopy.13 It is interesting to note that for Li7P3S1140 and Li7SiPS841 an opposite trend has been observed, i.e. a higher conductivity for the crystalline material. The dc conductivity dc represents the long-range transport of the Li ions while the PFG NMR measurements are probing somehow shorter length scales of about 0.5 m and T1-1 is sensitive to single Li ion jumps in the ns range. Therefore, the activation energy obtained from dc is slightly higher than those obtained from PFG NMR and T1-1. A similar trend has been observed for a Li3PS4 sample prepared via a solvent-based technique.12 19 ACS Paragon Plus Environment


In Figure 13, a summary of the jump rates, extracted from the T1-1, PFG NMR, and impedance spectroscopy data is shown for the calcined sample.

Figure 13: Temperature dependence of the Li jump rates as obtained from NMR methods as well as dc conductivity in the calcined sample. A good overall agreement is observed. The absolute values have to be taken with some care in view of the approximations included in these calculations. The T1-1 data could be measured in the largest temperature range and they reveal, as expected, a slightly smaller activation energy for the local hopping. The PFG NMR data and the impedance spectroscopy results were measured in a smaller temperature range and they are probing the long-rang-transport of the Li ions. This is in agreement with the slightly higher activation energy. BET Surface Area. A small specific surface area of 2 m2/g was determined by the BET method from nitrogen physisorption (see Supporting Information Figure S7) for both samples. From this it can be concluded that a nanoporous structure with high specific area is not necessary to obtain a high Li-ion conductivity. Morphology. SEM images of the amorphous and calcined samples are shown in Figure 14. Particles with a size between 1 and 15 μm are observed for both samples. In the calcined sample, in addition to the larger particles, a high proportion of smaller primary particles in the nano-size range interconnecting the larger particles is visible. These become apparent in a rougher surface observed for the calcined samples and probably result in a higher number of grain boundaries the Li ions have to overcome for 20 ACS Paragon Plus Environment

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the long-range transport. These additional nanoparticles might be the reason for the increase of the grain boundary resistance as seen via the EIS measurements described above. For the amorphous sample no agglomerated nanoparticles are present, therefore the grain boundary resistance might be stronger suppressed by the applied external pressure during the conductivity measurement. The average crystallite size of the nanoparticles is approximately 20 nm. A similar crystallite size was reported by Tsukasaki et al. using dark field TEM imaging at elevated temperatures between 453 K and 523 K.10 Furthermore, they performed impedance measurements and report on a strong decrease in ionic conductivity at temperatures above 453 K due to a higher degree of crystallinity (defined as the proportion of crystalline regions in an amorphous matrix). In contrast, from the PFG NMR measurements in this work, similar Li+ diffusivity for the amorphous and calcined sample was observed, at least on a length scale of about 0.5 m.

Figure 14: SEM images of the (a) amorphous and (b) calcined sample (523 K for 1 h).



Conclusion We were able to elucidate the formation mechanism during the amorphization process and to synthesize phase pure β-Li3PS4 by a subsequent calcination step. During ball milling of the mixture of the crystalline binary sulfides Li2S and P2S5, the amorphization process involves several anionic building blocks such as [PS4]3-, [P2S6]4- and [P2S7]4-, as confirmed by 31P MAS NMR and Raman spectroscopy. In situ HT-XRD on the amorphous 0.75Li2S-0.25P2S5 sample reveals phase pure crystallization to βLi3PS4 at temperatures up to 548 K. At higher calcination temperatures, α-Li3PS4 is formed which transforms into the γ-phase during cooling and leads to a decrease in the Li-ion conductivity. Therefore, a careful control of temperature is required. We were also able to study the crystallization kinetics and determined a time constant for the crystallization of τ = (4.1 ± 0.7) h at 453 K. For the amorphous and the calcined sample similar Li+ bulk conductivities were determined via PFG NMR measurements (σLi = 1.6 × 10-4 S/cm at 298 K), while impedance spectroscopy show a slightly lower overall Li-ion conductivity for the calcined sample in comparison to the amorphous sample. This suggests that in the 0.75Li2S-0.25P2S5 system a calcination step is not essential to obtain high Li-ion conductivity, since the formation of agglomerated nanocrystallites can have a negative impact on the conductivity. By physisorption measurements a small BET surface area of 2 m2/g was determined for both samples. Our results show that a certain nanostructure is not essential to obtain a good ionic conductivity in Li3PS4, not even a crystalline material is necessary. This might facilitate future upscaling of the synthesis procedure because a sintering step is not necessary. Supporting Information X-ray diffraction pattern (Mo Kα1) of P4S10 with Rietveld refinement, Raman spectra of initial compounds and empty capillary, in situ HT-XRD patterns (Molybdenum Kα1) of the amorphous sample. The patterns were recorded at temperatures from 298 K to 773 K (in red for increasing temperature) and at 298 K (in blue after cooling), in situ HT-XRD patterns (Molybdenum Kα1) of the Li2S-P2S5 glass, heated at 413 K and 433 K, Nyquist plot of electrochemical impedance for the (a) amorphous and (b) calcined sample at selected temperatures, Nitrogen physisorption isotherm of (a) amorphous and (b) calcined sample. 22 ACS Paragon Plus Environment

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Author Information *Email: [email protected] Notes

The authors declare no competing financial interests.

Acknowledgement This work contributes to the research performed at CELEST (Center for Electrochemical Energy Storage Ulm-Karlsruhe) and was financially supported from the Federal Ministry of Education and Research (BMBF) within the FELIZIA project (03XP0026G). This work has benefited from beamtime allocation at the synchrotron facility PETRA III (beamline P02.1), DESY. We thank Mariyam Susana Dewi Darma for performing the PDF measurements. We also thank Liuda Mereacre and Lihua Zhu for performing the TGA and Raman measurements as well as Sebastian Werner and Prof. Bernd Smarsly for their help with nitrogen physisorption.

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Amorphous vs. Crystalline Li3PS4: Local Structural Changes During

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