Crystalline Li-Ta-Oxychlorides with Lithium Superionic Conduction

Abstract

Nowadays, some amorphous and microcrystalline solid-state electrolytes (SSEs) with dual anions have attained high ionic conductivity and good compatibility with electrodes in all-solid-state lithium-ion batteries (ASSLIBs). In this work, crystalline SSEs of series A (Li_1+xTaO_1+xCl_4−x, −0.70 ≤ x ≤ 0.50) and B (LiTaO_2+yCl_2−2y, −1.22 ≤ y ≤ 0), having great application potential well over ambient temperatures, were prepared at 260–460 °C for 2–10 h using Li₂O, TaCl₅, and LiTaO₃ as the raw materials. The three-phase coexisting samples attained high σ values ranging from 5.20 to 7.35 mS cm⁻¹, which are among the reported high values of amorphous co-essential SSEs and other alloplasmatic crystalline ones. It is attributed to the synergistic effect of the polyanion trans-[O₂Cl₄] and cis-[O₄Cl₂] octahedra framework.

1. Introduction

Solid-state electrolytes (SSEs) in all-solid-state lithium-ion batteries (ASSLIBs) have been actively investigated recently. The systems were focused on oxide-, sulfide-, halide-, oxysulfide-, and oxyhalide-based compounds with amorphous, nanocrystalline, microcrystalline, or polycrystalline (ceramic) phases [1,2,3]. The optimized nanocrystalline LiTa(Nb)OCl₄ [4] reached a high level of Li-ion conductivity (σ > 10 mS cm⁻¹) with a specially low activation energy (E_a < 0.25 eV) at room temperature (RT; = 25 °C), being comparable to or surpassing that of organic liquid electrolytes (~10 mS cm⁻¹).

Li-Ta-oxychlorides (LTOCs)/Li-Nb-oxychlorides (LNOCs), constructed by introducing a mixed anion strategy in halide matrices, hold the following advantages [5]: (1) coupled weak/strong Coulombic forces between Cl⁻-Li⁺ and O²⁻-Li⁺, and wide/narrow Li⁺-transport channel formed by large/small anions ($r_{{\mathrm{C}\mathrm{l}}^{-}}$ = 0.181 nm, $r_{{\mathrm{O}}^{2-}}$ = 0.140 nm; r—effective ionic radius [5]); (2) better air stability and recoverability than sulfide-, halide-, and oxysulfide-based SSEs; and (3) coupled weak/strong bond strength with polyanions, softer/harder lattice structures, and high/low Cl⁻/O²⁻ polarizability to guarantee the SSEs’ good mechanical deformability for easily assembled ASSLIBs [1,2,3,4,6].

Amorphous Li-Ta(Nb)-halides and Li-Ta(Nb)-oxyhalides emerged as promising candidates for SSEs in ASSLIBs very recently within four categories: (i) ball-milled resultants from oxide-halide mixtures: xLi₂O-TaCl₅ (x = 1−2; σ = 6.6 mS cm⁻¹ @x = 1.6) [7] and 1.5Li₂O-0.9TaCl₅-0.1TaF₅ (2.3 mS cm⁻¹) [8]; (ii) halides: a-LiTaCl₆ (7.16 mS cm⁻¹ @30 °C [9]; 10.95 mS cm⁻¹ [10]; 11.0 mS cm⁻¹ (calculated) [11]), a-LiNbCl₆ (13.5 mS cm⁻¹ (calculated) [11]), a-LiTa_0.5Nb_0.5Cl₆ (15.7 mS cm⁻¹ (calculated) [11]), and a-LiTaCl₅F (4.37 mS cm⁻¹ [10]); (iii) nitrohalides: a-Li_3xTaCl₅N_x (1 ≤ 3x ≤ 1.4; 7.34 mS cm⁻¹ @3x = 1.25 & 30 °C) [12]; and (iv) oxyhalides: a-LiTaCl₅O_0.5 (7.75 mS cm⁻¹ [10]) and a-LiTaOCl₄ (5.03 mS cm⁻¹ @30 °C [13]).

The above amorphous/glassy SSEs were all synthesized via high-energy ball-milling (HEBM) or a mechanochemical process at 100–500 rpm/min for 2–200 h, while some were with post-annealed at mild temperatures (c.f. 120 °C × 2 h [10]).

Nanocrystalline LTOCs/LNOCs were a huge breakthrough by Tanaka et al. in 2023 [4]. They were synthesized from LiOH and TaCl₅/NbCl₅ by HEBM at 500 rpm/min for 40 h and post-annealed at 300/100 °C, respectively. They exhibited extremely high σ values at RT (12.4 and 10.4 mS cm⁻¹, respectively).

The structure of nanocrystalline LiNbOCl₄ was resolved and refined (based on LiVOF₄) to a Li-ordered orthorhombic system, i.e., o-LiNbOCl₄ (Cmc2₁, No. 36), which consist of 1D chains of apex-O-sharing [NbO₂Cl₄]-octahedra with (LiCl₄)-tetrahedra residing between those chains [4].

In 2024, Adams [14] argued the o-LiNbOCl₄ structure and revised it to t-LiNbOCl₄ (I4/m, No. 87) using ab initio molecular dynamics (AIMD) simulations. The structure consists of parallel (NbO₂/₂Cl₄)_∞ polyanion chains, and the pronounced Li-disorder leads to a more uniform packing within it.

Subsequently, Singh et al. [15] modified the above o-/t-LiNbOCl₄ structures to m-LiNbOCl₄ (Cc, No. 9) (4.4–10.7 mS cm⁻¹) by combining experiments with X-ray/neutron diffraction, pair distribution function (PDF) analysis, and density functional theory (DFT)/AIMD simulations. It crystallizes at ~120 °C and exhibits nanocrystalline behavior with a short coherence length (~10 nm). It features isolated 1D chains of [NbO₂Cl₄]-octahedra bound together by a relatively low concentration of highly delocalized Li-ions, as well as weak van der Waals interactions [15].

The processing conditions, such as HEBM rotation speed/time, annealing temperature/time etc., are closely correlated with the crystallinity, crystal structure, and ionic conductivity of SSEs. There are two controversial results for this mutual relationship.

(1) Case 1: Higher crystallinity leads to higher ionic conductivity.

Increasing the annealing temperature/time would increase the crystallinity and σ values of SSEs [4,15] (Equations (1) and (2)). Increasing the HEBM time would decrease the crystallinity and σ values of SSEs [15] (Equation (3)).

$$a-{\mathrm{LiTaOCl}}_4 (7.8 \mathrm{mS} {\mathrm{cm}}^{-1}) \stackrel{ 300 \mathrm{℃} , \mathrm{i}\mathrm{n} \mathrm{A}\mathrm{r}}{\to } o-{\mathrm{LiTaOCl}}_4 (12.4 \mathrm{mS} {\mathrm{cm}}^{-1})$$

(1)

$$a-{\mathrm{LiNbOCl}}_4 (9.1 \mathrm{mS} {\mathrm{cm}}^{-1}) \stackrel{ 100-120 \mathrm{℃} \times 0.5-12 \mathrm{h}, \mathrm{i}\mathrm{n} \mathrm{A}\mathrm{r}}{\to } o-{\mathrm{LiNbOCl}}_4 (10.4 \mathrm{mS} {\mathrm{cm}}^{-1})$$

(2)

$$a-{\mathrm{LiNbOCl}}_4 \mathrm{as}-\mathrm{milled}-8 \mathrm{h} (7 \mathrm{mS} {\mathrm{cm}}^{-1}) \to a-{\mathrm{LiNbOCl}}_4 \mathrm{as}-\mathrm{milled}-24 \mathrm{h} (4.4 \mathrm{mS} {\mathrm{cm}}^{-1})$$

(3)

Otherwise, the σ values of Li₃YBr₆ (Equation (4)) [16] and Li₃InCl₆ (Equation (5)) [17] were also consistent with this case.

$$m-{\mathrm{Li}}_3{\mathrm{YBr}}_6 (0.39 \mathrm{mS} {\mathrm{cm}}^{-1} \text{@}30 \text{°}\mathrm{C}) \stackrel{ 250/300/400/500/550 °\mathrm{C} \times 5 \mathrm{h}, \mathrm{i}\mathrm{n} \mathrm{A}\mathrm{r}}{\to } m-{\mathrm{Li}}_3{\mathrm{YBr}}_6 (1.18/1.26/1.42/3.31/1.76 \mathrm{mS} {\mathrm{cm}}^{-1} \text{@}30 \text{°}\mathrm{C})$$

(4)

$$m-{\mathrm{Li}}_3{\mathrm{InCl}}_6 (0.837 \mathrm{mS} {\mathrm{cm}}^{-1}) \stackrel{ 260 °\mathrm{C} \times 5 \mathrm{h}, \mathrm{i}\mathrm{n} \mathrm{A}\mathrm{r}}{\to } m-{\mathrm{Li}}_3{\mathrm{InCl}}_6 (1.49 \mathrm{mS} {\mathrm{cm}}^{-1})$$

(5)

This case corresponds to nano/microcrystalline structures with high temperature stability and slower crystallization. The structurally ordered (annealed) phase turns out to be superior in terms of ionic conductivity (σ_total) to the disordered (as-milled) one because ${\sigma }_{{\mathrm{L}\mathrm{i}}^{+}, \mathrm{a}\mathrm{s}-\mathrm{m}\mathrm{i}\mathrm{l}\mathrm{l}\mathrm{e}\mathrm{d}}$ ≈ ${\sigma }_{{\mathrm{L}\mathrm{i}}^{+}, \mathrm{a}\mathrm{n}\mathrm{n}\mathrm{e}\mathrm{a}\mathrm{l}\mathrm{e}\mathrm{d}}$ and ${\sigma }_{e^{-}, \mathrm{a}\mathrm{s}-\mathrm{m}\mathrm{i}\mathrm{l}\mathrm{l}\mathrm{e}\mathrm{d}}$ > ${\sigma }_{e^{-}, \mathrm{a}\mathrm{n}\mathrm{n}\mathrm{e}\mathrm{a}\mathrm{l}\mathrm{e}\mathrm{d}}$ [16,17,18].

(2) Case 2: Higher crystallinity leads to lower ionic conductivity.

In contrast with Case 1, increasing the annealing temperature/time would increase the crystallinity (in some cases, the crystal structure changes, such as Equation (6)) [19], but the σ values of SSEs are reduced (Equations (6) and (7); $P\overline{3}m1$, No. 164) [19,20]. Increasing the HEBM time (intense milling) would decrease the crystallinity and increase the σ values of SSEs. This contribution originates from factors (metastable nonperiodic features) beyond ideal and periodic atomic arrangements, such as defects, surface structure, particle size, strain, amorphous phase, etc. [19].

$$\mathit{trig}-{\mathrm{Li}}_2{\mathrm{ZrCl}}_6 (0.808 \mathrm{mS} {\mathrm{cm}}^{-1}) \stackrel{ 350 °\mathrm{C} \times 5 \mathrm{h}, \mathrm{i}\mathrm{n}\mathrm{A}\mathrm{r}}{\to } m-{\mathrm{Li}}_2{\mathrm{ZrCl}}_6 (0.00581 \mathrm{mS} {\mathrm{cm}}^{-1})$$

(6)

$$\begin{array}{c}\mathit{trig}-{\mathrm{Li}}_{2.6}{\mathrm{Er}}_{0.6}{\mathrm{Zr}}_{0.4}{\mathrm{Cl}}_6 (1.13 \mathrm{mS} {\mathrm{cm}}^{-1} \text{@}27 \text{°}\mathrm{C}) \stackrel{ 200/300/400/450 °\mathrm{C} \times 5 \mathrm{h}, \mathrm{i}\mathrm{n} \mathrm{A}\mathrm{r}}{\to } \mathit{trig}-{\mathrm{Li}}_{2.6}{\mathrm{Er}}_{0.6}{\mathrm{Zr}}_{0.4}{\mathrm{Cl}}_6 \\ (0.972/0.383/0.125/0.0881 \mathrm{mS} {\mathrm{cm}}^{-1} \text{@}27 \text{°}\mathrm{C})\end{array}$$

(7)

The structurally ordered (annealed) phase turns out to be inferior in terms of ionic conductivity (σ_total) to the disordered (as-milled) one because ${\sigma }_{{\mathrm{L}\mathrm{i}}^{+}, \mathrm{a}\mathrm{s}-\mathrm{m}\mathrm{i}\mathrm{l}\mathrm{l}\mathrm{e}\mathrm{d}}$ > ${\sigma }_{{\mathrm{L}\mathrm{i}}^{+}, \mathrm{a}\mathrm{n}\mathrm{n}\mathrm{e}\mathrm{a}\mathrm{l}\mathrm{e}\mathrm{d}}$ and ${\sigma }_{e^{-}, \mathrm{a}\mathrm{s}-\mathrm{m}\mathrm{i}\mathrm{l}\mathrm{l}\mathrm{e}\mathrm{d}}$ ≈ ${\sigma }_{e^{-}, \mathrm{a}\mathrm{n}\mathrm{n}\mathrm{e}\mathrm{a}\mathrm{l}\mathrm{e}\mathrm{d}}$ << ${\sigma }_{{\mathrm{L}\mathrm{i}}^{+}}$.

Essentially, the above cases depend on the thermodynamic and kinetic stability ranges of SSEs. Even though o-LiTaOCl₄/o-LiNbOCl₄ reached, respectively, the highest σ values of 12.4/10.4 mS cm⁻¹ in oxyhalide-based SSEs, there is a desperate need for deeper research on the corresponding maximal achievable crystalline phases. To increase the crystallinity, a thermodynamic stability range at mild temperatures (100–120 °C; < 150 °C for 0.5–12 h) was found for metastable a-LiNbOCl₄. Due to the slower crystallization kinetics from the amorphous precursor, the thermodynamically stable region for o-LiTaOCl₄ should locate at higher temperatures (>300 °C) [4,14,15].

Otherwise, an unknown phase was found in o-LiTaOCl₄, and LiCl (~2.8 wt.%) was found in o-LiNbOCl₄. During crystal structure resolution and refinement, passing incomplete crystallization from the amorphous phase led to very broad peaks, further limiting the ability to conclusively fit complex (orthorhombic-/tetragonal-/monoclinic-) structure models or to ascertain the presence of structural distortions [4,14,15].

Therefore, challenges exist in elaborate material synthesis, crystal structure construction/refinement, electrochemical stability, and applications well over ambient temperatures. Herein, we report new SSEs of LTOCs that exhibit outstanding ionic conductivity (>7 mS cm⁻¹) at RT.

2. Materials and Methods

2.1. Material Preparation

Thermodynamic deficiency (metastable nonperiodic features) existed in HEBM methods [4,14,15,19]. To synthesize crystalline LiTaOCl₄-type SSEs, a preliminary experiment was conducted in quartz tubes at 180–500 °C for 2–10 h using LiOH and TaCl₅ as the raw materials (Figure S1). Pure m-LiTaOCl₄ phase was attained via Equation (S1) by calcining at 260 °C for 2 h and 180 °C for 10 h (Figure S1a,b). But the obtained σ values were as low as 0.47 and 0.24 mS cm⁻¹, respectively (Figure S1c). Other samples were extremely impure (Figure S1a,b), which reacted as in Equation (S2). They attained σ values ranging from 0.01 to 1.48 mS cm⁻¹ (Figure S1c). Except for m-LiTaOCl₄, at least three phases (m-LiTaO₂Cl₂, o-Li₂Ta₂O₅Cl₂, and m-Li₂Ta₂O₅Cl₂) were found in the samples. Moreover, quite a few unreacted TaCl₅ crystals were found on the wall of the quartz tube after calcining.

Because the resultants included gaseous HCl while the raw material of TaCl₅ was extremely easy to sublimate and evaporate (m.p. = 223.75 °C; b.p. = 229.55 °C [21]), the closed circumstances of a quartz tube were not suitable to synthesize the target LiTaOCl₄-type SSEs using LiOH and TaCl₅ as the raw materials.

Therefore, a new multi-factor experimental design was adopted in this work (Figure S2; Table S1). Samples of series A (Li_1+xTaO_1+xCl_4−x, −0.70 ≤ x ≤ 0.50) and series B (LiTaO_2+yCl_2−2y, −1.22 ≤ y ≤ 0) were prepared in quartz tubes at different calcining temperatures (T = 260/380/460 °C) and dwell times (D = 2/10 h). By using Li₂O, TaCl₅, and LiTaO₃ as the raw materials, gaseous resultants such as HCl were avoided.

The commercial Li₂O (Aladdin Co., Ltd., Shanghai, China; 99.99 wt.%), TaCl₅ (Aladdin Co., Ltd., Shanghai, China; 99.99 wt.%), and LiTaO₃ (Energy Chem. Co., Ltd., Xiamen, China; 99.9 wt.%) powders were mixed, pressed into tablets (90 MPa), and sealed in quartz tubes (26 cm³/tube; 1 g/tube) under vacuum (5 Pa). After calcining, the tablets were taken out, ground to powder, and preserved in the glove box.

Firstly, the reaction process (Equation (8)) was designed for series A (Li_1+xTaO_1+xCl_4−x, −0.70 ≤ x ≤ 0.50).

$$0.4 c-{\mathrm{Li}}_2\mathrm{O}+0.2 h-{\mathrm{LiTaO}}_3+{\displaystyle \frac{ 0.8-0.2x }{1+x}}m-{\mathrm{TaCl}}_5 \stackrel{ 260-460 °\mathrm{C} \times 2-10 \mathrm{h}, 5 \mathrm{P}\mathrm{a} }{\to }{\displaystyle \frac{1}{ 1+x}} m-{\mathrm{Li}}_{1+x}{\mathrm{TaO}}_{1+x}{\mathrm{Cl}}_{4-x} (\mathrm{main})$$

(8)

Pure m-LiTaOCl₄ phase could be attained under some synthesis conditions. But the corresponding σ values were very low, similar to the results of the preliminary experiment. In some samples with high σ values (5.20−7.35 mS cm^–1), three phases (m-LiTaO₂Cl₂, o-Li₂Ta₂O₅Cl₂ and m-Li₂Ta₂O₅Cl₂) coexisted according to Equation (9), while m-LiTaO₂Cl₂ was the main phase.

$$\begin{array}{c}0.4 c-{\mathrm{Li}}_2\mathrm{O}+0.2 h-{\mathrm{LiTaO}}_3+{\displaystyle \frac{ 0.8-0.2x }{1+x}}m-{\mathrm{TaCl}}_5 \to m-{\mathrm{LiTaO}}_2{\mathrm{Cl}}_2 (\mathrm{main})+m-{\mathrm{LiTaOCl}}_4 (2\mathrm{nd} \mathrm{phase})+\\ o-{\mathrm{Li}}_2{\mathrm{Ta}}_2{\mathrm{O}}_5{\mathrm{Cl}}_2 (3\mathrm{rd} \mathrm{phase})\end{array}$$

(9)

Secondly, another reaction process (Equation (10)) was designed for Series B (LiTaO_2+yCl_2−2y, −1.22 ≤ y ≤ 0) because samples with high σ values in series A had m-LiTaO₂Cl₂ as the main phase.

$$0.2 c-{\mathrm{Li}}_2\mathrm{O}+{\displaystyle \frac{ 0.6+0.4y }{1-y}}h-{\mathrm{LiTaO}}_3+0.4 m-{\mathrm{TaCl}}_5 \stackrel{ 380-460 °\mathrm{C} \times 2-10 \mathrm{h}, 5 \mathrm{P}\mathrm{a} }{\to }{\displaystyle \frac{1}{ 1-y}} m-{\mathrm{LiTaO}}_{2+y}{\mathrm{Cl}}_{2-2y} (\mathrm{main})$$

(10)

There were two side-reactions (Equations (11) and (12)) accompanying Equation (10):

$$0.2 c-{\mathrm{Li}}_2\mathrm{O}+{\displaystyle \frac{ 0.6+0.4y }{1-y}}h-{\mathrm{LiTaO}}_3+0.4 m-{\mathrm{TaCl}}_5 \to m-{\mathrm{LiTaOCl}}_4 (\mathrm{main})+m-{\mathrm{LiTaO}}_2{\mathrm{Cl}}_2 (2\mathrm{nd} \mathrm{phase})$$

(11)

$$0.2 c-{\mathrm{Li}}_2\mathrm{O}+{\displaystyle \frac{ 0.6+0.4y }{1-y}}h-{\mathrm{LiTaO}}_3+0.4 m-{\mathrm{TaCl}}_5 \to h-{\mathrm{LiTaO}}_3 (\mathrm{main})+m-{\mathrm{LiTaO}}_2{\mathrm{Cl}}_2 (2\mathrm{nd} \mathrm{phase})$$

(12)

Samples in series A and B were labelled, respectively, as A_x/T-D^# and B_y/T-D^# where the ±|x| indicated over-/under-lithiated values in Series A and ±|y| were peroxy-/anoxic values in Series B, respectively. For example, sample A_−0.60/380-2^# represented Li_0.40TaO_0.40Cl_4.60 (x = −0.60) calcined at 380 °C for 2 h, sample B_−0.67/460-10^# represented LiTaO_1.33Cl_3.33 (y = −0.67) calcined at 460 °C for 10 h, and so on. Seven representative samples with high phase purity and/or high σ values were emphatically studied, where I^#—A_−0.60/380-2^#, II^#—A_−0.60/460-10^#, III^#—A₀/460-10^# (= B₋₁/460-10^#), IV^#—A_0.10/460-10^#, V^#—A_0.20/460-10^#, VI^#—B_−1.22/460-10^#, and VII^#—B_−0.67/460-10^#.

2.2. Material Characterization

The morphology and elemental distribution of the electrolyte powders were tested using a scanning electron microscope (SEM; TESCAN MIRA LMS, Brno, Southern Moravia, Czech Republic) equipped with an energy dispersive spectrometer (EDS; Xplore 30, Oxford Instr., Oxford, England, UK). Total elemental analyses except Cl (because of the chloric solvent) were performed using an inductively coupled plasma-optical emission spectrometer (ICP; Prodigy 7, Teledyne Leeman Labs. Inc., Mason, OH, USA). Single-element analysis (Li and Ta) was performed using an atomic absorption spectrophotometer (AAS; Agilent 200AA, Agilent Technologies Inc., Santa Clara, CA, USA).

Electrovalence was determined using an X-ray photoelectron spectrograph (XPS; ESCALAB 250Xi, Thermo Fisher Sci., Waltham, MA, USA) equipped with a focused monochromatized AlK_α X-ray source. Binding energy was calibrated by the adventitious C 1s peak (284.8 eV). The parameters for survey/narrow scans were as follows: pass energy 100/30 eV, spot size 600/300 nm, number of scans 1/5, dwell time 100/50 ms, and energy step size 1.00/0.10 eV, respectively. Spectra were analyzed with Advantage v.5 software using a Gaussian–Lorentzian peak shape function and removing the Shirley background contribution from the data.

Phase identification of the powder electrolytes was carried out using an X-ray diffractometer (XRD) with CuK_α1 radiation (λ = 0.1540598 nm, 40 kV, 40 mA) at a continuous scan rate of 0.02 °/s (Empyrean, PANalytical B.V., Almelo, The Netherlands). Samples were isolated from water and oxygen via a special sample stage (Deen Optics Co., Shanghai, China) with shielding argon and Kapton tape. Structure refinements were performed using the Rietveld method implemented in GSAS/EXIGUI software [22] with the models of m-NaNbO₂F₂ (CSD 23,238/pdf 22-1392, P1 2₁/c 1, No. 14) [23], m-LiNbOCl₄ (Cc, No. 9) [15], and o-Li₂Nb₂O₅F₂-I (Cmc2₁, No. 36) [24].

Electrochemical impedance spectra (EIS) were determined via assembled blocking cells of SS//LTOC//SS (SS—stainless steel) under a constant pressure of 90 MPa on an electrochemical workstation (CHI660e, Shanghai Chenhua Instr. Co., Ltd., Shanghai, China). Testing frequencies ranged from 1 Hz to 1 MHz with a voltage amplitude of 5 mV at 25−75 °C. Therein, a blocking cell was assembled as follows: First, 260 mg powder of LTOC (series A or B) was pressed into a sheet (ϕ12 × 0.8 mm) under 315 MPa for 1 min. Then, it was sandwiched between two pieces of SS foil in a homemade battery testing unit, which provided a water/air-proof, heating, pressing, and electrode contact environment [25]. The ionic conductivity (σ) of LTOCs was obtained from the EIS results and calculated according to the formula σ = L/(RS), where L represents the thickness of electrolyte plate, R is the resistance (Z′), and S the surface area of the electrode. The activation energy (E_a) was calculated using the Arrhenius formula σ = Aexp$(-{\displaystyle \frac{E_a}{{\mathrm{k}}_{\mathrm{B}}T}})$, where A is the pre-exponential factor, T is the absolute temperature, and k_B is the Boltzmann constant.

Initial charge–discharge curves were tested via assembled ASSLIBs of NCM//LTOC//LPSC//Li-In (NCM—composite cathode; LPSC—Li₆PS₅Cl) under a constant pressure of 90 MPa on a battery test system (CT2001A, Wuhan LAND Electronic Co., Ltd., Wuhan, Hubei, China). Testing cut-off voltages ranged from 2.0 to 3.7 V at 35 °C and the current density was 0.2 C (1 C = 200 mA g⁻¹). Therein, the NCM was prepared by grinding (for 1 h) powders of LTOC (one of the above-mentioned seven representative samples I^#−VII^#) and commercial LiNi_0.8Co_0.1Mn_0.1O₂ (NCM₈₁₁, 2–4 μm, 99 wt.%; Shenzhen Huaqing New Mater. Technol. Co., Ltd., Shenzhen, Guangdong, China) while LTOC:NCM₈₁₁ = 30:70 (wt.%). The dual-layer LTOC//LPSC was used as the electrolyte. An ASSLIB was assembled as follows: First, 140 mg powder of LTOC was pressed into a sheet (ϕ12 mm) under 90 MPa for 1 min. Then, 70 mg LPSC and 15 mg NCM were pressed successively on two sides of the LTOC sheet under 315 MPa for 1 min. Finally, a Li-In foil was pressed on the LPSC side of the sheet under 90 MPa. All processes were conducted in an argon-filled glove box.

3. Results

3.1. Chemical Compositions (ICP/AAS)

Samples I^#−VII^# had high contents of Li and Ta with trace Al, K, Ca and Ni (Figure 1). The ICP/AAS results were all close to the designed compositions, while the ICP values were nearer. The mole ratios of Li/Ta via ICP in I^# and II^# were 0.386 and 0.399, respectively, corresponding to the phase content results in XRD patterns (see later in Section 3.3) and Table S1. Namely, pure phases of m-Li_0.40TaO_0.40Cl_4.60 (A_−0.6/380-2^#) and m-Li_0.40TaO_0.40Cl_4.60 (A_−0.6/460-10^#) existed in I^# and II^#. Besides I^# and II^#, the mole ratios of Li/Ta via ICP in other samples were 0.956−1.214, also corresponding to the phase content results in XRD patterns (see later in Section 3.3) and Table S1. Even though three phases (m-LiTaO₂Cl₂, m-LiTaOCl₄ and o-Li₂Ta₂O₅Cl₂) coexisted in III^#−VII^#, their mole ratios of Li/Ta were all equal to 1. The above results also indicated that all of the raw material of TaCl₅, being extremely easy to sublimate and evaporate, had reacted completely.

3.2. Morphology, Elemental Distribution (SEM/EDS), and Ta-Ion-Valence (XPS)

SEM images (I^#(a,b)–VII^#(a,b)) and EDS elemental mappings (1 k×; I^#(c,d)–VII^#(c,d)) are shown in Figure 2. The morphology of the powders varied from strips to irregular nano/microparticles. There were more thicker strips and coarser blocks in the lower-temperature and shorter-time calcined powder (I^# @380 °C × 2 h) than in the higher-temperature and longer-time calcined powders (II^#–VII^# @460 °C × 10 h). Changes in particle shape caused by chemical composition were very small among series A (II^#–V^# @460 °C × 10 h) and B (VI^#–VII^# @460 °C × 10 h). Elements of Ta, O, and Cl were distributed homogeneously in all particles.

By using XPS (Figure S3), the main peaks of Ta 4f_7/2/4f_5/2 centered at 26.73/28.63 eV for sample II^# and 26.77/28.66 eV for sample VI^#, respectively, were determined. They were assigned to Ta⁵⁺ species, which demonstrated the formation of LTOCs [26,27].

3.3. Phases and Crystal Structures

Figure 3 and Tables S2 and S3 show the XRD patterns, phase contents, and refined crystal structures of series A (Li_1+xTaO_1+xCl_4−x, −0.70 ≤ x ≤ 0.50) and B (LiTaO_2+yCl_2−2y, −1.22 ≤ y ≤ 0), including seven representative samples I^#−VII^#.

In series A, samples A_−0.70~0/260-2^#, A_{−0.70~−0.10}/380-2^#, and A_{−0.70~−0.40}/460-10^# were obtained via reaction Equation (8), while m-LiTaOCl₄ was the main phase. In other words, these samples had matching nominal-realizable compositions and belonged to LiTaOCl₄-type SSEs.

A_−0.10~0.50/380-2^# and A_−0.40~0.50/460-10^# were obtained via reaction Equation (9), where three phases coexisted, i.e., m-LiTaO₂Cl₂ (main), m-LiTaOCl₄ (2nd phase), and o-Li₂Ta₂O₅Cl₂ (3rd phase). There were relatively large differences between their nominal and realizable compositions, belonging more to LiTaO₂Cl₂-type SSEs.

In series B, samples B₋₁/380-2^# and B_{−1.22~−0.67}/460-10^# were obtained via reaction Equation (10), having matching nominal-realizable compositions, while m-LiTaO₂Cl₂ was the main phase. Sample B_−1.22/460-10^# (VI^#) attained the highest σ value (7.35 mS cm^–1) with the highest content of m-LiTaO₂Cl₂ in the two series (see later in Section 3.4).

B_−1.22/380-2^# and B_−0.67/380-2^# were obtained via reaction Equation (11), having unmatched nominal-realizable compositions, while m-LiTaOCl₄ and m-LiTaO₂Cl₂ were the main/2nd phases, respectively.

B_−0.67~0/380-2^# and B_−0.67~0/460-10^# were obtained via reaction Equation (12), while unreacted h-LiTaO₃ and m-LiTaO₂Cl₂ were the main/2nd phases, respectively.

The refined monoclinic structure of LiTaOCl₄ [28] (Figure 3h) belongs to the space group Cc (No. 9) with a cell volume (V) of 0.6395(3) nm³ (Z = 4). It consists of 1Li, 1Ta, 1O, and 4Cl ions occupying 4a sites in the lattice (Table S2). A polyanion group of trans-[O₂Cl₄] octahedra constitute the isolated (in the bc-plane) 1D framework chains along the a-axis and, being remarkably flexible, are connected by O but lack bridging Ta-Cl-Ta bonds. The Li ions occupy interstitial sites along the 1D chains, which exhibit an energetically favorable orientation and extensive disorder. This leads to multiple octahedra orientations and creates a frustrated 3D energy landscape with fast Li-ion diffusivity [4,14,15]. The evidence is from very widely dispersed Cl-Ta-Cl bond angles, which may originate from the empty Ta⁵⁺ d-orbitals.

The refined monoclinic structure of LiTaO₂Cl₂ [29] (Figure 3h) belongs to the space group P1 2₁/c 1 (No. 14) with a cell volume (V) of 0.3388(5) nm³ (Z = 4). It consists of 1Li, 1Ta, 2O, and 2Cl ions occupying 4e sites in the lattice (Table S3). A polyanion group of cis-[O₄Cl₂] octahedra share only their 4O vertices with each other to build infinite {TaO₂Cl₂}_∞ layers with outward-facing 2Cl vertices. The Ta ions displace toward O ions located trans to Cl ones, increasing the overall bonding strength and enhancing inherent distortions of cis-[O₄Cl₂] octahedra [23,30]. The zigzag {TaO₂Cl₂}_∞ layers are stacked along the a-axis and separated from each other by Li ions located in an irregular environment—a distorted octahedra with only Cl ions or 7-fold coordinated Cl plus the nearest O ion.

In short, there exist highly disordered Li ions coupled with prominent trans-[O₂Cl₄] framework flexibility in m-LiTaOCl₄ and abundant ordered Li ions uniformly distributed around the cis-[O₄Cl₂] framework with strong ionic interactions in m-LiTaO₂Cl₂. The synergistic effect facilitates a low activation energy for Li-ion diffusion and high ionic conductivity.

3.4. Ionic Conductivity and Activation Energy

Figure 4a shows that the σ values are dependent on the composition of series A (Li_1+xTaO_1+xCl_4−x, −0.70 ≤ x ≤ 0.50). For sample A/260-2^#, the σ values were low and changed a little (0.12−0.87 mS cm⁻¹) when −0.70 ≤ x ≤ 0. For samples A/380-2^# and A/460-10^#, the σ values increased gradually from 0.24/0.14 mS cm⁻¹ to peak values of 5.44/6.56 mS cm⁻¹ when −0.70 ≤ x ≤ 0.10, respectively (there were small experimental errors around x = −0.10). Conversely, for samples with larger x values than 0.10, the σ values decreased to 2.50/3.92 mS cm⁻¹ until x = 0.50, respectively.

In short, when the lithiated value (x) varied between −0.70 and 0.10, larger x, higher temperature, and longer time, ranging from 260 °C/2 h to 460 °C/10 h, attained higher crystallinity and larger σ values for SSEs in Series A.

Figure 4b shows that the σ values are dependent on the composition of series B (LiTaO_2+yCl_2−2y, −1.22 ≤ y ≤ 0). For sample B/380-2^#, the σ value increased gradually from 2.36 mS cm⁻¹ to a peak value of 4.63 mS cm⁻¹ when −1.22 ≤ y ≤ −0.25. Then, the σ value decreased to 2.38 mS cm⁻¹ until y = 0. For sample B/460-10^#, a few TaCl₅ crystals were found on the wall of the quartz tube after calcining when y ≤ −1.22. The σ value reached a peak of 7.35 mS cm⁻¹ when y = −1.22. Then, the σ value decreased gradually from the peak value to 1.31 mS cm⁻¹ until y = 0 (there was a small error when y = −1).

In short, when the anoxic value (y) varied between −1.22 and −0.67, lower y, higher temperature, and longer time, ranging from 380 °C/2 h to 460 °C/10 h, attained higher crystallinity and larger σ values for SSEs in Series B.

Figure 4c and Table S1 summarize the σ values in relation to chemical composition, synthesis temperature (T = 260−460 °C), and dwell time (D = 2−10 h). The ionic conductivity for each sample increased with the increase in temperature (and crystallinity), obeying Case 1 in Section 1 (a few anomalies resulted from test errors).

Figure 4d, Table S1, and Figure S4 show that the σ and E_a values are dependent on the synthesis temperature for seven representative samples. The E_a values were 23.84/26.26 kJ mol⁻¹ for samples I^#−II^#, respectively. And the E_a values ranged from 5.22 to 12.43 kJ mol⁻¹ for samples III^#−VII^#. The low E_a values of SSEs facilitate Li-ion diffusion.

Samples I^# and II^# had matching nominal-realizable compositions (m-Li_0.40TaO_0.40Cl_4.60) and were synthesized at 380 °C for 2 h and 460 °C for 10 h, respectively (Figure 3f,g). But their σ values were as low as 1.21 and 0.82 mS cm⁻¹, respectively. This should be related to their large under-lithiated value (−|x|) of −0.60 (Figure 4a).

Other three-phase coexisting samples (III^#−VII^#) attained high σ values ranging from 5.20 to 7.35 mS cm⁻¹. The corresponding weight ratios were 40.8–80.6 wt.% for m-LiTaO₂Cl₂ (major phase), 10.2–57.2 wt.% for m-LiTaOCl₄ (2nd phase), and 2.0–17.9 wt.% for o-Li₂Ta₂O₅Cl₂ (3rd phase), respectively. High conductivity of SSEs was determined by high contents of the m-LiTaO₂Cl₂/m-LiTaOCl₄ mixture, while the 3rd role of o-Li₂Ta₂O₅Cl₂ was small. Sample VI^#, derived from series B having an anoxic value (−|y|) of −1.22 (Figure 4b), attained the highest σ value with the highest content of m-LiTaO₂Cl₂.

3.5. Initial Charge–Discharge Curves of ASSLIBs

An ASSLIB (NCM//LTOC//LPSC//Li-In) was assembled with a dual-layer SSE in which the LTOC was taken from one of seven representative samples I^#−VII^#. The initial charge–discharge curves were obtained under cut-off voltages of 2.0–3.7 V at 0.2 C and 35 °C (Figure 5). Results showed that all ASSLIBs (using different SSEs) delivered large battery capacities from 159.8 to 179.9 mAh g⁻¹, accompanying high initial Coulombic efficiencies (CEs) from 76.4% to 84.9% (inset i of Figure 5). The two series of SSEs in this work have great application potential for next-generation ASSLIBs.

4. Conclusions

As the other pole of amorphous SSEs with high σ values working at lower temperatures, the developed crystalline Li-Ta-oxychlorides (LTOCs) with high σ values are promising candidates to work at higher temperatures, particularly for high-power and fast charging applications in ASSLIBs.

In this work, crystalline SSEs of series A (Li_1+xTaO_1+xCl_4−x, −0.70 ≤ x ≤ 0.50) and B (LiTaO_2+yCl_2−2y, −1.22 ≤ y ≤ 0) were prepared for applications well over ambient temperatures.

Larger x, higher temperature, and longer time, ranging from 260 °C/2 h to 460 °C/10 h, attained higher crystallinity and larger σ values for SSEs in series A when the lithiated value (x) varied between −0.70 and 0.10. Lower y, higher temperature, and longer time, ranging from 380 °C/2 h to 460 °C/10 h, attained higher crystallinity and larger σ values for SSEs in series B when the anoxic value (y) varied between −1.22 and −0.67.

The optimized crystalline SSEs of LTOCs had lithium superionic conduction up to 7.35 mS cm⁻¹. High conductivity of SSEs was determined by high contents of the m-LiTaO₂Cl₂/m-LiTaOCl₄ mixture, while the 3rd role of o-Li₂Ta₂O₅Cl₂ was small. The synergistic effect facilitates a low activation energy for Li-ion diffusion and high ionic conductivity.