
==== Front
J Am Chem Soc
J Am Chem Soc
ja
jacsat
Journal of the American Chemical Society
0002-7863
1520-5126
American Chemical Society

39162691
10.1021/jacs.4c06263
Article
Modular MPS3-Based Frameworks for Superionic Conduction of Monovalent and Multivalent Ions
https://orcid.org/0000-0002-2226-9006
Iton Zachery W. B. †
https://orcid.org/0000-0002-6053-4763
Irving-Singh Zion ‡
Hwang Son-Jong ‡
Bhattacharya Amit §
https://orcid.org/0000-0003-1751-4908
Shaker Sammy ∥
https://orcid.org/0000-0002-3320-2157
Das Tridip ⊥
https://orcid.org/0000-0002-3611-1162
Clément Raphaële J. §
https://orcid.org/0000-0003-0097-5716
Goddard William A. III ⊥
https://orcid.org/0000-0002-0133-9693
See Kimberly A. *‡
† Department of Applied Physics and Materials Science, California Institute of Technology, Pasadena, California 91125, United States
‡ Division of Chemistry and Chemical Engineering, California Institute of Technology, Pasadena, California 91125, United States
§ Materials Department and Materials Research Laboratory, University of California, Santa Barbara, Santa Barbara, California 93106, United States
∥ Division of Biology and Biological Engineering, California Institute of Technology, Pasadena, California 91125, United States
⊥ Materials and Process Simulation Center (MSC), California Institute of Technology, Pasadena, California 91125, United States
* Email: ksee@caltech.edu.
20 08 2024
04 09 2024
146 35 2439824414
08 05 2024
08 08 2024
07 08 2024
© 2024 The Authors. Published by American Chemical Society
2024
The Authors
https://creativecommons.org/licenses/by-nc-nd/4.0/ Permits non-commercial access and re-use, provided that author attribution and integrity are maintained; but does not permit creation of adaptations or other derivative works (https://creativecommons.org/licenses/by-nc-nd/4.0/).

Next-generation batteries based on more sustainable working ions could offer improved performance, safety, and capacity over lithium-ion batteries while also decreasing the cost. Development of next-generation battery technology using “beyond-Li” mobile ions, especially multivalent ions, is limited due to a lack of understanding of solid state conduction of these ions. Here, we introduce ligand-coordinated ions in MPS3-based (M = Mn, Cd) solid host crystals to simultaneously increase the size of the interlayer spacing, through which the ions can migrate, and screen the charge-dense ions. The ligand-assisted conduction mechanism enables ambient temperature superionic conductivity of various next-generation mobile ions in the electronically insulating MPS3-based solid. Without the coordinating ligands, all of the compounds show little to no ionic conductivity. Pulsed-field gradient nuclear magnetic resonance spectroscopy suggests that the ionic conduction occurs through a hopping mechanism, where the cations are moving between H2O molecules, instead of a vehicular mechanism which has been observed in other hydrated layered solids. This modular system not only facilitates tailoring to different potential applications but also enables us to probe the effect of different host structures, mobile ions, and coordinating ligands on the ionic conductivity. This research highlights the influence of cation charge density, diffusion channel size, and effective charge screening on ligand-assisted solid state ionic conductivity. The insights gained can be applied in the design of other ligand-assisted solid state ionic conductors, which will be especially impactful in realizing solid state multivalent ionic conductors. Additionally, the ion-intercalated MPS3-based frameworks could potentially serve as a universal solid state electrolyte for various next-generation battery chemistries.

David and Lucile Packard Foundation 10.13039/100000008 NA Innovation and Technology Commission 10.13039/501100003452 NA Camille and Henry Dreyfus Foundation 10.13039/100001082 NA Arnold and Mabel Beckman Foundation 10.13039/100000997 NA Alfred P. Sloan Foundation 10.13039/100000879 NA National Institute of General Medical Sciences 10.13039/100000057 T32 GM008042 document-id-old-9ja4c06263
document-id-new-14ja4c06263
ccc-price
==== Body
pmcIntroduction

The landscape of modern battery technology has been dominated by lithium-ion batteries (LIBs). However, material availability, scalability, cost, and escalating global energy demands1−3 motivate the development of alternative battery chemistries. “Next-generation” mobile ions, such as Na+, K+, Mg2+, Ca2+, Zn2+, and Al3+, represent a promising frontier in battery technology due to their abundance and potential for high volumetric capacities.4 However, one of the major challenges in developing battery technology based on next-generation mobile ions is the difficulty in achieving the solid state conduction of those ions. Solid state conduction is crucial for ion transport in electrodes, interphases, and solid electrolytes. The larger size of ions like Na+, K+, and Ca2+ restricts their movement through the typically rigid migration pathways available in solid materials. Additionally, the mobility of ions with higher charge densities, such as Mg2+, Zn2+, and Al3+, is impeded by the strong electrostatic interactions between these mobile ions and other ions within the solid. These challenges are outlined in detail in our recent perspective.4 Difficulties with solid state ionic conduction are less pronounced in Li-based systems due to the relatively small size and low charge density of Li+.

Solid state conduction of larger cations has typically been achieved in structures based on open frameworks with large migration bottlenecks, like Prussian blue analogues,5−7 β-Alumina,8−11 or NASICON phases,12−14 whereas solid state conduction of charge-dense ions has been mostly limited to electronically conductive materials. In systems like Mo6S8, thiospinel Ti2S4, and MgSc2S4, the mobile electrons are hypothesized to facilitate ionic mobility by screening the charge-dense ions, decreasing the strength of electrostatic interactions within the solid.15−19 For electronically insulating materials, such as solid electrolytes or interphases, there is an absence of mobile electrons to screen the charge of the targeted ions. Only a handful of electronically insulating inorganic solids have been shown to conduct charge-dense ions, such as Zn2+ in ZnPS3,20 or Mg2+ in borohydrides,21−26 but each has its own challenges associated with low RT conductivities (σRT) or high activation energy (Ea).

One path to enable the conduction of large or charge-dense cations in electronic insulators is to introduce ligands, such as H2O, that can coordinate with the targeted cations within the host crystal. Solid state mobility of ligand-coordinated ions was extensively studied in mica clays such as vermiculite and montmorillonite.27−31 However, in clays the σ is generally <0.1 mS cm–1. In battery systems, the addition of H2O to increase multivalent ionic conduction has been primarily attempted for cathode materials, such as MnO2 and V2O5.32−35 Recently, this concept has been extended to electronically insulating inorganic solids, like Li2Sn2S536 and ZnPS3,37 as well as various MOFs and COFs,38−42 for solid state electrolyte applications, however in most cases the conductivity is still below practically useful values or the excess solvent leads to detrimental reactivity in the system.

Here, we leverage coordinated ligand molecules within solid frameworks based on MPS3 materials (M = Mn, Cd) to achieve two primary goals: (1) increasing the size of the interlayer spacing, through which the ions can migrate, and (2) screening charge-dense ions to decrease electrostatic interactions within the lattice. By combining the advantages of large interlayer spacing and charge screening, we can devise a universal framework for ambient temperature superionic conduction of various ions within an inorganic, electronically insulating solid.

CdPS3 and MnPS3 exhibit a peculiar mechanism for the intercalation of hydrated cations into the van der Waals gap.43 Unlike the redox-based intercalation in common battery electrodes in which incorporation of a cation coincides with a Faradaic reduction of the host,44 when cations are intercalated into MPS3 materials, charge balance is maintained by M2+ loss in an ion exchange mechanism.43,45,46 This mechanism is similar to that observed in clays.27,28,47,48 CdPS3 and MnPS3 form monoclinic (C2/m) layered compounds with a slightly distorted hexagonal network of edge-sharing Cd2+ or Mn2+ octahedra. The Cd2+ or Mn2+ bonds are coordinated by [P2S6]4– polyanions. The layers stack along the c-axis separated by a van der Waals gap of ∼3.5 Å.49 The basal spacing (d(001)) of CdPS3 is slightly larger than that of MnPS3 (6.55 vs 6.49 Å) due to the larger size of Cd2+ over Mn2+ (0.95 vs 0.83 Å).50

Here, we exploit the ability of CdPS3 and MnPS3 to host cations within the van der Waals gap to generate materials that contain various ions of interest through a sequential ion exchange strategy that is illustrated in Figure 1. First, a cation with a small hydrated radius (e.g., K+) is intercalated into the van der Waals gap of an MPS3 material.45 Charge neutrality is maintained through the loss of the labile metal (Cd2+ or Mn2+) from the metal layer (eq 1), resulting in negatively charged sheets of MPS3 sandwiching positively charged hydrated ions. The hydrated K+ ions can be further exchanged to introduce a larger hydrated cation (A), like Li+, Na+, Mg2+, Ca2+, Zn2+, and Al3+ (eq 2). Direct insertion of large hydrated cations is kinetically limited; as such, this sequential exchange method is preferred:1

2

Figure 1 Schematic of the ion and ligand exchange used in the present study. In an aqueous KCl solution, CdPS3 intercalates hydrated K+ ions into the van der Waals gap and maintains charge neutrality by losing Cd2+ ions from the metal layer. After K0.5Cd0.75PS3 is formed one can perform either: (1) second ion exchange to introduce desired monovalent or multivalent mobile ions, or (2) ligand exchange to replace H2O in the system.

Additionally, we demonstrate that the H2O molecules can be exchanged for aprotic ligands, e.g., acetonitrile (MeCN), and tetrahydrofuran (THF). The ligand exchange serves several purposes: (1) to probe the effect of different ligand molecules on ionic mobilities, (2) to demonstrate that the mobile ions conduct in the absence of H+, and (3) to broaden the applicability of these frameworks to H2O-incompatible systems.

The ability of MPS3 materials to intercalate a wide variety of guest ions and molecules has been useful for diverse applications. The areas of research span from nonlinear optics,51,52 photoluminescence,53 hosts for biomolecules54 or polymers,55,56 and superconducting magnets.57 However, the ionic conductivity of these materials has been less rigorously explored. Some studies have investigated the mobility of hydrated Na+, K+, and Cs+ in CdPS3.58−61 These studies concluded that Na+ was slightly mobile (10–5 S cm–1) but that K+ and Cs+ were immobile. Another study investigated the conductivity of trivalent cations in MPS3 in the dried state, finding no significant ionic conduction (10–8–10–10 S cm–1).62 Notably, nanosheet-based membranes using Cd0.85PS3Li0.15H0.15 were found to have exceedingly high 2D H+ conductivity (300 mS cm–1 at RT and 98% relative humidity (RH)).63

Recently, Yu and Ren reported on CdPS3-based nanosheet membranes with various intercalated cations.64 The membranes boasted impressive 2D ionic conductivity but importantly, the conductivity was similar irrespective of the intercalated cation (170–330 mS cm–1 at RT and 98% RH). In fact, the ion-intercalated membranes also displayed similar behavior with regard to layer expansion and activation energy. This is likely due to excess H2O content within the restacked membranes which could lead to significant contribution from H+ conduction through a Grotthus mechanism,63,65 or direct diffusion of solvated ions in a confined liquid electrolyte. Such a mechanism is distinct from conduction of ligand-coordinated ions within a host crystal. The transition from ligand-coordinated solid state conduction to the conduction of solvated ions within a confined liquid in the solid framework at excess ligand/solvent concentrations has been noted in previous studies of hydrated Li2Sn2S5, MOFs, and clays.36,66−68

To minimize contributions from H+ conduction and to reveal inherent structure–property relationships that can be drowned out in the presence of excess H2O, here we investigate bulk, ion-intercalated CdPS3 and MnPS3 materials primarily at ambient RH (∼40 to 55%), and with coordinating ligands that range from H2O to aprotic, neutral solvent molecules. This allows us to understand the nuanced differences between the effects of various frameworks, intercalated ions, and coordinating ligands on the ionics, thereby deepening our understanding of ionic conduction in solids. Additionally, the effective bulk pellet conduction enabled by the inherent 2D conduction channels of randomly oriented particles in the polycrystalline samples in this study can be directly applied to practical battery applications. We also introduce MnPS3 as an environmental friendly alternative with comparable performance, and demonstrate nonaqueous analogues that would be more compatible with desirable electrode materials.

We employ several characterization techniques to study the structural and chemical changes after ion exchange as well as the resulting ionic mobility. These techniques include electrochemical impedance spectroscopy (EIS), X-ray diffraction (XRD), thermogravimetric analysis (TGA), inductively coupled plasma mass spectrometry (ICP-MS), scanning electron microscopy (SEM), energy dispersive X-ray spectroscopy (EDS), quantum mechanics simulations, solid state magic angle spinning nuclear magnetic resonance spectroscopy (MAS NMR), and pulsed-field gradient nuclear magnetic resonance spectroscopy (PFG NMR). At ambient temperature and RH, all of the H2O-coordinated interlayer cations, barring Al3+-intercalated MnPS3, exhibit “superionic”, or practically useful, bulk conductivity (>0.1 mS cm–1) and relatively low Ea. Notably, in the polycrystalline samples studied here, both σRT and Ea vary significantly depending on the identity of the intercalated cation. The ligand exchange to aprotic molecules generally results in decreased σRT and increased Ea, but achieving practically useful conductivity is still possible using ligands like MeCN.

Results and Discussion

Chemical and Structural Characterization after Ion Exchange

Elemental Analysis

MPS3 materials can undergo ion exchange processes as outlined in eqs 1 and 2. The ion exchanges have been well characterized in previous reports.45,52 However, the amount of M that is replaced by A, described by “x” in eqs 1 and 2, differs between various studies—ranging from 0.1 ≤ x ≤ 0.25.43,52,60,64 Intuitively, x can be controlled by varying the MPS3:ACln(aq) ratio, as illustrated in Figure S1. Here, we aimed for high A content (x ≃ 0.25 for CdPS3-based samples and ≃0.20 for MnPS3-based samples) to incorporate the largest number of charge carriers, which we hypothesize will lead to optimal conductivity. Table 1 shows the results of elemental analysis on the ion-intercalated MPS3 compounds. The amount of A intercalated (x) and remaining M (y) is determined using ICP-MS and is normalized to the measured P content. The measured S content is inaccurate due to H2S evolution during the sample digestion process but is also presumed to stay constant. The standard deviations for the ICP-MS data are representative of a minimum of five different synthesis batches per material. Due to synthesis and measurement inaccuracies the precise stoichiometry of the A2x/nCd1–xPS3 compounds sometimes contains ≤10% excess Cd, however, a comparison of “ideal” batches and those with excess Cd shows that there is no noticeable impact on the σRT and Ea (Figure S2). The H2O content is determined by a combination of TGA (Figure S3) and changes in pellet masses pre- and postdrying, measured using an analytical balance. The mass loss measured by TGA can be regained by re-equilibrating the material at the appropriate RH, suggesting that the TGA mass loss is measuring H2O loss For the rest of this paper, the compounds are referred to by their nominal formula for simplicity. The differences in the degree of ion exchange that occurs, both for different ions within a given framework and between the frameworks (CdPS3 vs MnPS3), are discussed in Supplementary Note 1.

Table 1 Stochiometry of Each A2x/nMyPS3 Compound, x and y Are Measured by ICP-MS and Normalized to P, Which Is Assumed To Be Constanta

compound	x	y	H2O/f.u.	H2O/A	nominal formula	
KxCdyPS3	0.50 ± 0.02	0.79 ± 0.02	1.0 ± 0.1	2	K0.5Cd0.75PS3·H2O	
LixCdyPS3	0.47 ± 0.01	0.85 ± 0.03	2.0 ± 0.7	4.4	Li0.5Cd0.75PS3·2H2O	
NaxCdyPS3	0.51 ± 0.02	0.81 ± 0.02	2.0 ± 0.5	3.9	Na0.5Cd0.75PS3·2H2O	
MgxCdyPS3	0.24 ± 0.01	0.78 ± 0.05	1.9 ± 0.0	7.9	Mg0.25Cd0.75PS3·1.9H2O	
CaxCdyPS3	0.24 ± 0.02	0.81 ± 0.02	2.0 ± 0.5	8.3	Ca0.25Cd0.75PS3·2H2O	
ZnxCdyPS3	0.41 ± 0.09	0.76 ± 0.24	0.25	0.63	Zn0.4Cd0.6PS3·0.25H2O	
AlxCdyPS3	0.13 ± 0.01	0.81 ± 0.08	2.5 ± 0.6	14.5	Al0.17Cd0.75PS3·2.3H2O	
KxMnyPS3	0.40 ± 0.05	0.80 ± 0.03	0.8 ± 0.2	2.1	K0.4Mn0.8PS3·0.8H2O	
LixMnyPS3	0.34 ± 0.03	0.84 ± 0.03	1.4 ± 0.2	4.0	Li0.4Mn0.8PS3·1.4H2O	
NaxMnyPS3	0.59 ± 0.08	0.71 ± 0.04	1.7 ± 0.6	2.7	Na0.6Mn0.7PS3·1.7H2O	
MgxMnyPS3	0.19 ± 0.01	0.82 ± 0.01	1.8 ± 0.1	9.5	Mg0.2Mn0.8PS3·1.8H2O	
CaxMnyPS3	0.25 ± 0.02	0.76 ± 0.02	1.5 ± 0.1	6	Ca0.25Mn0.75PS3·1.5H2O	
ZnxMnyPS3	0.46 ± 0.06	0.62 ± 0.06	1.6 ± 0.1	2.6	Zn0.4Mn0.6PS3·1.6H2O	
AlxMnyPS3	0.12 ± 0.00	0.84 ± 0.01	2.4 ± 0.4	16	Al0.13Mn0.8PS3·2.2H2O	
a H2O content of each compound, which is determined with a combination of TGA and measured mass pre- and postdrying using an analytical balance. The amount of H2O per intercalated ion in each material and the nominal formula of each compound were based on the measured cation content.

Interlayer Expansion after Ion Exchange

The changes to the crystal structure of the MPS3 materials after ion exchange are investigated by using XRD and Raman spectroscopy. Previous reports of ion or molecule intercalated MPS3 materials, and other hydrated layered compounds like clays and hydrated MX2 phases, largely focus on the interlayer spacing increases driven by the coordinating ligands after ion exchange.28,36,45,47,53,57,58,69,71−73,92 To the best of our knowledge, a complete structural solution has not been provided for any of the ion-intercalated MPS3 phases, likely due to the low crystallinity or structural complexity due to the disordered, mobile interlayer cations and ligands. However, some reports have shown that the ion-exchanged compounds can be indexed to a monoclinic crystal structure like the pristine material, but with an increased layer spacing,45,58 and have suggested that the structure within the metal layers is largely maintained after ion exchange.52,58,69

Figure 2a,b shows the low Q regions of representative XRD patterns of ion-exchanged CdPS3-based and MnPS3-based compounds, respectively. Here, we also focus on the changes to the layer spacing after intercalation as we expect this to have the most significant impact on the ionics, which is the primary focus of this study. We perform Le Bail analysis on the XRD patterns (examples shown in Figures S4 and S5) and find that the resulting materials can be indexed to monoclinic or orthorhombic phases. The significantly larger basal spacing, determined by the d(001) reflection, compared with the parent phases in accordance with previous reports. Some materials show a peak at 0.6 Å–1 that is difficult to index with a Le Bail fit, suggesting that it is associated with a small impurity. A complete structural determination is difficult due to the presence of small impurity phases, remnant parent phases, and disorder of the dynamic coordinated ions in the interlayer. The Raman spectra of the CdPS3-based compounds (Figure S6) support that the polyanion remains intact and its environment in the metal layer is largely unchanged, which supports the structural similarity of the metal layer before and after intercalation. Furthermore, we conduct a computational structural analysis to begin visualizing what the local order in the interlayer may look like (vide infra), and we hope this will spur further experimental studies into the structure of these complex materials. The full synchrotron XRD patterns are provided along with the other relevant data and are shown in Figure S7.

Figure 2 Characterization of the A2x/nM1–xPS3·yH2O materials using powder XRD. XRD patterns of (a) CdPS3-based and (b) MnPS3-based compounds. (c) d(001) spacing of A2x/nM1–xPS3·yH2O. The error bars reflect the standard deviation of several XRD measurements, the differences are due to the range of ambient humidity (≈40 to 55%). The gray regions highlight materials with a monolayer of H2O and the purple regions highlight materials with a bilayer of H2O. Note, the pattern of Zn0.4Mn0.6PS3·1.6H2O is collected using Cu Kα radiation because the sample inadvertently dehydrated during preparation for the synchrotron measurement.

Figure 2c shows the extracted d(001) values for all of the compounds. Error bars are shown to reflect the deviation associated with fluctuating RH at the time of replicate measurements. The original (001) reflection of the pristine MPS3 phase occurs at Q = 0.95 and 0.97 for M = Cd and Mn, respectively. The expanded (001) reflection of the ion-exchanged materials appears between Q = 0.49 and 0.67 Å, and the subsequent (002) reflections can be seen around Q ≃ 1.05 Å. Layer expansion suggests incorporation of the ions into the van der Waals gap, and SEM confirms that the layered platelet morphology of the particles is maintained. EDS confirms that the intercalated ions are homogeneously dispersed within the particles (Figure S8).

It is clear from Figure 2c that the d(001) spacing after ion exchange can be grouped into two main categories. The layer spacing of the hydrated structures increases by either ∼2.8 or ∼5.6 Å. These expansions correspond to the van der Waals radius of one or two H2O molecules, respectively. Therefore, the layer spacing increase is likely due to the formation of either a monolayer or bilayer of H2O around the intercalated cations in the van der Waals gap rather than following the expected hydrated cation radii of the inserted ions. The formation of hydrated MPS3 materials containing mono or bilayer H2O has been noted in previous studies on monovalent cation-intercalated MPS3 materials.45,52,60 However, to the best of our knowledge, the hydrated structure of MPS3 phases containing other ions has not been explored.

The occurrence of mono or bilayer H2O structures with the other ions is similar to the results in clays, transition metal dichalcogenides (MS2), or MXenes, which also exhibit hydrated structures containing mono or bilayer H2O with a wide range of interlayer cations.28,67,70−73

Whittingham, in addition to Lerf and Schöllhorn, independently investigated the hydration of interlayer ions in MS2 compounds, (M = Ti, Nb, Ta).71,72 These studies showed that the resulting structure of the hydrated compounds is governed by cation-ligand interaction. Specifically, the ability of the cation-ligand interaction offsets the loss in electrostatic lattice energy incurred by separating the cations from the anionic sulfide layers. Lerf and Schöllhorn found that this behavior can be described empirically by the charge/radius ratio (e/r) of the interlayer cations, which is correlated with the hydration energy. These rules translated well to other layered hydrated structures, like MXenes.73 Cations for which e/r < 1 (large and low charge, e.g., K+) can only stabilize a monolayer of H2O, whereas if e/r > 1 (small and high charge, e.g., Li+ and multivalents) then a bilayer of H2O can be stabilized. In Figure 2c, we show the e/r value for each mobile ion and confirm the relationship between e/r and the layer spacing in the ion-intercalated MPS3 materials. The e/r of Na+ is ≃1, suggesting that it can stabilize either a monolayer or bilayer of H2O. Indeed, isolated forms of monolayer or bilayer H2O, or even a mixture of both, can be achieved when A = Na+ depending on the RH (Figure S9). Furthermore, the TGA and derivative thermogravimetry (DTG) results (Figure S3) show that Na0.5Cd0.75PS3·2H2O is the only CdPS3-based compound with two distinct H2O loss events, corresponding to the transition between the stable monolayer and bilayer H2O structures. The H2O content of materials containing multivalents, e.g. Ca0.25Mn0.75PS3, changes as the RH is varied but a phase transition to a structure with monolayer H2O is not observed for multivalent-intercalated compounds at any RH due to the higher hydration energies of the multivalent ions. The solvation shells of multivalent ions are likely incomplete at low RH, but the overall bilayer structure is maintained (Figure S10). This also agrees with the findings of Lerf and Schöllhorn for MS2 compounds.71 Notably, in contrast to the structure observed by Yu et al. in CdPS3-based membranes,64 bulk K0.5Cd0.75PS3 could not be forced into a bilayer H2O structure even at 96% RH (Figure S11), perhaps due to the more rigid constraints of the polycrystalline powders.

Structural Effects Associated with the Mobile Ion

Next, we consider the structural effects associated with (1) the identity of the mobile ion and (2) the CdPS3 vs the MnPS3 host structure. First, we consider a given MPS3 framework with different ions. As outlined above, of the ions selected for this study, only K+ stabilizes a hydrated structure with a monolayer of H2O, while the other ions stabilize a bilayer at ambient RH in both MPS3-based frameworks. Although two main categories of layer spacings exist as the mobile ion is changed, there is a measurable deviation of the layer spacing within the bilayer regime depending on the intercalated cation. We plot the charge density of the mobile ion for structures that reside in the bilayer regime vs the d(001) layer spacing (shown in Figure S12) and find that the layer spacing is inversely related to the charge density of the intercalated ions. More charge-dense ions, such as Mg2+, interact more strongly with the H2O ligands, creating shorter bonds and, therefore, a smaller total diameter of the hydrated cation complex. Al3+ is an outlier in this trend likely because the very high charge density of the cation attracts more H2O which expands the layer. In fact, the Al-containing material contains more mole equivalents of H2O (∼16:1 H2O:Al) then can be accommodated in a first shell coordination environment. However, the H2O does not behave as “free” water as evidenced by 1H NMR (vide infra).

Ion exchange with Zn2+ in both structures behaves differently compared with most of the other ions. In both host materials, reflections associated with the pristine MPS3 phase dominate the diffraction patterns, though a small peak associated with an expanded lattice is observed for the MnPS3 framework. We note that Zn is indeed incorporated into the materials (Table 1). For the CdPS3-based material, we interpret these data to mean that most of the Zn2+ prefers to occupy the vacancies in the metal layer, which would suggest that the material does not contain substantial ligand-coordinated Zn2+. Indeed, the Zn0.4Cd0.6PS3 material does not contain a significant amount of H2O (Table 1). The Zn0.6Mn0.4PS3, which shows a small expanded reflection, does contain some water due to the small ligand-coordinated Zn2+ content. Zn2+ is thermodynamically stable in the metal layer as evidenced by the ZnPS3, which is a stable endmember that is isostructural to CdPS3 and MnPS3.74

Difference in Structural Changes between Frameworks

Next, we discuss trends between the CdPS3 and MnPS3 frameworks. A comparison of the two frameworks reveals that generally, the CdPS3-based compounds have a larger d(001) than the MnPS3-based compounds, which is consistent with the pristine materials (Figure 2c). However, Mg0.2Mn0.8PS3·1.8H2O has a larger basal spacing than Mg0.25Cd0.75PS3·1.9H2O. The reasons for this should be explored in a dedicated structural study. Na0.6Mn0.7PS3·1.6H2O forms a two phase mixture of mono and bilayer H2O at ambient RH while only the bilayer is formed with the CdPS3 host at ambient RH. Even though the Mn analogue has more intercalated Na+, it absorbs less H2O at a given RH. This is likely because the smaller lattice of MnPS3 demands a larger energy cost for expansion, therefore a higher driving force for H2O absorption (i.e., higher RH) is required to stabilize the bilayer structure.

Density Functional Theory (DFT) Simulations

The computationally stable monolayer hydrated structure of K0.5Mn0.75PS3·H2O and bilayer hydrated structure of Na0.5Mn0.75PS3·2H2O are investigated using DFT simulations to gain further insight into the organization of H2O molecules and interlayer cations since the structure of the cations and ligated H2O has not been addressed in previous reports, and a complete structural solution was not developed in this study. Such information is difficult to obtain experimentally due to the disorder and high mobility of the interlayer cations and ligated H2O.

Figure 3a shows the relaxed structure of K0.5Mn0.75PS3·H2O and 3b shows the relaxed structure of Na0.5Mn0.75PS3·2H2O. The interlayer cation and H2O organization is highlighted and viewed down the c axis and down the a axis. In K0.5Mn0.75PS3·H2O, some of the H2O is bridging between two K+ ions, while the others are more isolated to one K+ ion. The isolated H2O ligands are hydrogen bonded to neighbors and located approximately within the same ab plane, supporting the experimentally measured d(001) spacing that suggests roughly a monolayer of H2O. Many other configurations of K0.5Mn0.75PS3·H2O are similarly stable and are shown and discussed in Figure S13. The variety of structures with similar energy suggests multiple possible configurations, or combinations thereof, likely coexist at RT and align with the idea of disordered mobile interlayer cations.

Figure 3 Computationally relaxed structures of (a) K0.5Mn0.75PS3·H2O and (b) Na0.5Mn0.75PS3·2H2O. The left images show the view down the c axis after removing the metal layer above the hydrated interlayer cations; the metal layer below is made partially transparent for clarity. The right images show the view down the a axis. Hydrogen bonds between H2O molecules are shown as dotted black lines. The H2O molecules are inaccurately depicted as being smaller than the interlayer cation for the sake of clarity.

On the other hand, the structure of Na0.5Mn0.75PS3·2H2O contains no bridging H2O ligands. The H2O coordinates to Na+ and participates in hydrogen bonding with nearby H2O ligands. The higher density of H2O molecules in Na0.5Mn0.75PS3·2H2O forces the H2O to distort in the c direction and results in more than a monolayer in the ab plane, again supporting the experimentally observed d(001) spacing that suggests bilayer H2O in the interlayer when two H2O per formula unit are introduced. The d(001) spacing of the monolayer and bilayer relaxed structures are 8.9 and 10.3 Å respectively, which are smaller than the XRD measured values of 9.3 and 12.0 Å for K0.5Mn0.75PS3·H2O and Na0.5Mn0.75PS3·2H2O, respectively, likely due to thermal expansion at RT compared to the calculated structure at 0 K. We note, however, that the diffraction patterns simulated from the DFT relaxed structures do not perfectly match the experimental synchrotron diffraction patterns likely due to many defects in the real materials, such as stacking faults, and significant disorder within the van der Waals gap.

Characterization with MAS NMR

We employ MAS NMR to directly probe the chemical environments of the H2O and some of the interlayer cations in A2x/nCd1–xPS3·yH2O. Performing MAS NMR measurements on MnPS3-based compounds is challenging due to the presence of paramagnetic Mn2+, however, we anticipate that the results of the CdPS3-based compounds can be extended to the MnPS3-based analogues.

Figure 4a shows 1H MAS NMR spectra of all A2x/nCd1–xPS3·yH2O materials zooming in on the centerband. The chemical shifts are provided next to the corresponding spectra and are listed in Table S2 along with the fwhm of the centerband. For the A2x/nCd1–xPS3·yH2O compounds, the 1H NMR spectra all contain a single resonance appearing within a very narrow range of resonant frequencies, indicating the presence of a single 1H environment or of an average 1H environment due to fast chemical exchange between available sites in these structures on the NMR time scale. The observed chemical shifts are within the range expected for 1H in H2O molecules (around 5 ppm)77−80 but are noticeably shifted and broader than the signal of “bulk” free H2O, which is characterized by a sharp peak at ∼4.75 ppm.81,82 The 1H chemical shift values trend with the charge density of the intercalated cation. Cations with higher charge density experience stronger interactions with the O in H2O and are thus stronger bases, resulting in more deshielded hydrogen nuclei. Additionally, the appearance of spinning sidebands in the 1H NMR spectra of A2x/nCd1–xPS3·yH2O materials suggests that H+/H2O molecules are bound to the cations and exhibit limited mobility. The full 1H MAS NMR spectra are magnified in Figure S14 to highlight the spinning sidebands.

Figure 4 MAS NMR study of A2x/nCd1–xPS3·yH2O materials. (a) 1H MAS NMR spectra of A2x/nCd1–xPS3·yH2O compounds. (b) 7Li spectra of Li0.5Cd0.75PS3·2H2O, (c) 23Na spectra of Na0.5Cd0.75PS3·2H2O, and (d) 27Al MAS NMR spectra of Al0.17Cd0.75PS3·2.3H2O. The metal MAS NMR spectra are collected on the materials at ambient RH (hydrated), and after drying under vacuum at 200 °C for ≈10 h (dried). * marks the spinning sidebands.

Notably, although divalent and trivalent ions have ∼8 and ∼16 H2O molecules per cation, respectively, the NMR spectra of these compounds suggest that this “excess” H2O is not free H2O. These compounds all exhibit d001 spacings that are consistent with bilayer H2O. Therefore, for the Mg2+ and Ca2+ compounds, we speculate that the arrangement resembles Figure 3b, but with half of the metal sites unoccupied. The empty interlayer metal sites result in contraction of the H2O molecules around the cation forming a sort of second shell within the H2O bilayer. Alternatively, the H2O could be arranged in a single shell of 8-coordinate cubic arrangements around each cation. Either of these interpretations is consistent with the ionic conductivity study of the Ca-intercalated MPS3 compounds (vide infra). This hypothesis can be extended to the Al0.17Cd0.75PS3·2.3H2O material, which has more H2O than can fit directly around each Al3+ ion in the bilayer structure in either a 4- or 8-coordinate arrangement, in this case, formation of a second shell of coordinated H2O is the most likely explanation for the lack of free H2O. Exchange of the H2O in the different solvation shells could result in averaging of the 1H environments into the broad signal that we observe for the Al compound in Figure 4a. In general, the chemical shift trending with the charge density and the lack of a resonance that resembles free water indicate that the H2O molecules in A2x/nCd1–xPS3·yH2O are bound to the intercalated cations, in agreement with the TGA results.

We also investigated the chemical environment of 7Li, 23Na, and 27Al in the respective ion-intercalated CdPS3-based compounds. Note, conducting metal cation NMR analyses on the K+, Mg2+, Ca2+, and Zn2+ materials is prohibitively challenging due to the low natural abundance of their NMR-active isotopes and quadrupolar effects. Figure 4b–d shows 7Li, 23Na, and 27Al MAS NMR spectra, respectively, in both hydrated and dehydrated states in order to probe the environment of these intercalated cations with and without H2O ligands. The comparison between the hydrated and dehydrated materials can be found below in the discussion section on dried A2x/nCd1–xPS3.

The chemical shifts for 7Li and 23Na are 0.1 and 1.34 ppm, respectively, which are consistent with previous reports of interlayer ions in bilayer H2O-containing hydrated structures of Li2Sn2S5 and Na-containing mica clays.36,83−85 The 23Na spectrum of the hydrated phase is consistent with previous 23Na MAS NMR studies on Na0.50Cd0.75PS3·2H2O.58,60 In the case of Al0.17Cd0.75PS3·2.3H2O, the 27Al MAS NMR spectrum contains a single resonance at −1.5 ppm, which is consistent with Al3+ that is octahedrally coordinated by H2O (∼0 ppm).86,87 In summary, the MAS NMR results further support the claim that the intercalated ions are located within the interlayer spacing and are coordinated by the H2O molecules.

Electrochemical Characterization of the Ionic Mobility

The elemental and structural analyses provide evidence that the H2O-coordinated ions are introduced into MPS3 frameworks and occupy the interlayer space. These results are consistent with previous work on alkali metal ion-intercalated MPS3 materials,43,45,52,58 and show that the concept extends to many other ions. Due to the increased interlayer spacing, which is directly related to the change in d(001) and likely results in a wider ionic conduction channel, as well as the screened Coulombic interactions between the interlayer cation and anion framework, we hypothesize that the intercalated ions should exhibit enhanced mobility.

To probe the ionics of the materials, the ionic conductivity and activation energy of all compounds are measured using EIS on cold-pressed pellets with a 6 mm diameter. The pellets are assembled in PTFE Swagelok cells with symmetric ion-blocking electrodes. For each material, EIS is measured at a range of temperatures from RT to 70 °C. Representative Nyquist plots for all compounds are shown in Figures S15 and S16, and an example of a fit using an equivalent circuit model for a sample where the high-frequency semicircle can be resolved is shown in Figure S17. The Ea is determined using the well-established Arrhenius-type relationship that governs the thermally activated ionic conduction process.4,75Figures S18 and S19 show average Arrhenius-type plots of ln(σT) vs T–1 for all studied ion-exchanged compounds. The σRT values and Ea values for all studied compounds are plotted in Figure 5a,b, respectively. Figure 5a,b shows the mean and standard deviations obtained from at least three replicate cells, each from a different synthesis batch, and these values are also listed in Table S1. In general, all of the compounds containing solvated ions in the interlayer show high total ionic conductivities at RT, and by 70 °C all exhibit practically useful conductivities >0.1 mS cm–1.

Figure 5 (a) Room temperature ionic conductivity (σRT) and (b) activation energy (Ea) of all ion-intercalated materials at ambient RH. Zn0.4Cd0.6PS3·0.25H2O is not shown here because the σRT is too low (10–9 S cm–1).

Detailed analysis of the data set enabled by this modular framework facilitates the development structure–property relationships for ligand-coordinated ion conduction in rigid solids.

Mobility of Various Ions in a Given Framework

The differences in performance between various ions in a given framework elucidate the effect of charge density and interlayer spacing size on mobility. It is evident that the ions with higher charge density generally exhibit a lower σRT and a higher Ea. Thus, the charge density of the mobile ion is a strong factor in determining its mobility, suggesting that the coordinating H2O is not completely screening the charge. In an aqueous solution, ions with greater charge density attract more H2O molecules to screen the charge. However, within a rigid solid framework, there is a limit to the number of H2O molecules that can surround the ion due to spatial constraints. For example, Mg0.25Cd0.75PS3·1.9H2O and Ca0.25Cd0.75PS3·2H2O both contain bilayer H2O and both have ∼8H2O per cation. If the mobile ions migrate with the bound H2O via a “vehicular mechanism” (discussed in detail later), the mobile complex to be considered is the An+·yH2O species. In this case, assuming that each H2O molecule provides the same degree of charge screening, the mobile complex Mg2+·8H2O is more charge-dense than the Ca2+·8H2O complex. Alternatively, if the mobile ions conduct via a hopping mechanism between H2O molecules, the charge density differences between the cations must still be considered.

Furthermore, the charge density of the intercalated ion will affect the size of the hydrated complex and, therefore, also change the interlayer spacing. As mentioned previously, the migration channel in these materials is likely dependent on interlayer spacing. Mg2+–OH2 bonds are stronger than the Ca2+–OH2 bonds, as indicated by the higher peak mass loss temperature observed in TGA and DTG analyses (Figure S3). The stronger bond in the Mg2+·8H2O complex pulls the H2O molecules closer to the cation, leading to less layer expansion and thus a more narrow interlayer channel in Mg0.25Cd0.75PS3·1.9H2O compared to Ca0.25Cd0.75PS3·2H2O. Again, Al3+ is the outlier to this trend due to the excess water absorption, as discussed previously.

The combination of the higher charge density of the mobile complex and a narrower interlayer spacing explains the lower mobility of the charge-dense ions. These trends hold when comparing Na+ and K+ in the monolayer structures as well (Figure S20).

Ionic Mobility Differences between Frameworks

In most cases, the performances of a given ion between frameworks are similar (K+, Li+, Mg2+, and Al3+). In these cases, slight differences in performance between the two frameworks can be understood in the context of interlayer spacing size. The MnPS3-based Zn-intercalated compound exhibits significantly better performance than the CdPS3 analogue, while the MnPS3-based compounds for Na+ and Ca2+ ions both perform notably worse than the CdPS3 analogue. Exploring these pronounced differences between frameworks allows us to highlight the impact of ligand-coordinated interlayer ions, the hydration state, and charge screening by coordinated ligands.

Effect of the Interlayer Spacing

In the cases where the frameworks have a similar performance for a given ion, K+, Li+, Na+, Mg2+, and Al3+, the subtle differences can be explained by considering the interlayer spacing size. In the framework with the larger interlayer spacing, the σRT is slightly higher and Ea is slightly lower. As discussed previously, typically the CdPS3-based framework has the larger interlayer spacing and thus we observe that most ion-exchanged CdPS3 phases have higher σRT and lower Ea than the MnPS3 analogues. The exception is again the Mg2+ intercalated compounds, which show the opposite trend in the d(001) spacing. This is reflected in the higher σRT and lower Ea for the MnPS3 host compared to those of CdPS3. The observed trend in performance highlights the impact of the interlayer spacing on ion mobility, with larger interlayer spacings facilitating improved ion transport. The observed correlation of the interlayer spacing to the conductivity suggests that it plays an important role in defining the conduction pathway for the mobile ions in these materials.

Effect of Ligand-Coordinated Interlayer Ions

The σRT of Zn0.4Cd0.6PS3 is on the order of 10–9 S cm–1 (not shown in Figure 5), which is 4 orders of magnitude lower than that of Zn0.6Mn0.4PS3. The superior performance of Zn0.6Mn0.4PS3 is due to the presence of some hydrated Zn2+ ions in the van der Waals gap, whereas in Zn0.4Cd0.6PS3 the Zn2+ ions are immobilized in the metal-layer lattice sites. This illustrates that the solvated interlayer cations are crucial to achieving high ionic conductivity.

Effect of Hydration State

The disparity in Na+ conductivity can be explained by the fact that Na0.6Mn0.7PS3·1.6H2O is a two-phase mixture of the monolayer and bilayer hydrated structures. The monolayer structure will have a much lower Na+ mobility than the bilayer structure due to the significantly decreased interlayer spacing size and less effective charge screening (this is discussed in detail in the SI). The performance of Na0.6Mn0.7PS3·1.6H2O reflects the combined properties of both hydration states. By equilibrating the material at 75% RH, the bilayer structure (Na0.6Mn0.7PS3·2H2O) can be isolated. Na0.6Mn0.7PS3·2H2O exhibits a conductivity more similar to that of Na0.5Cd0.75PS3·2H2O as expected (Figure S20).

Effect of Charge Screening with Ligand Molecules

Similarly to the Na-containing frameworks, Ca0.25Mn0.75PS3·1.5H2O absorbs less H2O than Ca0.25Cd0.75PS3·2H2O at ambient RH. However, despite the lower H2O content, the layer spacing associated with the bilayer phase is maintained and does not collapse into the monolayer structure. Instead, we assume the solvation shells around the Ca2+ ions in the MnPS3 host are incomplete relative to the CdPS3 host at ambient RH. The maintenance of a bilayer H2O structure despite H2O loss allows us to isolate the effect of charge screening from the impact of the interlayer spacing size. Clearly, the difference in H2O content and solvation significantly impacts the ionic conduction, as evidenced by the 25% higher Ea in Ca0.25Mn0.75PS3·1.5H2O (362 meV) compared to Ca0.25Cd0.75PS3·2H2O (287 meV). Notably, the Ca0.25Mn0.75PS3·1.5H2O also contains 25% less water compared to Ca0.25Cd0.75PS3·2H2O. Therefore, the Ea increase and corresponding σRT decrease in the Mn-based framework is attributed to the less effective screening of Ca2+ in Ca0.25Mn0.75PS3·1.5H2O due to the lower H2O per cation. This highlights the crucial role of charge screening in facilitating high ionic mobility in MPS3 frameworks.

Identification of the Mobile Ion

Though we can draw several structure–property relationships using the range of materials discussed above while assuming that the charge carriers are primarily the hydrated cations, the use of H2O as the coordinating ligand introduces the possibility of H+ conduction. It has long been debated whether the ionic conductivity of ion-intercalated hydrated clays is a result of primarily H+ conduction or migration of the intercalated ions.28−30,68 Therefore, care must be taken to identify the majority charge carrier in A2x/nM1–xPS3·yH2O materials.

Mobile H+ in A2x/nM1–xPS3·yH2O could potentially arise through either surface acidity,76 or hydrolysis of H2O bound to intercalated cations that act as Lewis bases.29,30 Since S-based materials tend to adsorb less surface H2O than O-based materials, due to weaker hydrogen bonds between H2O and S, surface acidity likely plays a less significant role in MPS3-based frameworks than in clays. To estimate the contribution of H+ from surface acidity to the total conductivity, we investigate the ionic conductivity of pure MPS3 phases under ambient relative humidity (the EIS results are shown in Figure S21). The σRT of MnPS3 and CdPS3 are 3 × 10–9 and 4 × 10–10 S cm–1, respectively, which is 4 or 5 orders of magnitude lower than that of the worst performing A2x/nM1–xPS3·yH2O compounds. Therefore, the observed ionic conductivity in A2x/nM1–xPS3·yH2O materials is likely either due to mobile H+ from the acidity of the ion-solvating H2O or due to the conduction of the interlayer cations themselves.

Several characteristics of A2x/nM1–xPS3·yH2O suggest that the intercalated cations are the majority charge carriers. The fact that the ions can be inserted and removed by ion exchange in an aqueous solution indicates their inherent mobility within the structure. Additionally, the strong correlation between the Ea and σRT with the identity of the intercalated ion underscores the importance of these ions on charge transport. The Ea for H+ conduction through a Grotthuss-type mechanism, where H+ is being exchanged by neighboring H2O molecules, is typically <400 meV.63,65 The variation in Ea measured in the materials reported here is between 216 and 622 meV, which is more consistent with the conduction of mobile complexes of different charge densities, as discussed above. The observed trends in σRT also do not align with H+ being the majority charge carrier. The H+ concentration would depend on the acidity of the absorbed H2O, which is influenced by the charge density of the intercalated cation. Therefore, the trend in pKa of the hydrated cations can be used to predict the trend in H+ concentration in A2x/nM1–xPS3·yH2O. In general, the more charge-dense the cation, the more acidic the bound H2O will be due to A–O bond strength. The pKa trend of the intercalated ions is overlaid onto the σRT data in Figure S22. The A2x/nM1–xPS3·yH2O compounds with more acidic interlayer ions (e.g., Mg2+, Zn2+, Al3+) should have the highest H+ concentration but show the lowest measured ionic conductivities. Thus, the observed electrochemical performance is not adequately explained by assuming that H+ is the majority charge carrier.

Investigation of Ionic Conduction with PFG NMR

To quantify and compare the relative contributions from H+ and Li+ species to the ionic conductivity observed in a representative member of this new class of superionic conductors, the self-diffusion coefficients and Ea of Li and H species in Li0.5Cd0.75PS3·2H2O are measured using 7Li and 1H PFG NMR. The normalized echo signal attenuation data obtained between 25 and 60 °C are analyzed using both the Tanner–Stejskal equation for 3D diffusion,88,89 and a 2D diffusion model described by Stoll and Kimmerle et al.90,91 for comparison. For both species, a biexponential fit using the Tanner–Stejskal equation provided the best results (Figure S23), as discussed in more detail in Supplementary Note 2. Briefly, the 2D diffusion model90,91 was quickly eliminated (despite individual Li0.5Cd0.75PS3·2H2O exhibiting 2D diffusion) because it led to a poor fit of the 1H data, while the 7Li data could be fit with this model the results were nonsensical based on the layered structure of this material, as the 2D model suggests the out of plane diffusivity (in the c direction) is an order of magnitude higher than the in plane diffusivity (ab plane). The suitability of the biexponential model using the Tanner–Stejskal equation provides strong evidence for the presence of two distinct 1H and 7Li diffusing species in these materials, which are not significantly affected by 2D confinement in the layered structure. We tentatively attribute those two 1H/7Li diffusion environments to ion transport at the grain boundaries and in the bulk. Hatz et al. also found that 7Li PFG NMR data collected on lithium tin sulfide nanosheets could only be fit using a biexponential 3D diffusional model and the Tanner–Stejskal equation.77

Figure 6 shows Arrhenius-type plots of ln(D) vs 1000/T for both the Li- and H-containing species, where D is the diffusivity of the diffusing components. The RT diffusivities of the Li components are 1.62 × 10–10 and 3.09 × 10–11 m2 s–1, which is associated with 57 and 43% of the Li species measured, respectively. The diffusivities of the H species are 1.13 × 10–12 and 1.55 × 10–14 m2 s–1, which is associated with 18 and 82% of the H species, respectively. The second, low diffusivity component is not shown in Figure 6b. Since the Li diffusivity is a few orders of magnitude higher than that of H, we can deduce not only that Li are the majority charge carriers but also that the Li+ ions conduct through a hopping mechanism between H2O molecules, instead of a vehicular mechanism in which Li+ conducts with its solvation shell. Interestingly, this mechanism is in contrast to that suggested in the case of hydrated Li2Sn2S5 in which the Li and H diffusivities were comparable, leading to the conclusion of a vehicular ionic conduction mechanism.36

Figure 6 Arrhenius-type relationships of the diffusivity (D) measured with (a) 7Li and (b) 1H PFG NMR. The second component associated with the H-containing species is not shown due to the low diffusivity.

The Ea values are obtained from the Arrehnius-type plots in Figure 6. The PFG NMR determined Ea for Li diffusion is 180 and 150 meV for components one and two, respectively. These Ea values are consistent with that measured by EIS (217 ± 14 meV). The Ea measured for the H-containing component is even lower than that for the Li components at 120 meV, though we note this value is more of an estimate due to the low measured diffusivity. Though the Ea for the H-containing component is low, the much lower diffusivity suggests that the majority charge carrier measured with EIS is Li.

The observed correlation between electrochemical properties and the intercalated ions and the significantly lower mobility of H vs Li as evidenced by PFG NMR strongly suggests that the ionic mobility observed in these compounds is predominantly due to the intercalated ions.

Structural and Electrochemical Characterization of Dried A2x/nCd1–xPS3

Although the primary charge carriers are likely the intercalated ions rather than mobile H+, the H2O ligands play a critical role in enabling high conductivity. We use MAS NMR and XRD to investigate the local and long-range structural changes that occur after drying in a vacuum oven at T ≥ 100 °C for at least 10 h, and EIS to probe the electrochemical performance.

The XRD patterns of A2x/nCd1–xPS3 pre- and postdrying are shown in Figure S24. Lattice contraction is observed after drying all samples. For the majority of the dry samples, the basal spacing resembles that of the pristine CdPS3 phase. Only dried K0.5Cd0.75PS3 and Na0.5Cd0.75PS3 materials display a larger basal spacing than does the pristine CdPS3 material.

The removal of the majority of water in the dried A2x/nCd1–xPS3·yH2O compounds is confirmed by the significant reduction of the 1H signal (an example is shown in Figure S25).

Without H2O, the intercalated cations are destabilized within the van der Waals gap, causing them to occupy vacant M sites. This has been shown previously for Li+-intercalated MnPS3,52 as well as A3+ ion-intercalated CdPS3.92 For monovalent ions, there are twice as many intercalated ions as there are vacancies in the metal layer; therefore, at least half of the ions may remain trapped in the van der Waals gap. In the case of K-intercalated MPS3, the K+ is too large to occupy the vacant sites in the metal layer, and therefore all of the K+ likely remains in the van der Waals gap. This is evidenced by the lack of a reflection resembling the pristine material lattice spacing in dried K0.5Cd0.75PS3. In the case of dried Na0.5Cd0.75PS3, the XRD pattern shows a reflection corresponding to the spacing of pristine CdPS3 in addition to a reflection indicating a slightly expanded lattice. The 23Na MAS NMR of dried Na0.5Cd0.75PS3, shown in Figure 4c, contains two distinctive resonances, which could be assigned to even distributions of Na in two different chemical environments. The 23Na spectrum of dried Na0.5Cd0.75PS3 strongly resembles the 23Na spectrum of Na4P2S6.93 The −0.7 ppm resonance likely corresponds to the interlayer Na+ and the 27.7 ppm resonance is attributed to the Na+ that occupies the metal layer. These values are slightly different than Na4P2S6 (5.6 and 18.2 ppm respectively) due to the structural and chemical changes imparted by the Cd ions in Na0.5Cd0.75PS3.

Although 0.25 equiv of Li+ likely remains in the van der Waals gap in dried Li0.5Cd0.75PS3,52 layer expansion is not observed because Li is small enough to occupy the interlayer without expanding the d(001) spacing, as is the case in Li4P2S6. However, typically 7Li MAS NMR does not permit differentiation between the different chemical environments.94−96

For multivalent ion-intercalated samples, no layer expansion is observed in the dried materials because the number of intercalated ions is less than or equal to the number of vacancies; therefore, all of the ions occupy the metal layer. As evidenced in Figure 4d, after drying, the 27Al resonance shifts to 36 ppm, which is more reminiscent of Al octahedrally coordinated by S.87

The dried multivalent-intercalated compounds do not exhibit any meaningful ionic conductivity (<10–9 S cm–1); however, the remnant interlayer monovalent ions in the dried monovalent-intercalated compounds enable the study of metal cation mobility in MPS3 frameworks in the absence of solvating ligands.

Figure 7a shows the basal spacing of A2x/nCd1–xPS3 (A = Li, Na, K) pre- and postdrying. In contrast to the hydrated phases, where the basal spacing is determined first by the amount of H2O absorbed and second by the charge density of the ion and by extension the size of the hydrated complex, the d(001) spacing of the dried phases trends with the intercalated ionic radius. In the absence of the interlayer expansion and charge screening, the σRT drops by up to 6 orders of magnitude, and the Ea increases up to threefold. In this case, the trend in electrochemical behavior is in line with what is expected of traditional solid state electrolytes. Li0.5Cd0.75PS3 exhibits the highest conductivity, due to the inherent mobility of the small, low charge density Li+ cation. Notably, this is consistent with the MAS NMR data, wherein the hydrated state of the 7Li and 23Na MAS NMR spectra of the respective samples both show relatively narrow line widths and an absence of spinning sidebands, suggesting that the cations are sufficiently mobile to average out anisotropic interactions. After drying, the 7Li and 23Na MAS NMR spectra show broadening of the signals and emergence of spinning sidebands, indicating reduced mobility.

Figure 7 (a) d(001) spacing of monovalent ion-intercalated CdPS3 in the hydrated and dried state, and (b) σRT and Ea of dried monovalent ion-intercalated CdPS3.

Ligand Exchange of A2x/nM1–xPS3·yH2O

The absorbed H2O ligands are vital for enabling high ionic mobility. However, in many battery systems, H2O can cause undesired side reactions that impair performance. To extend the applicability of these materials to nonaqueous systems, the H2O can be replaced with more stable organic molecules. Further, the ligand exchange helps to elucidate the effect of the ligands on ionic conduction, specifically by comparing the impact of widening the interlayer and the degree of charge screening. As a proof of concept, we exchange H2O ligands for other solvent molecules that are frequently used in next-generation battery liquid electrolytes: MeCN and THF. These molecules cover a range of dielectric constants, as shown in Figure 8a, allowing us to further probe the effect of charge screening.

Figure 8 (a) Bar chart showing the reported dielectric constants of H2O, MeCN, and THF, and comparisons of (b) d(001), (c) σRT, and (d) Ea of K0.5Cd0.75PS3· Z (Z = H2O, MeCN, or THF).

The solvent exchange is done by first exposing K0.5Cd0.75PS3 to a vacuum to at least partially dry the material followed by stirring the K0.5Cd0.75PS3 in dried solvent in a glovebox for 1 h. To confirm that the organic solvent is incorporated, Raman is measured on the resulting materials. In all cases, the Raman modes associated with the new ligand are observed (Figure S26). To further confirm ligand exchange, we analyze the basal spacing after the exchange, which is shown in Figure 8b (XRD patterns available in Figure S27). The basal spacing increases in the order of H2O < MeCN < THF, correlating with the size of the ligand molecules. The increase in basal spacing with THF (4.5 Å) approximately matches the radius of a flat THF molecule (4.2 Å).97Figure 8c,d shows the σRT and Ea, respectively, of K0.5Cd0.75PS3 ·Z (Z = H2O, MeCN, THF). Replacing H2O with MeCN or THF decreases the σRT by 1 or 3 orders of magnitude and increases the Ea by ∼70 or 200 meV, respectively. Given that the framework and mobile ion remain constant, the variation in mobility is attributed to the interplay between an increase in interlayer spacing size and a decrease in the extent of charge screening. Despite THF creating the largest interlayer spacing, the THF-containing compounds show low σRT and high Ea due to either ineffective charge screening by THF or a channel that is effectively too large. However, the conductivity with absorbed THF is still 2 orders of magnitude higher than dried K0.5Cd0.75PS3. Although MeCN absorption results in a smaller interlayer spacing than THF, its superiority in terms of screening the mobile ion charge leads to a significantly higher σRT and lower Ea. The variation in the dielectric constant of the selected ligands (53% decrease for MeCN, 90% decrease for THF) is larger than the resulting variation in interlayer spacing size (5% increase for MeCN, 18% increase for THF). Therefore, variation in screening ability plays the dominant role in governing the mobility. Consequently, K0.5Cd0.75PS3·MeCN exhibits a high ionic conductivity of 0.18 mS cm–1 in the absence of H2O, enabling potential application in nonaqueous systems. In addition, achieving high ionic conductivity with an aprotic ligand provides further evidence that the majority of charge carriers are the interlayer cations and not H+. Future work will explore the integration of aprotic solvents with high dielectric constants into MPS3 frameworks with various next-generation mobile ions.

Ionic Conduction in A2x/nM1–xPS3 Compared to Previous Reports

Now that we have discussed in-depth the structure–property relationships associated with ionic conduction in the A2x/nM1–xPS3 materials, we will next discuss our work in the context of similar studies.

In the late 90s, Jeevanandam et al. measured ionic conductivity using admittance measurements of K0.5Cd0.75PS3·H2O and Na0.5Cd0.75PS3·2H2O.58,59 The studies concluded that Na+ is mobile while K+ is not. However, the EIS analysis presented here conclusively shows very high ionic conductivity for K+ in K0.5Cd0.75PS3·H2O (Nyquist plots can be seen in Figure S15). Additionally, we find that the σRT of Na0.5Cd0.75PS3·2H2O is 2 orders of magnitude higher than in their previous reports. These discrepancies may arise from sample preparation or EIS measurement errors, but it is difficult to pinpoint the exact cause by using the experimental conditions that were reported.

The recent study by Yu and Ren, in which exceedingly high 2D ionic conductivity was observed in CdPS3-based membranes, provides an opportunity to compare solvent-assisted ionic conductivity to that of confined liquid electrolytes. In the study by Yu and Ren, there is no discussion of the quantity of H2O present in each membrane. However, since the basal spacing, σRT, and Ea are largely uncorrelated to the identity of the intercalated ion, reminiscent of a confined liquid electrolyte we speculate that the membranes likely contain excess amounts of H2O. In the case of ligand-assisted ionic conduction in bulk A2x/nM1–xPS3 materials, although the intercalated ions are screened they still interact with the framework, explaining the strong dependence of the electrochemical performance on intercalated ion identity. A comparison of the results of their 2D σRT and Ea with the findings of this study is shown in Figure S28. The measured conductivity of CdPS3-based membranes is between 2 and 4 orders of magnitude higher than the bulk polycrystalline A2x/nM1–xPS3 samples reported here. This is logical because the inherent conduction channels in this structure are 2D (within the interlayer); conduction throughout the bulk is achieved in polycrystalline pellets through a series of 2D conduction processes within randomly oriented particles. However, particularly for energy storage applications considering general materials processing and cell assembly steps, the bulk pellet measurement is more representative of the materials' performance. Furthermore, operating within the ligand-assisted regime combines the benefits of high ionic mobility with the advantages of using a solid state electrolyte. Additionally, there is limited free H2O in the material, which minimizes H2O-related undesired side reactions and limits the impact of H+ conduction. However, in the confined liquid electrolyte regime the materials may become softer and sticky due to excess H2O,36 and the free H2O can lead to undesired reactivity and increased H+ conduction in battery systems.

Conclusions

The introduction of ligand molecules into solids can drastically increase the ionic mobility of both larger cations (Na+, K+, Ca2+) and charge-dense cations (Mg2+, Zn2+, Al3+) by expanding the interlayer spacing, which likely widens the possible conduction pathways and screening charge-dense mobile ions. In MPS3-based materials, ion-intercalated compounds of the form A2x/nM1–xPS3·yH2O can be obtained containing hydrated interlayer A cations. These intercalated structures contain either a monolayer H2O, K+, and Na+ (at low RH); or bilayer H2O:Li+, Na+, Mg2+, Zn2+, Ca2+, and Al3+, depending on the hydration energy of the intercalated cation. All A2x/nM1–xPS3·yH2O materials show exceptionally high bulk conductivity (generally >0.1 mS cm–1) at RT and ambient RH. Therefore, leveraging ligand-assisted ionic conduction in solid state electrolytes is a promising avenue to achieving high ionic mobility of next-generation mobile ions, which has historically been very challenging, particularly for multivalent ions.36,38,39,98,99 Furthermore, A2x/nM1–xPS3·yH2O materials exhibit bulk ionic conductivities that are among the highest reported values of any electronically insulating inorganic solid for all of the mobile ions studied. Notably, these compounds exhibit higher conductivities than other reported solids containing solvated ions (e.g., clays and MOFs), due in part to the more polarizable S-based anion framework in MPS3-based materials. Employing ligand-assisted ionic conduction seems to be more effective for achieving exceedingly high conductivity for less charge-dense ions, such as Li+, Na+, K+, and Ca2+, but still provides suitable results for charge-dense ions.

This work highlights the impact of expanding the conduction channels by showing that with identical ions and ligands, the framework with the larger interlayer spacing has a higher σRT and lower Ea. Furthermore, the impact of charge screening is illustrated by the systematic increase in Ea and decrease in σRT for similar structures when some of the solvating H2O ligands are removed.

Since the introduction of H2O introduces the possibility of H+ conduction, we provide evidence that the intercalated cations are the majority carriers. The observed electrochemical behavior is much more adequately explained by mobile intercalated ions than by mobile H+. PFG NMR demonstrates that in the case of Li0.5Cd0.75PS3·2H2O the Li diffusivity is 2 to 4 orders of magnitude higher than that of H. This result further suggests that the intercalated cation is the majority charge carrier and that the conduction occurs through a hopping mechanism, where the cations are moving between H2O molecules, instead of a vehicular mechanism. Additionally, 7Li, 23Na, and 27Al MAS NMR show a narrower line width in the respective hydrated structures than the dried structures, indicating higher mobility of those cations in the former.

Although H+ ions are likely not mobile ions in the studied materials, the drastically inferior performance of the dried phases emphasizes that H2O ligands are crucial for the high mobility of the intercalated ions. Finally, exchanging H2O for aprotic ligands like MeCN and THF in K0.5Cd0.75PS3·H2O provides further evidence that H+ is not the majority carrier. The ionic mobility of the intercalated cations with these ligands depends on the interplay between conduction channel expansion and effective charge screening. Specifically, K0.5Cd0.75PS3·MeCN exhibits a high ionic conductivity, demonstrating the potential application of A2x/nM1–xPS3 to nonaqueous systems.

A2x/nM1–xPS3 materials represent a modular system in which we can change the mobile ion, the framework (M = Cd for slightly higher performance vs Mn for environmental friendliness), or the ligand molecules to tailor the performance for specific applications. Fundamentally, this modular framework allows us to develop structure-property relationships to better understand the solid state ionic conduction of next-generation mobile ions, particularly ligand-assisted ionic conduction. Additionally, the high performance of A2x/nM1–xPS3 materials suggests that these, or similar compounds, could be used as a “universal” solid electrolyte for a variety of battery chemistries with different mobile ions.

Experimental Methods

Material Preparation

Synthesis

The MPS3 (M = Mn, Cd) materials were prepared using traditional solid state methods from Mn (Alfa Aesar, 99.3%) or Cd (Thermo Scientific, 99.99%) metal powder, elemental S (Acros Organics, >99.5%), and 10% excess P2S5 (Acros Organics, >98%) in an Ar-filled glovebox without further purification.

The M, P2S5, and S8 were combined in a 2:1.1:1/8 molar ratio and ground thoroughly using a mortar and pestle. The reactants were then pressed into pellets with an Arbor press and sealed in a vitreous silica ampule under vacuum (<10 mTorr). The reaction vessel was placed in a box furnace, heated to 650 °C at a rate of 1 °C min–1 (K min–1) and allowed to react at 650 °C for 24 h. After the reaction was complete, the tube was allowed to cool to ambient temperature inside the furnace. The resulting green (MnPS3) or off-white (CdPS3) powder was collected and handled in an Ar filled glovebox.

Ion-Exchange Reaction

The ion-exchange reactions were conducted by stirring powder MPS3 samples in an aqueous solution of the appropriate metal chloride. A typical K-exchange reaction involved stirring 400 mg of MnPS3 or 500 mg of CdPS3 in 10 mL of a 3 or 2 M solution, respectively, of aqueous KCl at room temperature (RT) for 3 h. For CdPS3, 0.1 M EDTA, in a 1 M K2CO3/KHCO3 buffer solution, is added as a complexing agent. After the reaction was complete, the mixture was filtered using a fritted glass vacuum filter, washed three times with water and once with ethanol, and allowed to dry for at least 30 min. The second ion exchanges were conducted on the material obtained from the first exchange. In this case, 100–150 mg of K-exchanged MPS3 was added to 10 mL of 1 M solutions of the relevant metal chloride, and no complexing agents were added. The drying and washing procedures were the same as the first exchange. The Li-exchanged samples were allowed to dry overnight. After the ion exchanges were complete the samples were stored in vials at ambient conditions. A hygrometer was used to measure the ambient RH, if there was a period in which the humidity was <40%, the materials were stored in a humidity chamber maintained at 53% RH using super-saturated solutions of magnesium nitrate.

Ligand-Exchange Reaction

The ligand exchange was carried out in a dry N2-filled glovebox. Samples of K0.5Cd0.75PS3·H2O were partially dried through three 10 min vacuum-backfill cycles in the glovebox antechamber. MeCN (99.9%, Fisher Scientific) and THF (99.9%, Fischer Scientific) were dried on a solvent purification system (Pure Process Technology) and transferred into the glovebox, without exposure to air, and stored over 3 Å molecular sieves. Before use, the measured water content of both solvents was less than 20 ppm via KF titration. In the glovebox, K0.5Cd0.75PS3 was added to 5 mL of the solvent in a scintillation vial and mixed on a magnetic stir plate for 1 h. The mixture was vacuum filtered over a fritted glass filter for a couple of seconds until the material visibly changed from its dark gray “wet” state to the drier light gray state. Quickly, the material was scraped off of the filter and put into Swagelok cells for impedance characterization or an empty scintillation vial for absorbed solvent characterization.

Absorbed solvent was characterized by mass loss measurements. An initial mass was obtained from the sample immediately following removal from the filter. This sample was then allowed to passively lose solvent in a dry atmosphere.

Dried Sample

To prepare dried ion-exchanged materials for structural and electrochemical characterization, the powders were transferred into a Ar-filled glovebox and placed in a vacuum oven at 120 °C for 10 h.

Pelletization

Between 15 and 30 mg of powder was pressed into a pellet in an ambient atmosphere using a 12 ton hydraulic press from Carver (unit 3912). The powder was pressed using a 6 mm stainless steel die set at 2 tons for 5 min. The resulting pellets were between 0.2 and 0.6 mm thick. In most cases, the pellets were sputtered with Au at 40 mA for 60 s on both flat surfaces using a Ted Pella 108 Auto Sputter Coater in an Ar filled glovebox. Then, the mixture was removed from the glovebox and allowed to re-equilibrate at ambient RH for at least 2 days.

Material Characterization

Powder XRD

After ion exchange, all of the materials are vacuum filtered until dry and then equilibrated at ambient RH before XRD is taken. High-resolution synchrotron powder XRD patterns were collected on samples sealed in 1.0 mm (o.d.) glass capillaries (to prevent changes in RH). The samples were measured on beamline 28-ID-1 (λ = 0.1665 Å) at the National Synchrotron Light Source II at Brookhaven National Laboratory.

Additional XRD data were collected by using a Rigaku SmartLab diffractometer (CuKα). The hydrated samples were placed on a glass slide sample holder at ambient RH, while the dried samples were prepared in a glovebox and measured in a Rigaku air-free sample holder. All patterns were collected from 5° to 60° 2θ as a step size of 0.03° and 5° per minute.

Raman Spectroscopy

Raman spectroscopy was measured using a Horiba Instruments XplorRA PLUS Raman Spectrometer equipped with a 532 nm laser. The sample was mounted on a glass microscope slide. The signal was averaged over 200 acquisitions lasting 1 s each with a 50 μm slit and 500 μm hole. The laser power used was either 1 or 10% to prevent local heating and sample degradation.

TGA

TGA was performed by using a TA Instruments TGA 550. Powder samples (5–30 mg) were loaded into a tared high-temperature pan composed of an Inconel coated bail wire and platinum pan as a flat, evenly distributed layer and heated under a nitrogen flow (25 mL/min) at 5 °C/min from RT (19–25 °C) to 200 °C, at which it was held constant for 1 min. The instrument was calibrated using a nickel Curie temperature standard as per the manufacturer’s directions.

ICP-MS

ICP-MS was performed on an Agilent 8800. About 2 mg of each synthesized batch of material was digested in 2 mL of 70% nitic acid at 80 °C for 4 h. After the initial digestion, the solutions were diluted twice in 5% nitric acid to reach x2500 dilution. Five different concentrations of standard solutions were made from stock solutions of Cd, Mn, P, S, Li, Na, Mg, Ca, Zn, and Al to generate a calibration curve.

SEM and EDS

SEM was performed on select materials using a ZEISS 1550VP field emission SEM with an acceleration voltage of 10 kV at 5, 10, and 30 kX magnification. Before SEM was performed, the materials were sputtered with Pt for 5 s at 40 mA to avoid charging during the measurement. EDS data were collected using an Oxford X-MAX SSD system with an acceleration voltage of 10 or 20 kV.

EIS

EIS measurements were collected using a BioLogic VSP300 multichannel potentiostat with ultralow current probes. Typically, symmetric cells were assembled with Au-sputtered, in 0.25″ i.d. PTFE spring-loaded Swagelok cells. EIS was measured at different temperatures, controlled by a convection oven. The cell temperature was allowed to equilibrate for 30 min at each temperature. The temperature series were terminated at 70 °C. The EIS spectra were collected using a sinusoidal voltage amplitude of 50 mV in a frequency range of 3 MHz to 1 Hz and averaged over 10 measurements. At least three successive measurements were taken to ensure that the response was stable. Equivalent results were obtained without pelletization, by making a pellet in situ in the PTFE Swagelok cell by applying at least 3 kN of force with a vice. In this case, either polished Au foil or stainless steel plungers were used as electrodes. This method was primarily utilized for the Al-exchanged samples, which took a long time to re-equilibrate after pelletization, and it was also used for replicates of other samples to ensure consistency. Since for the vast majority of samples the high-frequency semicircle in the Nyquist plot could not be resolved even at RT, the x intercept of the Nyquist plot is taken as a “worst case” approximation for the total impedance for all samples, which we conservatively approximate as the electrolyte impedance. For the few samples where a high-frequency semicircle could be resolved, the data were fit to an equivalent circuit using ZFit in the EC-Lab software, e.g., Figure S17, to ensure that the capacitance of this feature corresponded with that expected for bulk ionic conductivity in solids.

Solid State MAS NMR Spectroscopy

Multinuclear MAS NMR experiments were performed using a Bruker Avance I-500 MHz spectrometer and using a Bruker 4 mm MAS NMR probe. A powder sample was packed into a zirconia (ZrO2) rotor at ambient conditions and spun at 10 kHz. 1H NMR (500.2 MHz) spectra were acquired after 4 μs-90 deg pulse. NMR signals of metal ions (quadrupole nuclei) were recorded after applying short tip angle rf pulses (1/12π for I = 3/2 nuclei (7Li and 23Na) or 1/18π pulse for 27Al) and strong 1H decoupling pulse. Chemical shifts were externally calibrated to TMS for 1H, and a 1 M aqueous solution of LiCl, NaCl, and Al(NO3) for 7Li, 23Na, and 27Al nuclei, respectively. For NMR measurements after dehydration, a 4 mm rotor containing packed powder sample was inserted into an 80 mm long-5 mm glass NMR tube, and the glass NMR tube was attached to a 1/4″-Cajon-VCR-T fitting and the side arm was connected to a vacuum manifold for high-temperature evacuation overnight. In this special setup, a glass rod with a sealing Kel-F rotor cap at the end was attached at the top of the 1/4″-Cajon-VCR-T. The rod was able to slide down, closing the 4 mm rotor with a tightly fit O-ring. The rotor underwent heating in a 10 mm-cylinder furnace with evacuation. The setup was filled with Ar gas before sealing, resulting in complete avoidance of exposure to air for the dried powder sample.

PFGNMR Spectroscopy

Li0.5Cd0.5PS3·2H2O sample was packed and sealed inside a 4 mm ZrO2 rotor, which was itself placed in an airtight 5 mm NMR tube. 1H and 7Li PFG NMR measurements were conducted on a 7.05 T (1H, 300 MHz) Bruker Avance III super wide-bore NMR spectrometer equipped with a Diff50 probe under static conditions. Diffusion measurements were performed from low (25 °C) to high temperature (60 °C) after a 30 min equilibration period at each temperature, and the temperature was regulated by a heater and using N2 gas flowing at a rate of 800 L/h. The sample temperature was calibrated using dry ethylene glycol solution. Self-diffusion coefficients were measured using a stimulated echo pulse sequence88 with variable magnetic field gradient pulses. The PFG NMR data were processed with TOPSPIN 4.3.0 and fitted using Python 3.12.4 using the various models discussed in detail in the Supporting Information.

Theoretical Methods

Structure Relaxations

We used DFT implemented in the Vienna Ab Initio Simulation Package (VASP 6.4.2)100−102 along with projector augmented wave (PAW)103,104 pseudopotentials to determine the lowest energy configurations of dehydrated and hydrated K0.5Mn0.75PS3 and Na0.5Mn0.75PS3. The stoichiometries were rounded to the stated values for the sake of computational simplicity. PAW potentials were used with valence configurations of 3s23p64s1 for K, 2s22p4 for O, 3s23p3 for P, 3s23p4 for S, 4s13d6 for Mn, and 1s1 for H to describe the valence electrons. Our calculations employed the Perdew–Burke–Ernzerhof105 generalized gradient approximation and included vdW corrected DFT-D3 (Becke-Johnson)106,107 along with empirical dispersion corrections. To ensure accuracy, we included all plane waves of energy up to 650 eV and set electronic minimization energy criteria to 10–6 eV. Ionic relaxation was stopped when the norms of all Hellmann–Feynman forces were less than 10–2/Å, and we utilized a 4 × 2 × 4 Γ-centered k-mesh for Brillouin zone integration within the unit cells so that a k-spacing of less than 0.2 Å–1 was sufficient for the required accuracy for electronic energy minimization.

The structures in Figure 3 were obtained through a heating and cooling approach using ab initio molecular dynamics (AIMD) and the Nosé and Hoover NVT ensemble in VASP 6.4.2.108−111 The structures were brought from 20 to 300 K, held at 300 K then cooled back to 20 K and were relaxed with DFT minimization from there. Initial estimate structures were optimized with the parameters described above. All AIMD simulations were run with 1 fs time steps, a fixed cell shape and volume, and the same parameters as described above except as noted. The heating and cooling were both performed over 2 ps with a 1 × 1 × 1 Γ-centered k-mesh. Between these steps, the simulation was held at 300 K over 5 ps with a 2 × 1 × 2 Γ-centered k-mesh.

Data Availability Statement

The data that support the findings of this study are openly available in CaltechDATA at https://dx.doi.org/10.22002/a58b4-m6e21.

Supporting Information Available

The Supporting Information is available free of charge at https://pubs.acs.org/doi/10.1021/jacs.4c06263.Differences in the degree of ion exchange that occurs, both for different ions within a given framework and between the frameworks; ligand exchange procedure, challenges, and results; XRD, EIS, Arrhenius relationships, TGA, DTG, Raman, SEM, EDS of MPS3-based intercalated compounds; XRD of select compounds at different RH; additional computational structural investigations of K0.5Mn0.75PS3·H2O; EIS of pristine MPS3 materials; XRD of dried MPS3-based ion-exchanged compounds; MAS NMR of select compounds in hydrated and dried states; XRD and Raman of compounds after ligand exchange; a comparison between electrochemical results of ligand-assisted ionic conduction and a confined liquid electrolyte; and electrochemical properties of each sample and the chemical shifts and fwhm from MAS NMR (PDF)

Supplementary Material

ja4c06263_si_001.pdf

The authors declare no competing financial interest.

Acknowledgments

This research was supported by the Arnold and Mabel Beckman Foundation through the Beckman Young Investigator Award. K.A.S. also acknowledges support from the Packard Fellowship for Science and Engineering, the Alfred P. Sloan Foundation, and the Camille and Henry Dreyfus Foundation. The ICP-MS data were collected at the Water and Environment Lab at Caltech. The authors thank Dr. Nathan Dalleska for assistance with ICP-MS sample preparation and data collection. The authors thank Michelle D. Qian and Gihan Kwon for their assistance in preparing samples for and measuring synchrotron XRD. S.S. thanks the UCLA-Caltech Medical Scientist Training Program (MSTP) through the National Institutes of Health (NIH) NIGMS training grant T32 GM008042 for support. The work of TD and WAG was supported by the Hong Kong Quantum AI Lab, AIR@InnoHK of the Hong Kong Government.
==== Refs
References

Grey C. P. ; Tarascon J. M. Sustainability and in Situ Monitoring in Battery Development. Nat. Mater. 2017, 16 , 45–56. 10.1038/nmat4777.
Olivetti E. A. ; Ceder G. ; Gaustad G. G. ; Fu X. Lithium-Ion Battery Supply Chain Considerations: Analysis of Potential Bottlenecks in Critical Metals. Joule 2017, 1 , 229–243. 10.1016/j.joule.2017.08.019.
Sun X. ; Hao H. ; Hartmann P. ; Liu Z. ; Zhao F. Supply Risks of Lithium-Ion Battery Materials: An Entire Supply Chain Estimation. Materials Today Energy 2019, 14 , 100347 10.1016/j.mtener.2019.100347.
Iton Z. W. B. ; See K. A. Multivalent Ion Conduction in Inorganic Solids. Chem. Mater. 2022, 34 , 881–898. 10.1021/acs.chemmater.1c04178.
Lundgren C. A. ; Murray R. W. Observations on the Composition of Prussian Blue Films and Their Electrochemistry. Inorg. Chem. 1988, 27 , 933–939. 10.1021/ic00278a036.
Itaya K. ; Uchida I. ; Neff V. D. Electrochemistry of Polynuclear Transition Metal Cyanides: Prussian Blue and Its Analogues. Acc. Chem. Res. 1986, 19 , 162–168. 10.1021/ar00126a001.
Lee H.-W. ; Wang R. Y. ; Pasta M. ; Woo Lee S. ; Liu N. ; Cui Y. Manganese Hexacyanomanganate Open Framework as a High-Capacity Positive Electrode Material for Sodium-Ion Batteries. Nat. Commun. 2014, 5 , 5280 10.1038/ncomms6280.25311066
Farrington G. Divalent Beta″-Aluminas: High Conductivity Solid Electrolytes for Divalent Cations. Solid State Ionics 1982, 7 , 267–281. 10.1016/0167-2738(82)90023-6.
Kummer J. β-Alumina Electrolytes. Prog. Solid State Chem. 1972, 7 , 141–175. 10.1016/0079-6786(72)90007-6.
Yao Y.-F. Y. ; Kummer J. Ion Exchange Properties of and Rates of Ionic Diffusion in Beta-Alumina. J. Inorg. Nucl. Chem. 1967, 29 , 2467–2475. 10.1016/0022-1902(67)80301-4.
Whittingham M. S. ; Huggins R. A. Measurement of Sodium Ion Transport in Beta Alumina Using Reversible Solid Electrodes. J. Chem. Phys. 1971, 54 , 414–416. 10.1063/1.1674623.
Goodenough J. ; Hong H.-P. ; Kafalas J. Fast Na+-Ion Transport in Skeleton Structures. Mater. Res. Bull. 1976, 11 , 203–220. 10.1016/0025-5408(76)90077-5.
Pal S. K. ; Saha R. ; Kumar G. V. ; Omar S. Designing High Ionic Conducting NASICON-type Na3Zr2Si2PO12 Solid-Electrolytes for Na-Ion Batteries. J. Phys. Chem. C 2020, 124 , 9161–9169. 10.1021/acs.jpcc.0c00543.
Guin M. ; Tietz F. ; Guillon O. New Promising NASICON Material as Solid Electrolyte for Sodium-Ion Batteries: Correlation between Composition, Crystal Structure and Ionic Conductivity of Na3+xSc2SixP3-xO12. Solid State Ionics 2016, 293 , 18–26. 10.1016/j.ssi.2016.06.005.
Aurbach D. ; Lu Z. ; Schechter A. ; Gofer Y. ; Gizbar H. ; Turgeman R. ; Cohen Y. ; Moshkovich M. ; Levi E. Prototype Systems for Rechargeable Magnesium Batteries. Nature 2000, 407 , 724–727. 10.1038/35037553.11048714
Sun X. ; Bonnick P. ; Nazar L. F. Layered TiS2 Positive Electrode for Mg Batteries. ACS Energy Lett. 2016, 1 , 297–301. 10.1021/acsenergylett.6b00145.
Sun X. ; Bonnick P. ; Duffort V. ; Liu M. ; Rong Z. ; Persson K. A. ; Ceder G. ; Nazar L. F. A High Capacity Thiospinel Cathode for Mg Batteries. Energy Environ. Sci. 2016, 9 , 2273–2277. 10.1039/C6EE00724D.
Canepa P. ; Bo S.-H. ; Sai Gautam G. ; Key B. ; Richards W. D. ; Shi T. ; Tian Y. ; Wang Y. ; Li J. ; Ceder G. High Magnesium Mobility in Ternary Spinel Chalcogenides. Nat. Commun. 2017, 8 , 1759 10.1038/s41467-017-01772-1.29170372
Glaser C. ; Wei Z. ; Indris S. ; Klement P. ; Chatterjee S. ; Ehrenberg H. ; Zhao-Karger Z. ; Rohnke M. ; Janek J. To Be or Not to Be – Is MgSc2Se4 a Mg-Ion Solid Electrolyte?. Adv. Energy Mater. 2023, 13 , 2301980 10.1002/aenm.202301980.
Martinolich A. J. ; Lee C.-W. ; Lu I.-T. ; Bevilacqua S. C. ; Preefer M. B. ; Bernardi M. ; Schleife A. ; See K. A. Solid-State Divalent Ion Conduction in ZnPS3. Chem. Mater. 2019, 31 , 3652–3661. 10.1021/acs.chemmater.9b00207.
Higashi S. ; Miwa K. ; Aoki M. ; Takechi K. A Novel Inorganic Solid State Ion Conductor for Rechargeable Mg Batteries. Chem. Commun. 2014, 50 , 1320–1322. 10.1039/C3CC47097K.
Yan Y. ; Grinderslev J. B. ; Jo̷rgensen M. ; Skov L. N. ; Skibsted J. ; Jensen T. R. Ammine Magnesium Borohydride Nanocomposites for All-Solid-State Magnesium Batteries. ACS Appl. Energy Mater. 2020, 3 , 9264–9270. 10.1021/acsaem.0c01599.
Yan Y. ; Dononelli W. ; Jo̷rgensen M. ; Grinderslev J. B. ; Lee Y.-S. ; Cho Y. W. ; Černý R. ; Hammer B. ; Jensen T. R. The Mechanism of Mg2+ Conduction in Ammine Magnesium Borohydride Promoted by a Neutral Molecule. Phys. Chem. Chem. Phys. 2020, 22 , 9204–9209. 10.1039/D0CP00158A.32232248
Le Ruyet R. ; Fleutot B. ; Berthelot R. ; Benabed Y. ; Hautier G. ; Filinchuk Y. ; Janot R. Mg3(BH4) 4(NH2)2 as Inorganic Solid Electrolyte with High Mg2+ Ionic Conductivity. ACS Appl. Energy Mater. 2020, 3 , 6093–6097. 10.1021/acsaem.0c00980.
Kisu K. ; Kim S. ; Inukai M. ; Oguchi H. ; Takagi S. ; Orimo S.-I. Magnesium Borohydride Ammonia Borane as a Magnesium Ionic Conductor. ACS Appl. Energy Mater. 2020, 3 , 3174–3179. 10.1021/acsaem.0c00113.
Roedern E. ; Kühnel R.-S. ; Remhof A. ; Battaglia C. Magnesium Ethylenediamine Borohydride as Solid-State Electrolyte for Magnesium Batteries. Sci. Rep. 2017, 7 , 46189 10.1038/srep46189.28387305
Whittingham S. Transport Properties of the Mineral Vermiculite. Solid State Ionics 1989, 32–33 , 344–349. 10.1016/0167-2738(89)90239-7.
Whittingham M. Sodium Ion Conduction in Single Crystal Vermiculite. Solid State Ionics 1987, 25 , 295–300. 10.1016/0167-2738(87)90193-7.
Maraqah H. ; Li J. ; Whittingham M. S. Ion Transport in Single Crystals of the Clay-Like Aluminosilicate Vermiculite. MRS Proc. 1990, 210 , 351 10.1557/PROC-210-351.
Slade R. Conduction and Diffusion in Exchanged Montmorillonite Clays. Solid State Ionics 1987, 24 , 289–295. 10.1016/0167-2738(87)90135-4.
Ruiz-Hitzky E. ; Casal B. Crown Ether Intercalations with Phyllosilicates. Nature 1978, 276 , 596–597. 10.1038/276596a0.
Nam K. W. ; et al. The High Performance of Crystal Water Containing Manganese Birnessite Cathodes for Magnesium Batteries. Nano Lett. 2015, 15 , 4071–4079. 10.1021/acs.nanolett.5b01109.25985060
Novák P. ; Desilvestro J. Electrochemical Insertion of Magnesium in Metal Oxides and Sulfides from Aprotic Electrolytes. J. Electrochem. Soc. 1993, 140 , 140–144. 10.1149/1.2056075.
Song J. ; Noked M. ; Gillette E. ; Duay J. ; Rubloff G. ; Lee S. B. Activation of a MnO2 Cathode by Water-Stimulated Mg2+ Insertion for a Magnesium Ion Battery. Phys. Chem. Chem. Phys. 2015, 17 , 5256–5264. 10.1039/C4CP05591H.25608277
Kundu D. ; Adams B. D. ; Duffort V. ; Vajargah S. H. ; Nazar L. F. A High-Capacity and Long-Life Aqueous Rechargeable Zinc Battery Using a Metal Oxide Intercalation Cathode. Nat. Energy 2016, 1 , 16119 10.1038/nenergy.2016.119.
Joos M. ; Schneider C. ; Münchinger A. ; Moudrakovski I. ; Usiskin R. ; Maier J. ; Lotsch B. V. Impact of Hydration on Ion Transport in Li2Sn2S5·xH2O. J. Mater. Chem. A 2021, 9 , 16532–16544. 10.1039/D1TA04736A.
Iton Z. W. B. ; Lee B. C. ; Jiang A. Y. ; Kim S. S. ; Brady M. J. ; Shaker S. ; See K. A. Water Vapor Induced Superionic Conductivity in ZnPS 3. J. Am. Chem. Soc. 2023, 145 , 13312–13325. 10.1021/jacs.3c03368.37294168
Yoshida Y. ; Yamada T. ; Jing Y. ; Toyao T. ; Shimizu K.-I. ; Sadakiyo M. Super Mg2+ Conductivity around 10–3 Scm–1 Observed in a Porous Metal–Organic Framework. J. Am. Chem. Soc. 2022, 144 , 8669–8675. 10.1021/jacs.2c01612.35507008
Yoshida Y. ; Kato K. ; Sadakiyo M. Vapor-Induced Superionic Conduction of Magnesium Ions in a Metal–Organic Framework. J. Phys. Chem. C 2021, 125 , 21124–21130. 10.1021/acs.jpcc.1c05250.
Park S. ; Kristanto I. ; Jung G. Y. ; Ahn D. B. ; Jeong K. ; Kwak S. K. ; Lee S.-Y. A Single-Ion Conducting Covalent Organic Framework for Aqueous Rechargeable Zn-ion Batteries. Chem. Sci. 2020, 11 , 11692–11698. 10.1039/D0SC02785E.34123199
Miner E. M. ; Park S. S. ; Dincă M. High Li+ and Mg2+ Conductivity in a Cu-Azolate Metal–Organic Framework. J. Am. Chem. Soc. 2019, 141 , 4422–4427. 10.1021/jacs.8b13418.30773017
Iliescu A. ; Andrews J. L. ; Oppenheim J. J. ; Dincă M. A Solid Zn-Ion Conductor from an All-Zinc Metal–Organic Framework Replete with Mobile Zn2+ Cations. J. Am. Chem. Soc. 2023, 145 , 25962–25965. 10.1021/jacs.3c10326.38010994
Clement R. A Novel Route to Intercalation into Layered MnPS3. J. Chem. Soc., Chem. Commun. 1980, 647 10.1039/c39800000647.
Whittingham M. S. The Role of Ternary Phases in Cathode Reactions. J. Electrochem. Soc. 1976, 123 , 315–320. 10.1149/1.2132817.
Clement R. ; Garnier O. ; Jegoudez J. Coordination Chemistry of the Lamellar MPS3 Materials: Metal-Ligand Cleavage as the Source of an Unusual “Cation-Transfer” Intercalation Process. Inorg. Chem. 1986, 25 , 1404–1409. 10.1021/ic00229a022.
Clément R. ; Lagadic I. ; Léaustic A. ; Audière J. P. ; Lomas L. In Chemical Physics of Intercalation II; Bernier P. ; Fischer J. E. ; Roth S. ; Solin S. A. , Eds.; Springer US: Boston, MA, 1993; Vol. 305 ; pp 315–324.
Fan Y. A New Family of Fast Ion Conductor-Montmorillonites. Solid State Ionics 1997, 93 , 347–354. 10.1016/S0167-2738(96)00446-8.
Suzuki M. ; Wada N. ; Hines D. ; Whittingham M. S. Hydration States and Phase Transitions in Vermiculite Intercalation Compounds. Phys. Rev. B 1987, 36 , 2844–2851. 10.1103/PhysRevB.36.2844.
Ouvrard G. ; Brec R. ; Rouxel J. Structural Determination of Some MPS3 Layered Phases (M = Mn, Fe, Co, Ni and Cd). Mater. Res. Bull. 1985, 20 , 1181–1189. 10.1016/0025-5408(85)90092-3.
Shannon R. D. Revised Effective Ionic Radii and Systematic Studies of Interatomic Distances in Halides and Chalcogenides. Acta Cryst. A 1976, 32 , 751–767. 10.1107/S0567739476001551.
Coradin T. ; Clément R. ; Lacroix P. G. ; Nakatani K. From Intercalation to Aggregation: Nonlinear Optical Properties of Stilbazolium Chromophores-MPS3 Layered Hybrid Materials. Chem. Mater. 1996, 8 , 2153–2158. 10.1021/cm960060x.
Lagadic I. ; Lacroix P. G. ; Clément R. Layered MPS 3 (M = Mn, Cd) Thin Films as Host Matrixes for Nonlinear Optical Material Processing. Chem. Mater. 1997, 9 , 2004–2012. 10.1021/cm970155e.
Clement R. ; Leaustic A. ; Marney K. ; Francis A. Synthesis and Luminescence Properties of CdPS3 Intercalated with Rare Earth Cations. J. Phys. Chem. Solids 1994, 55 , 9–16. 10.1016/0022-3697(94)90178-3.
Coradin T. ; Coupé A. ; Livage J. Intercalation of Biomolecules in the MnPS3 Layered Phase. J. Mater. Chem. 2003, 13 , 705–707. 10.1039/b210514d.
Lagadic I. ; Léaustic A. ; Clément R. Intercalation of Polyethers into the MPS3 (M = Mn, Cd) Host Lattice. J. Chem. Soc., Chem. Commun. 1992, 0 , 1396–1397. 10.1039/C39920001396.
Oriakhi C. O. ; Nafshun R. L. ; Lerner M. M. Preparation of Nanocomposites of Linear Poly(Ethylenimine) with Layered Hosts. Mater. Res. Bull. 1996, 31 , 1513–1520. 10.1016/S0025-5408(96)00142-0.
Clement R. ; Lomas L. ; Audiere J. P. Intercalation Chemistry of Layered Iron Trithiohypophosphate (FePS3). An Approach toward Insulating Magnets below 90 K. Chem. Mater. 1990, 2 , 641–643. 10.1021/cm00012a009.
Jeevanandam P. ; Vasudevan S. Preparation and Characterization of Cd0.75PS3A0.5(H2O)y [A = Na, K and Cs]. Solid State Ionics 1997, 104 , 45–55. 10.1016/S0167-2738(97)00411-6.
Jeevanandam P. ; Vasudevan S. Conductivity and Dielectric Response in the Ion-Exchange Intercalated Mono- and Double-Layer Hydrates Cd0.75PS3Na 0.5(H2O) y, y = 1, 2. J. Phys. Chem. B 1998, 102 , 3082–3089. 10.1021/jp972945h.
Arun N. ; Jeevanandam P. ; Vasudevan S. ; Ramanathan K. V. Motion of Interlamellar Hydrated Sodium Ions in Layered Cd0.75PS3Na0.5 (H2O)2. J. Chem. Phys. 1999, 111 , 1231–1239. 10.1063/1.479308.
Arun N. ; Vasudevan S. ; Ramanathan K. V. Orientation and Motion of Interlamellar Water: An Infrared and NMR Investigation of Water in the Galleries of Layered Cd0.75PS3K0.5(H2O) y. J. Am. Chem. Soc. 2000, 122 , 6028–6038. 10.1021/ja991357+.
Ruiz-León D. ; Manríquez V. ; Kasaneva J. ; Avila R. Insertion of Trivalent Cations in the Layered MPS3 (Mn, Cd) Materials. Mater. Res. Bull. 2002, 37 , 981–989. 10.1016/S0025-5408(02)00719-5.
Qian X. ; Chen L. ; Yin L. ; Liu Z. ; Pei S. ; Li F. ; Hou G. ; Chen S. ; Song L. ; Thebo K. H. ; Cheng H.-M. ; Ren W. CdPS3 Nanosheets-Based Membrane with High Proton Conductivity Enabled by Cd Vacancies. Science 2020, 370 , 596–600. 10.1126/science.abb9704.33122384
Yu X. ; Ren W. 2D CdPS3-Based Versatile Superionic Conductors. Nat. Commun. 2023, 14 , 3998 10.1038/s41467-023-39725-6.37414802
Agmon N. The Grotthuss Mechanism. Chem. Phys. Lett. 1995, 244 , 456–462. 10.1016/0009-2614(95)00905-J.
Hou T. ; Xu W. ; Pei X. ; Jiang L. ; Yaghi O. M. ; Persson K. A. Ionic Conduction Mechanism and Design of Metal–Organic Framework Based Quasi-Solid-State Electrolytes. J. Am. Chem. Soc. 2022, 144 , 13446–13450. 10.1021/jacs.2c03710.35700972
García N. Conductivity in Na+- and Li+-Montmorillonite as a Function of Equilibration Humidity. Solid State Ionics 1996, 92 , 139–143. 10.1016/S0167-2738(96)00447-X.
Fan Y. Cation Diffusion and Conduction in Solid Electrolytes Li, Na-montmorillonites. Solid State Ionics 1988, 28–30 , 1596–1601. 10.1016/0167-2738(88)90426-2.
Barj M. ; Lucazeau G. Raman Spectra of Lamellar CdPS3 Intercalated with Alkali Ions. Solid State Ionics 1983, 9–10 , 475–479. 10.1016/0167-2738(83)90279-5.
Schöllhorn R. ; Meyer H. Cathodic Reduction of Layered Transition Metal Chalcogenides. Mater. Res. Bull. 1974, 9 , 1237–1245. 10.1016/0025-5408(74)90042-7.
Lerf A. ; Schoellhorn R. Solvation Reactions of Layered Ternary Sulfides AxTiS2, AxNbS2, and AxTaS2. Inorg. Chem. 1977, 16 , 2950–2956. 10.1021/ic50177a057.
Whittingham M. The Hydrated Intercalation Complexes of the Layered Disulfides. Mater. Res. Bull. 1974, 9 , 1681–1689. 10.1016/0025-5408(74)90162-7.
Ghidiu M. ; Halim J. ; Kota S. ; Bish D. ; Gogotsi Y. ; Barsoum M. W. Ion-Exchange and Cation Solvation Reactions in Ti3C2 MXene. Chem. Mater. 2016, 28 , 3507–3514. 10.1021/acs.chemmater.6b01275.
Prouzet E. ; Ouvrard G. ; Brec R. Structure Determination of ZnPS3. Mater. Res. Bull. 1986, 21 , 195–200. 10.1016/0025-5408(86)90206-0.
Famprikis T. ; Canepa P. ; Dawson J. A. ; Islam M. S. ; Masquelier C. Fundamentals of Inorganic Solid-State Electrolytes for Batteries. Nat. Mater. 2019, 18 , 1278–1291. 10.1038/s41563-019-0431-3.31427742
Frenkel M. Surface Acidity of Montmorillonites. Clays and Clay Minerals 1974, 22 , 435–441. 10.1346/CCMN.1974.0220510.
Hatz A.-K. ; Moudrakovski I. ; Bette S. ; Terban M. W. ; Etter M. ; Joos M. ; Vargas-Barbosa N. M. ; Dinnebier R. E. ; Lotsch B. V. Fast Water-Assisted Lithium Ion Conduction in Restacked Lithium Tin Sulfide Nanosheets. Chem. Mater. 2021, 33 , 7337–7349. 10.1021/acs.chemmater.1c01755.
Ye G. ; Janzen N. ; Goward G. R. Solid-State NMR Study of Two Classic Proton Conducting Polymers: Nafion and Sulfonated Poly(Ether Ether Ketone)s. Macromolecules 2006, 39 , 3283–3290. 10.1021/ma0523825.
Ratcliffe C. ; Ripmeester J. ; Tse J. NMR Chemical Shifts of Dilute 1H in Inorganic Solids. Chem. Phys. Lett. 1985, 120 , 427–432. 10.1016/0009-2614(85)85634-7.
Volkov V. I. ; Chernyak A. V. ; Slesarenko N. A. ; Avilova I. A. Ion and Molecular Transport in Solid Electrolytes Studied by NMR. IJMS 2022, 23 , 5011 10.3390/ijms23095011.35563404
Kwon B. J. ; et al. High Voltage Mg-Ion Battery Cathode via a Solid Solution Cr–Mn Spinel Oxide. Chem. Mater. 2020, 32 , 6577–6587. 10.1021/acs.chemmater.0c01988.
Gun’ko V. M. ; Turov V. V. Structure of Hydrogen Bonds and 1 H NMR Spectra of Water at the Interface of Oxides. Langmuir 1999, 15 , 6405–6415. 10.1021/la9809372.
Hayashi S. ; Mizuno M. Proton Dynamics in Cs2(HSO4)(H2PO4) Studied by 1H NMR. Solid State Ionics 2005, 176 , 745–754. 10.1016/j.ssi.2004.10.008.
Sanz J. ; Sobrados I. ; Robert J.-L. Influence of Hydration on 23Na, 27Al, and 29Si MAS-NMR Spectra of Sodium Saponites and Sodium Micas. Am. Mineral. 2015, 100 , 1076–1083. 10.2138/am-2015-4832.
Ohkubo T. ; Saito K. ; Kanehashi K. ; Ikeda Y. A Study on Hydration Behaviors of Interlayer Cations in Montmorillonite by Solid State NMR. Sci. Technol. Adv. Mater. 2004, 5 , 693–696. 10.1016/j.stam.2004.02.016.
Malicki N. ; Beccat P. ; Bourges P. ; Fernandez C. ; Quoineaud A.-A. ; Simon L. J. ; Thibault-Starzyk F. Studies in Surface Science and Catalysis. Elsevier 2007, 170 , 762–770. 10.1016/S0167-2991(07)80918-9.
Haouas M. ; Taulelle F. ; Martineau C. Recent Advances in Application of 27Al NMR Spectroscopy to Materials Science. Prog. Nucl. Magn. Reson. Spectrosc. 2016, 94–95 , 11–36. 10.1016/j.pnmrs.2016.01.003.
Tanner J. E. Use of the Stimulated Echo in NMR Diffusion Studies. J. Chem. Phys. 1970, 52 , 2523–2526. 10.1063/1.1673336.
Stejskal E. O. ; Tanner J. E. Spin Diffusion Measurements: Spin Echoes in the Presence of a Time-Dependent Field Gradient. J. Chem. Phys. 1965, 42 , 288–292. 10.1063/1.1695690.
Stoll M. ; Kaess U. ; Majer G. ; Barnes R. Nuclear Magnetic Resonance Studies of Hydrogen Diffusion in the Layer-Structured System ZrClHx. J. Alloys Compd. 1997, 253–254 , 435–440. 10.1016/S0925-8388(96)03080-0.
Kimmerle F. ; Majer G. ; Kaess U. ; Maeland A. ; Conradi M. ; McDowell A. NMR Studies of Hydrogen Diffusion in ZrBe2H1.4. J. Alloys Compd. 1998, 264 , 63–70. 10.1016/S0925-8388(97)00273-9.
Manríquez V. ; Galdámez A. ; Villanueva A. ; Aranda P. ; Galván J. C. ; Ruiz-Hitzky E. Insertion of In(III) and Ga(III) into MPS3 (M = Mn, Cd) Layered Materials. Mater. Res. Bull. 1999, 34 , 673–683. 10.1016/S0025-5408(99)00059-8.
Scholz T. ; Schneider C. ; Eger R. ; Duppel V. ; Moudrakovski I. ; Schulz A. ; Nuss J. ; Lotsch B. V. Phase Formation through Synthetic Control: Polymorphism in the Sodium-Ion Solid Electrolyte Na 4 P 2 S 6. J. Mater. Chem. A 2021, 9 , 8692–8703. 10.1039/D0TA11008F.
Haber S. ; Leskes M. What Can We Learn from Solid State NMR on the Electrode–Electrolyte Interface?. Adv. Mater. 2018, 30 , 1706496 10.1002/adma.201706496.
Ramsey N. F. Magnetic Shielding of Nuclei in Molecules. Phys. Rev. 1950, 78 , 699–703. 10.1103/PhysRev.78.699.
Komoroski R. A. Applications of 7Li NMR in Biomedicine. Magn. Reson. Imaging 2000, 18 , 103–116. 10.1016/S0730-725X(99)00116-2.10722969
Beguin F. ; Setton R. ; Beguin F. ; Setton R. ; Hamwi A. ; Touzain P. The Reversible Intercalation of Tetrahydrofuran in Some Graphite-Alkali Metal Lamellar Compounds. Materials Science and Engineering 1979, 40 , 167–173. 10.1016/0025-5416(79)90186-1.
Miner E. M. ; Dincă M. Metal- and Covalent-Organic Frameworks as Solid-State Electrolytes for Metal-Ion Batteries. Philos. Trans. R. Soc. A 2019, 377 , 20180225 10.1098/rsta.2018.0225.
Kharod R. A. ; Andrews J. L. ; Dincă M. Teaching Metal-Organic Frameworks to Conduct: Ion and Electron Transport in Metal-Organic Frameworks. Annu. Rev. Mater. Res. 2022, 52 , 103–128. 10.1146/annurev-matsci-080619-012811.
Kresse G. ; Hafner J. Ab Initio Molecular Dynamics for Liquid Metals. Phys. Rev. B 1993, 47 , 558–561. 10.1103/PhysRevB.47.558.
Kresse G. ; Furthmüller J. Efficiency of Ab-Initio Total Energy Calculations for Metals and Semiconductors Using a Plane-Wave Basis Set. Comput. Mater. Sci. 1996, 6 , 15–50. 10.1016/0927-0256(96)00008-0.
Kresse G. ; Furthmüller J. Efficient Iterative Schemes for Ab Initio Total-Energy Calculations Using a Plane-Wave Basis Set. Phys. Rev. B 1996, 54 , 11169–11186. 10.1103/PhysRevB.54.11169.
Kresse G. ; Joubert D. From Ultrasoft Pseudopotentials to the Projector Augmented-Wave Method. Phys. Rev. B 1999, 59 , 1758–1775. 10.1103/PhysRevB.59.1758.
Blöchl P. E. Projector Augmented-Wave Method. Phys. Rev. B 1994, 50 , 17953–17979. 10.1103/PhysRevB.50.17953.
Perdew J. P. ; Burke K. ; Ernzerhof M. Generalized Gradient Approximation Made Simple. Phys. Rev. Lett. 1996, 77 , 3865–3868. 10.1103/PhysRevLett.77.3865.10062328
Grimme S. ; Antony J. ; Ehrlich S. ; Krieg H. A Consistent and Accurate Ab Initio Parametrization of Density Functional Dispersion Correction (DFT-D) for the 94 Elements H-Pu. J. Chem. Phys. 2010, 132 , 154104 10.1063/1.3382344.20423165
Grimme S. ; Ehrlich S. ; Goerigk L. Effect of the Damping Function in Dispersion Corrected Density Functional Theory. J. Comput. Chem. 2011, 32 , 1456–1465. 10.1002/jcc.21759.21370243
Nosé S. A Unified Formulation of the Constant Temperature Molecular Dynamics Methods. J. Chem. Phys. 1984, 81 , 511–519. 10.1063/1.447334.
Nosé S. Constant Temperature Molecular Dynamics Methods. Prog. Theor. Phys. Suppl. 1991, 103 , 1–46. 10.1143/PTPS.103.1.
Hoover W. G. Canonical Dynamics: Equilibrium Phase-Space Distributions. Phys. Rev. A 1985, 31 , 1695–1697. 10.1103/PhysRevA.31.1695.
Frenkel D. ; Smit B. Understanding Molecular Simulation; Elsevier, 2002; pp 139–163.
