Polyanion-stabilized amorphous halide electrolytes with low lithium content for all-solid-state lithium batteries

Domain: chemistry
Published: 2026-02-28
Editorial outcome: minor_revision

Abstract
Lithium halide solid electrolytes have garnered significant attention owing to their high ionic
conductivity and positive electrode compatibility. However, achieving target ionic conductivity
typically requires high lithium concentration (> 4.3 wt%) within optimal structure, which raises
costs and exacerbates air sensitivity. Here, we leverage anion clusters to synthesize a series of
amorphous halide electrolytes xLi SO -ZrCl , with optimal ionic conductivities of 1.5 mS cm-1 at
2 4 4
30 ℃ and a significantly reduced lithium content of 2.4 wt%, alongside good air stability. Through
neutron/synchrotron X-ray experiments, first-principles calculations and machine learning-
accelerated molecular dynamics simulations, we reveal a disordered backbone of [Zr Cl (SO )]2-
a 4a 4
(1 ≤ a ≤ 4) that enables fast Li-ion diffusion via under-coordinated oxygen sites. All-solid-state
lithium batteries employing these electrolytes and LiNi Co Mn O positive electrode exhibit
0.8 0.1 0.1 2
81.1% capacity retention after 1400 cycles at 1 C (60 min) and 30 ℃. Our findings reveal anion-
cluster chemistry as an approach that transforms solid electrolyte design for advanced batteries,
bridging materials science with practical energy storage innovation.

Introduction
Next-generation electronics and electric vehicles demand advanced all-solid-state lithium
batteries (ASSLBs) for their safety, high energy density, and long cycle life1, 2. The cornerstone of
ASSLB technology is the development of Li-ion solid electrolytes (SEs).3 Oxides, sulfides and
halides, have been explored over the past decades4, 5. Oxides (e.g., Li La Zr O ) exhibit excellent
7 3 2 12
thermal and oxidative stability, but require high-temperature sintering due to their rigidity
structure6-8. Sulfides (e.g., Li GeP S , Li PS Cl ), while offering high room-temperature
10 2 12 6-x 5-x 1+x
ionic conductivity via cold pressing, have limited electrochemical stability and poor environmental
resistance9-11. Halide SEs, commonly represented by ternary Li-M-X systems (M = Y, In, Sc, Zr,
Sm, La, Al, Ta, or Nb; X = F, Cl, Br, or I), have garnered significant research interest due to their
high ionic conductivity, wide electrochemical window, and reasonable deformability12, 13,14-27.
Lithium concentration is a key factor influencing ionic conduction in Li-M-X SEs. In general,
a high lithium content (> 4.3 wt%) is required in these high-conducting halide SEs, as shown in
Fig. 1c. For instance, Li ZrCl 28is a promising low-cost SE but has a limited ionic conductivity of
2 6
~0.4 mS cm-1. Substitution of ZrI4+ by trivalent cations (e.g., Fe3+, In3+, Sc3+, or Al3+)28-30 increases
lithium content, enhancing ionic conductivity from 0.4 mS cm-1 to typically > 1.0 mS cm-1.
Theoretical study suggests that an optimal 40-60% Li octahedral occupancy (corresponding to a
lithium weight percentage of ~4.8-6.7% in representative halides, Supplementary Table 1) is
crucial for achieving high ionic conduction in close-packed halide SEs 31. Very recently, this
finding has been verified experimentally32. However, achieving high conductivity while reducing
lithium content remains a significant challenge.

The high lithium content will increase the SE cost, especially when considering the limited
global lithium reserves (Supplementary Fig. 1) and their upward fluctuations33. For example, the
price of LiCl (anhydrous, 98%) in 2024 (Supplementary Tables 2 and 3) is about 20 times higher
than that in 202134. Therefore, it is crucial to explore new strategies for decreasing a SE's ultimate
cost by decreasing lithium content, owing to the multiple variables like supply-demand
fluctuations, synthesis and manufacturing complexities, etc5. Additionally, reducing the Li-ion
content is crucial for improving the air stability of SE, as the high charge density of small Li+
cations leads to strong electrostatic interactions with polar H O molecules35. As a result,
2
spontaneous hydrolysis reactions occur when halide SEs are exposed to air, leading to a drastic
decay in performance, as seen in all reported halide SEs36-38. This issue presents a significant
challenge for large-scale production and processing.

In this work, we report a family of low-lithium-content amorphous SEs by combining clusters
of [SO ]2- and [ZrCl ] using a simple ball milling method. By regulating the anion proportions,
4 4
these materials achieve a significantly reduced Li content of 2.4 wt% while maintaining high Li+
conductivity of 1.5 mS cm-1 at 30 ℃. To uncover the structural origin of the high conductivity in
this low-lithium-content amorphous system, we performed pair distribution function (PDF)
analysis based on the neutron total scattering and synchrotron X-ray data, integrated with multi-
scale modeling under a self-constructed accurate and efficient workflow, including first-principles
calculations and molecular dynamics simulations (MD) based on machine-learning force field
(MLFF). We revealed an interesting, disordered backbone consisting of four complex anions
clusters of [Zr Cl (SO )]2- (1 ≤ a ≤ 4) interconnected by sulfate ions, forming a percolation
a 4a 4
network for Li-ion with frustrated energy landscapes. The ionic conduction in the amorphous
electrolyte is primarily attributed to Li-ion hopping between sites with low oxygen coordination.
ASSLBs utilizing the SEs and LiNi Co Mn O positive electrode demonstrate good cycling
0.8 0.1 0.1 2
stability for over 1400 cycles at 1.0 C and 30 ℃. In addition, the SEs deliver decent air stability
under 30% relative humidity. Our work redefines the design paradigm for halide SEs,
demonstrating that strategic anion-cluster manipulation in amorphous systems can circumvent
traditional lithium-content limitations, thereby advancing cost-effective, high-energy-density all-
solid-state batteries.

Results
Synthesis and conduction properties
The xLi SO -ZrCl SEs (0 < x  1.0) with different lithium content were mechanochemical
2 4 4
synthesized from a stoichiometric mixture of Li SO and ZrCl . Arrhenius plots of xLi SO -ZrCl
2 4 4 2 4 4
SEs, derived from the electrochemical impedance spectroscopy (EIS) at different temperatures,
were compared for varying compositions (x) in Supplementary Figs. 2-3. The resultant ionic
conductivities and activation energies (E ) are calculated as shown in Fig. 1a. The ionic
a
conductivity increased with the gradually increasing Li SO -to-ZrCl ratio, reaching to the highest
2 4 4
value of 1.5 mS cm-1 when x = 0.5, and then decreased with further increasing Li SO content.
2 4
Correspondingly, 0.5Li SO -ZrCl show the lowest activation energy (0.33 eV), which is
2 4 4
comparable with typical halide SEs like Li ScCl (0.33 eV)19 and Li InCl (0.347 eV)39.
3 6 3 6

A series of 2xLiCl-ZrCl SEs (0 < x  1.0) with the same lithium molar ratio were also
4
prepared for comparison. As shown in Fig. 1b, the ionic conductivities of 2xLiCl-ZrCl SEs are
4
generally lower than that of xLi SO -ZrCl , and increased with the increase of x. The highest ionic
2 4 4
conductivity of 0.58 mS cm-1 is achieved at x = 1.0 (2LiCl-ZrCl ), which is identical with the
4
reported results34, 40. When 2xLiCl-ZrCl system adopts the same low Li/Zr ratio as that in
4
0.5Li SO -ZrCl , that is LiCl-ZrCl (x = 0 .5), the ionic conductivity is only 0.24 mS cm-1, one
2 4 4 4
order of magnitude lower than that of 0.5Li SO -ZrCl . These results suggest that reduced lithium
2 4 4
content in Li-Zr-Cl system is impractical for achieving high ionic conductivity, which however is
achievable in the dual-anion clustered structure proposed in this work.

It’s worth noting that the lithium content in our 0.5Li SO -ZrCl is only 2.4 wt%, substantially
2 4 4
lower than that in representative halide (5~6 wt%) and sulfide (typically > 7 wt%) SEs with the
desirable high ionic conductivity, as shown in Fig. 1c and Supplementary Tables 2-8. Notably, Zr-
based halide SEs inherently possess a cost advantage over halide SEs based on other metals15, and
our design further amplifies this benefit through two key features: first, the low-lithium-content
configuration reduces reliance on uncertain lithium resources; Second, unlike the precursors of
other selected polyanion-based SEs, the use of cost-effective Li SO as a precursor enables our
2 4
electrolyte to maintain stable ionic conductivity even when lower-purity raw materials are
employed (Supplementary Figs. 4-5). This distinctive combination not only drives a significant
reduction in material costs but also enhances the practicality of 0.5Li SO -ZrCl for future large-
2 4 4
scale applications, strengthening its competitiveness in the SE landscape.

In addition, the electronic conductivities of xLi SO -ZrCl SEs were measured by direct
2 4 4
current (DC) method and displayed in Supplementary Fig. 6. The electronic conductivities of
xLi SO -ZrCl SEs are at the level of 10-9 S cm-1, which are significantly lower than their ionic
2 4 4
conductivities and supports the application of these materials as a separator for ASSLBs.

Air stability
Beyond conduction property, practical application of these materials requires environmental
stability. Conventional halide-based SEs often suffer from moisture sensitivity, which induces by-
product formation and reduces ionic conductivity, thereby hindering scalable production and
storage. Encouragingly, the 0.5Li SO -ZrCl SE demonstrates good air stability. As shown in Fig.
2 4 4
1d and Supplementary Figs. 7-8, the fresh 0.5Li SO -ZrCl SE didn’t display noticeable change
2 4 4
in both ionic conductivity and phase after exposure in dry room or Ar with a relative humidity (RH)
of ~1% at room temperature (dew point: -40 ℃), suggesting the structural robust under dry air.
Moreover, 2LiCl-ZrCl reference sample was also tested under the same condition for comparison.
4
It has been recognized for the decent air stability among all the halide SEs34, and displayed
similarly high air stability in dry room. However, under a harsher air condition (e.g., 30% RH),
2LiCl-ZrCl shows a more pronounced conductivity decrease (Fig. 1d), accompanied with a faster
4
moisture absorption (Fig. 1e-f), reflecting the higher air stability of 0.5Li SO -ZrCl .
2 4 4
Moreover, in-situ XRD patterns reveal striking differences in moisture stability between the
two electrolytes. 0.5Li SO -ZrCl exhibits minimal changes in XRD patterns even after ~60 min
2 4 4
of exposure, with only trace amounts of Li SO and ZrOCl ·6H O (PDF#00-018-1498) detected
2 4 2 2
(Supplementary Fig. 9a-b). In contrast, 2LiCl-ZrCl undergoes complete phase transition within
4
~15 min (Supplementary Fig. 9c-d). Analysis of decomposition products after air exposure
(Supplementary Fig. 9 and Supplementary Note 2) showed a significant amount of lithium salts
decomposed from 2LiCl-ZrCl , which correlates with its lower ionic conductivity. Given its better
4
air stability with relative low lithium content (Supplementary Fig. 10-11, and Supplementary Note
2), 0.5Li SO -ZrCl alleviates the stringent requirements for low dew points typically necessary
2 4 4
for the fabrication and storage of halide SEs. Coupled with its highly competitive cost, 0.5Li SO -
2 4
ZrCl emerges as a promising SE candidate for large-scale, industrially viable production.
4

Amorphous structure analysis
Given that xLi SO -ZrCl possess high ionic conductivity, the relationship between this
2 4 4
property and their structure is of interest. X-ray diffraction (XRD) patterns of the as-prepared
xLi SO -ZrCl are shown in Fig. 1g. The XRD pattern of the 0.5Li SO -ZrCl sample displays no
2 4 4 2 4 4
distinct Bragg diffraction peaks except for a weak Li SO peak (Supplementary Fig. 12), indicating
2 4
the dominating amorphous structure. An impurity phase of trigonal Li ZrCl appeared with x > 0.5,
2 6
while ZrCl reactant remained with x < 0.5. The evolution of phase in products and the
4
corresponding ionic conductivities described above indicate that the higher ionic conductivity is
most likely arise from the amorphous structure, which is further verified by the decreased ionic
conductivity of calcinated xLi SO -ZrCl with higher degree of crystallinity (Supplementary Fig.
2 4 4
13).For comparison, the single-anion system 2xLiCl-ZrCl maintain crystalline phase under
4
varying x (Supplementary Fig. 14), with significantly lower ionic conductivities compared to
0.5Li SO -ZrCl .
2 4 4

Furthermore, the microstructure of 0.5 Li SO -ZrCl was analyzed by transmission electron
2 4 4
microscopy (TEM) measurement. No obvious lattice fringe can be found in the TEM images and
the corresponding selective area electron diffraction (SAED) pattern shows merged rings with halo
feature (Fig. 1h), confirming the amorphous nature of the sample. TEM energy dispersive spectra
(EDS) elemental mapping for 0.5Li SO -ZrCl in Supplementary Fig. 15 demonstrates that the
2 4 4
elements Zr, S, Cl and O are homogeneously distributed. Synchrotron-based XRD (sXRD) further
confirm the amorphous nature (Supplementary Fig. 16). These results further demonstrate that the
raw materials fully react with each other, and formed a highly amorphous structure of 0.5Li SO -
2 4
ZrCl .
4

The amorphous structure of 0.5Li SO -ZrCl was then investigated by using Raman
2 4 4
spectroscopy. As shown in Fig. 2a, Raman spectra of 0.5Li SO -ZrCl represents a weak polarized
2 4 4
band at 404 cm-1 and a depolarized band of medium intensity in the range of 50-150 cm-1, similar
to the A modes of zirconium chloride octahedron in ZrCl 41. It is reasonable to speculate that the
g 4
octahedron units from the ZrCl precursors are likely to be presented in the amorphous 0.5Li SO -
4 2 4
ZrCl . Notably, ZrCl adopts monoclinic structure with a P2/c (C 4) space group, composing of
4 4 2h
zigzag polymer chains of [(ZrCl )Cl ] with bridged octahedra. Three pairs of distinct Zr–Cl
4/2 2 n
bonds (two bridging bonds and one terminal bond) form as a result of the highly distorted octahedra
with edge sharing42. The bands at 281 and 217 cm-1 are attributed to the two bridging Zr-Cl
b
bonds43, while the band at ~409 cm-1 for terminal Zr–Cl is similar to that (~412 cm-1) of the vapor
t
complex ZrAl Cl in which Zr is 6-fold coordinated44. Fig. 2c exhibits the symmetric stretching
2 8
vibrations (A ), the antisymmetric stretching vibrations of A and B , and the bending motion (A ,
g g g g
B ) of six-coordinated zirconium chloride octahedron45. The vibration bands of ZrCl after ball-
g 4
milling (BM-ZrCl ) were maintained but became broader, with a slight shift to lower energies in
4
the 0.5Li SO -ZrCl product. These observations suggest the existence of octahedral fragments of
2 4 4
ZrCl in 0.5Li SO -ZrCl that however has a more disordered structure27. Meanwhile, the bands at
4 2 4 4
~161 cm-1 and ~325 cm-1, characteristic of 2LiCl-ZrCl with face-sharing zirconium chloride
4
octahedron, are likely still present. In addition, Raman vibration band in a range of 400-1250 cm-1
of 0.5Li SO -ZrCl displays weak signals compared to Li SO (Fig. 2b). The vibrational modes
2 4 4 2 4
of Li SO in Fig. 2d shows a tetragonal T symmetry for S-O tetrahedron46, including symmetric
2 4 d
S–O stretching mode (ν ) at about 1010 cm-1, symmetric O–S–O bending mode (ν ) around 400–
1 2
500 cm-1, asymmetric S–O stretching mode (ν ) at 1050~1250 cm-1, and asymmetric O–S–O
3
bending mode (ν ) at 500~700 cm-1. Among them, the ν and ν modes represent the bending
4 2 4
vibration modes related to the rotation of polyanion SO 2-. It can be seen that these two vibration
4
bands for BM-Li SO and 0.5Li SO -ZrCl are highly broadened compared to those of Li SO
2 4 2 4 4 2 4
starting material, indicating that SO 2- undergoes order-to-disorder transition during ball-milling
4
in order to form the interconnected backbone47.

The local structure of amorphous 0.5Li SO -ZrCl was further explored by pair distribution
2 4 4
function (PDF) analysis based on the synchrotron X-ray and neutron total scattering data. As
shown in Fig. 2e, the synchrotron PDF (sPDF) pattern of 0.5Li SO -ZrCl exhibits strong
2 4 4
oscillation in the atomic pair distance range of r < 8.0 Å, and become significantly damped beyond
8.0 Å, indicating the completely random structure with loss of long-range order (LRO) in the
amorphous product. In contrast, ZrCl and 2LiCl-ZrCl display sharp peaks in the sPDF pattern
4 4
across a wide range of r (1 to 50 Å, Fig. 2e and Supplementary Fig. 17), indicating the presence
of LRO. These observations align well with the previously discussed XRD and Raman results.
Moreover, the PDF of the amorphous samples can be divided into two main regions of A and B.
Region A (r < 5.31 Å) represents the internal local structure of the primary constituents or building
blocks, suggesting a similar short-range order (SRO) structure consistent with that of ZrCl and
4
2LiCl-ZrCl . Region B (5.31 Å < r < 8 Å) provides insight into the interplay (coherence) between
4
these building blocks. A slight oscillation at 7.79 Å is still observable, consistent with the Cl–Cl
distance (7.89 Å) in the edge-sharing [ZrCl ] dimer. This finding suggests that despite the less
4 2
well-ordered state, the fundamental structure of amorphous 0.5Li SO -ZrCl remains primarily
2 4 4
composed of these dimer units.

The comparison of nPDF and sPDF G(r) experimental results of 0.5Li SO -ZrCl and 2LiCl-
2 4 4
ZrCl , respectively, are displayed in Fig. 2f, and the simulated nPDF G(r) values of Li ZrCl , ZrCl
4 2 6 4
and Li SO (see Methods) are shown in Supplementary Fig. 18. Compared to the sPDF, the nPDF
2 4
shows less resolved features owing to the strong incoherent scattering cross-sections of Cl (5.3
barn) and its high relative content48. Therefore, the more obvious Zr-Cl bond distance was
observed in both sPDF tests, while the more obvious Cl–Cl bond distance was observed in nPDF,
which indicates that the SRO structure of 0.5Li SO -ZrCl is similar to that of 2LiCl-ZrCl
2 4 4 4
with face-sharing Zr-center octahedra. However, the Zr–Cl of 2.41Å and Cl–Cl of 3.52Å bond distances
in 0.5Li SO -ZrCl become shorter relative to 2LiCl-ZrCl , which implies a more contracted
2 4 4 4
polyhedral structure. S–O of around 1.5Å bond distance for 0.5Li SO -ZrCl is observed,
2 4 4
indicating the strong bond in SO 2- cannot be easily destroyed. In both nPDF and sPDF tests, the
4
peaks corresponding to the Li–O and Li–Cl coordination cannot be observed, which is related to
the contraction of the adjacent Zr–Cl bond. The contraction of Zr–Cl bond enlarges the size of the
Li+ transport channel and ultimately promotes rapid ion diffusion.

To further elucidate the local coordination environment around Zr, we performed Zr K-edge
X-ray absorption spectroscopy (XAS), including X-ray absorption near-edge structure (XANES)
and extended X-ray absorption fine structure (EXAFS) measurements. Supplementary Fig. 19a
compares the spectral features of 0.5Li SO -ZrCl along with reference spectra of ZrCl and ZrO .
2 4 4 4 2
The whiteline of 0.5Li SO -ZrCl showed three obvious splits, which were almost identical with
2 4 4
the whiteline shape for ZrCl 49. Thus, the amorphous structure in 0.5Li SO -ZrCl could be
4 2 4 4
regarded as the fragmented ZrCl domains with three kinds of distinct bond lengths. Moreover, the
4
more integrated whiteline shape of 0.5Li SO -ZrCl indicates that bond lengths heterogeneity
2 4 4
within Zr–Cl/O polyhedrons becomes less pronounced after O substitution (from SO 2-). To
4
quantify the coordination environment around central Zr atoms, we employed Zr K-edge Fourier-
transformed EXAFS (FT-EXAFS) spectroscopy (Supplementary Fig. 19b). Using single scattering
information from ZrO (Zr–O bonds) and ZrCl (Zr–Cl bonds), we identified peaks corresponding
2 4
to O and Cl scatterers around Zr. The Zr-O peak intensity increases, while the Zr–Cl peak intensity
decreases, demonstrating that Cl in the immediate vicinity of Zr are replaced by O (from SO 2-).
4
In-depth analysis of phase-uncorrected wavelet transform (WT)-EXAFS (Supplementary Fig. 19c)
further confirms the sequential coordination of Zr by O (~1.5 Å), Cl (~2 Å), and Zr (~3 Å). EXAFS
fitting (Supplementary Fig. 19d) yield precise average coordination parameters: each central Zr
atom is surrounded by 2.6 O atoms (2.23 Å) and 3.0 Cl atoms (2.45 Å) (Supplementary Table 9).
This result is consistent with the contracted Zr–Cl polyhedral structure observed in PDF analysis,
reinforcing the reliability of our structural interpretations.

In order to determine the atomic-level amorphous structure, we further developed a parameter-
free, multi-scale modeling workflow (Supplementary Fig. 20), which is broadly applicable for
constructing various amorphous materials. The workflow began by acquiring the initial chemical
reaction between Li SO to ZrCl , referred to as cluster units, through ab-initio molecular
2 4 4
dynamics (AIMD) simulations (Fig. 2g). Then, spatially arranged cluster units were relaxed using
MD based on accurate machine-learning force field (MLFF, Supplementary Fig. 21), ultimately
leading to the formation of final amorphous structure. The resulting amorphous structures are
dense and free of voids (Fig. 2g) and the calculated density of 2.04 g/cm3 closely aligns with the
experimental value of 2.05 g/cm3.

The simulated amorphous structures of 0.5Li SO -ZrCl demonstrate good agreement with
2 4 4
the experimental PDF analysis (Fig. 2f). The theoretical sPDF and nPDF for the amorphous
structure were calculated by averaging over 100 configurations equidistantly sampled from the
MLFF-MD trajectory spanning 300 picoseconds (refer to Methods for details). All the peak
positions are in alignment with the experimental values in Fig. 2f. As illustrated in Fig. 3a-d and
Supplementary Fig. 22, the corresponding atomic pair contributions in the PDF are clearly resolved.
The peak at 1.48 Å corresponds to S–O bonding in SO 2-. Peaks at 2.15 Å represent Zr–O bonding
4
between SO 2- and ZrCl . Peaks around 2.4 Å signify Zr–Cl bonding and the peak at 3.5 Å
4 4
represents the Cl–Cl pair in the monomer of ZrCl . The peak at 4 Å indicates Zr–Zr distribution in
4
oligomers of [ZrCl ] (n > 1), as clarified in Fig. 3c. The relatively weak peak observed between
4 n
4.5 Å and 5 Å is predominantly attributed to Zr–Cl and Cl–Cl atomic pairs. In summary, the PDF
analysis indicates that SO 2- are embedded within the ZrCl through the Zr–O bonds.
4 4

As depicted in Fig. 3a-e, four principal anion clusters consist of the amorphous structure
frameworks: 1) Monomeric SO-Zr : the SO 2- are edge-sharing with ZrCl via two oxygen atoms,
1 4 4
2) Dimeric SO-Zr : the SO 2- bridges either an edge-sharing [ZrCl ] dimer or two isolated [ZrCl ]
2 4 4 2 4
clusters through oxygen coordination; 3) Trimeric SO-Zr : SO 2- bridges an edge-sharing [ZrCl ]
3 4 4 2
dimer while simultaneously binding to a third [ZrCl ] unit by chlorine, 4) Tetrameric SO-Zr : SO 2-
4 4 4
bridges a corner-sharing [ZrCl ] trimer (pre-interconnected via chloride bridges) to a fourth [ZrCl ]
4 3 4
cluster through oxygen coordination. Notably, the monomers SO-Zr, dimers SO-Zr and trimers
2
SO-Zr constitute nearly 99% of this amorphous structure, whereas the tetramers SO-Zr are
3 4
present in much lesser amounts (Fig. 3e). Apart from the involvement of SO 2-, these units share a
4
common feature: the coordination number of all zirconium atoms remains stable at 6
(Supplementary Fig. 23).

Li+transport mechanism
The Arrhenius relation for the amorphous structure was calculated based on the MLFF-MD
(see Methods). As shown in Fig. 3g, by analyzing the region where the mean squared displacement
(MSD) curves exhibit a linear time dependence (Supplementary Fig. 24–25), the room-temperature
ionic conductivity based on a 3-ns MLFF-MD simulation is about 1.7 mS cm-1, close to the
experimental value of 1.5 mS cm-1. The activation energy derived from the simulation is 0.32 eV,
which is in good agreement with experimental results (0.33 eV), further confirming the reliability
of our computational results.

To elucidate the conduction mechanism of Li+ in the disordered framework, we initially
examined the coordination environment of Li+. As shown in Fig. 3f, the coordination environment
is complex and can be classified into four main types based on the number of coordinating oxygen
atoms: (i) oxygen-free sites labeled as Li-Cl (x = 1-5), (ii) single-oxygen-coordinated Li-Cl O (x
x x
= 1-4), (iii) double-oxygen-coordinated Li-Cl O (x = 1-4), and (iv) triple-oxygen-coordinated Li-
x 2
Cl O (x = 0-2). These coordination environments evolve over time due to the mobility of Li+ that
x 3
is dictated by the kinetic energy of Li+ and the potential energy surface, potentially giving rise to
more than a dozen distinct coordination environments, including Li-Cl , Li-Cl , Li-Cl O, and Li-
5 4 3
Cl O, among others. All these local Li+ environments lack the periodically arranged LiCl
2 6
sublattice observed in Li ZrCl , instead adopting a fully disordered configuration. Notably, the
2 6
presence of diverse Li⁺ coordination environments creates a frustrated energy landscape50, which
promotes fast ionic conduction by facilitating percolation networks for Li⁺ transport. The spatial
distribution of Li+ probability at 300 K over the 3-ns trajectory in Fig. 3h reveals a percolating
isotropic diffusion pattern, indicating efficient Li⁺ transport through interconnected pathways.
Furthermore, in this amorphous framework, as shown in Supplementary Fig. 26a-b, approximately
50% of Li+ exhibit diffusion distances exceeding 2 Å, alongside 5% achieving distances greater
than 8 Å in this frustrated energy landscape over the 3-ns trajectory. These diffusion processes are
correlated with oxygen coordination. The shorter Li–O bond length (~2.0 Å compared to ~2.5 Å
of Li–Cl from radial distribution functions in Supplementary Fig. 26e) arises from stronger Li–O
interactions, which suggests that lower oxygen coordination environments favor Li+ diffusion, a
trend corroborated by statistical analysis in Supplementary Fig. 26c. In summary, the amorphous
framework creates various Li+ coordination environments that homogenizes the energy landscape,
thereby enabling room-temperature superionic conductivity.

Cell performance
In addition to the high ionic conductivity, the 0.5Li SO -ZrCl SE also possess favorable
2 4 4
deformability and good oxidation stability, which are two critical attributes for achieving high-
energy density ASSLBs. The powder and pellet morphologies of 0.5Li SO -ZrCl were examined
2 4 4
by scanning electron microscopy (SEM). As displayed in Supplementary Fig. 27, the powder
samples, hundreds of nanometers to a few micrometers in size, can be easily cold-pressed into
dense pellets, indicating good deformability. Atomic force microscopy (AFM) was further
employed to assess the Young’s modulus of 0.5Li SO -ZrCl , revealing a low average Young’s
2 4 4
modulus of ~2.0 GPa (Supplementary Fig. 28), a value comparable to the representative soft
electrolytes of AlOCl-LiCl51 and Li-Ta-Cl52, and lower than that of sulfide SEs (20 GPa)53. This
low Young’s modulus ensures tight contact between SE particles, which is essential for minimizing
interfacial resistance in ASSLBs. Moreover, the electrochemical stability of 0.5Li SO -ZrCl SE
2 4 4
was evaluated through linear sweep voltammetry (LSV), as shown in Supplementary Fig. 29. The
oxidation stability of 0.5Li SO -ZrCl is ~4.4 V vs. Li/Li+, which is identical to the calculated
2 4 4
value (Supplementary Figs. 30-31), and higher than that of sulfide SEs (~2.4 V vs. Li/Li+) and
competitive compared to the reported halides (~4.3 V vs. Li/Li+)14, 54.

As a proof of concept, bulk-type ASSLBs utilizing 0.5Li SO -ZrCl SE were constructed to
2 4 4
demonstrate the electrochemical performance of the amorphous halide-based SE. As shown in Fig.
4a, full cell with LiNi Mn Co O (NCM811) positive electrode was measured between 2.8 and
0.8 0.1 0.1 2
4.3 V vs. Li/Li+ at 0.1 C (1 C = 190 mA g-1), delivering a high initial discharge capacity of 209.9
mAh g-1 and a high Coulombic efficiency of 96.7%, indicating a reliable stability of 0.5Li SO -
2 4
ZrCl SE toward NCM811 positive electrode. As the specific current increased stepwise from 0.1
4
to 1.0 C, high discharge capacities can be maintained. Fig. 4b-c exhibits the specific discharge
capacities of 209.9, 188.2, 172.28, 157.5 and 121.9 mAh g-1 at the rates of 0.1, 0.2, 0.33, 0.5 and
1.0 C, respectively. Additionally, the specific discharge capacity of the full cell can recover to
194.7 mAh g-1 when the current density is returned to 0.1 C. Moreover, the full cell also shows
good long-term cycling stability at 1.0 C, with a capacity retention of 81.1% w.r.t first cycle after
1400 cycles and maintaining a discharge capacity of 101.7 mAh g-1 after 2500 cycles with an
average Coulombic efficiency over 99.99% (Fig. 4d).

To evaluate the practical relevance, we tested full cells with a high NCM811 loading of 38.95
mg cm-2. These achieved a high initial areal capacity of 6.4 mAh cm-2, and a robust capacity
retention of 82.4% after 300 cycles (5.3 mAh cm-2). Moreover, full cell with the 0.5Li SO -ZrCl
2 4 4
SE also can even endure a higher charge cut-off voltage of 4.6 V (Fig. 4f), delivering an initial
capacity as high as 218.2 mAh g-1 under 0.2 C and maintaining a capacity retention of 70.0% over
100 cycles. The SE also enabled stable cycle performance with LiCoO positive electrode
2
(Supplementary Figs. 32-35 and Supplementary Table 11). Overall, the performance of 0.5Li SO -
2 4
ZrCl -based cells is competitive with other representative halide SEs (Supplementary Table 10),
4
underscoring its potential for high-energy-density ASSLBs.

In conclusion, we present a dual-anion clustered amorphous SE, xLi SO -ZrCl (x = 0.5) that
2 4 4
combines a low lithium content of 2.4 wt% with good performance. The 0.5Li SO -ZrCl SE
2 4 4
exhibits a high ionic conductivity of 1.5 mS cm-1, at 30 ℃. This conductivity can be ascribed to
the non-periodic structure of the anion clusters introduced via ball milling. An ab initio workflow
integrating first-principles calculations and machine-learning molecular dynamics determined the
amorphous atomic structure without empirical parameters. The calculated pair distribution
function and temperature-dependent ionic conductivity show good agreement with experimental
data. Within this structure, the migration of lithium ions is facilitated by the dual-anion clustered
[Zr Cl (SO )]2- framework, where a reduced oxygen coordination is associated with enhanced
a 4a 4
ionic mobility. ASSLBs employing this SE with the NCM811 positive electrode demonstrate stable
long-term cycling performance, with 81.1% capacity retention after 1400 cycles at 1.0 C, and
achieve a high-loading performance with a stable capacity of 5.3 mAh cm-2 after 300 cycles at 0.1
C. Furthermore, 0.5Li SO -ZrCl also showcases improved air stability compared to that of 2LiCl-
2 4 4
ZrCl at 30% RH. This anion-clustering strategy addresses key challenges of cost and stability for
4
halide SEs, offering a viable design approach for practical high-energy-density batteries.

Methods
Synthesis of electrolytes
SEs with the general formula xLi SO -ZrCl (0 < x ≤ 1.0) were synthesized via a mechanochemical
2 4 4
ball-milling. Stoichiometric mixture of Li SO (Aladdin, 98.5%, titration) and ZrCl (Macklin,
2 4 4
99.5%) were used as the raw materials. For comparative studies, LiCl (Aladdin, 99.5%) replaced
Li SO to prepare 2xLiCl-ZrCl . The starting materials were mixed in an argon-filled glovebox
2 4 4
(H O < 0.01ppm, O < 0.01ppm). Then, each mixture (1 g) was processed in a zirconia jar with 30
2 2
g grinding media using a FRITSCH Pulverisette 7 planetary mill. The milling protocol consisted
of 500 rpm operation for 15 min, followed by 5 min pauses, with this cycle repeated for a total of
50 times. After ball-milling, the products were transferred into the glovebox for further use.

Cost analysis
The raw material costs for large-scale production (1000 kg) were estimated using the quantity-
price scaling model55. This model relates the unit price (P) and purchase (Q) quantity through:
𝑙𝑜𝑔 𝑃 = 𝑙𝑜𝑔 𝑎+𝑏×𝑙𝑜𝑔 𝑄 (1)
10 10 10
where a and b are constants for a given chemical. Laboratory-scale price data (P, Q) were obtained
from Aladdin (https://www.aladdin-e.com/) and extrapolated to industrial quantities to calculated
the scaled cost.

Structure characterization
Lab-based XRD was performed on a Rigaku Ultima IV diffractometer (Cu Kα radiation,  =
1.5406 Å). Samples were sealed in Kapton films to prevent air/moisture exposure during
measurement. The neutron total scattering data were acquired on the multi-physics instrument
using a high-intensity time-of-flight diffractometer at the China Neutron Spallation Source (CSNS).
The high-energy X-ray total scattering test was conducted at BL08, SPring-8 in Japan, utilizing a
16-inch two-dimensional detector (XRD1621) with a wavelength of 0.1087 Å. The simulated
nPDF and sPDF G(r) values from CIF files are calculated by PDF gui software56. Scanning electron
microscopy (SEM) images were acquired using a Hitachi S-4800 field emission SEM equipped
with energy dispersive spectroscopy (EDS). Microstructures were further observed via
transmission electron microscopy (TEM) (Talos F200X G2). X-ray photoelectron spectroscopy
(XPS) data were recorded using a Thermo Scientific K-Alpha Spectrometer with a monochromatic
Al Kα source. Raman spectra were obtained using a HORIBA Scientific LabRAM HR Raman
spectrometer system (532.4 nm laser). X-ray absorption spectroscopy (XAS) measurements were
performed at BL11B of the Shanghai Synchrotron Radiation Facility (SSRF). An atomic force
microscopy (AFM) system (Bruker Dimension Icon) housed in an argon-filled glovebox was used
to characterize the topography and modulus images of cold-pressed solid electrolyte (SE) pellets,
which were formed under a pressure of 2 tons. For force spectroscopy measurements, an AFM tip
(Bruker Corp.) with a spring constant (k) of 200 N·m-1 and a resonance frequency (f ) of 525 kHz
0
was utilized.

Conductivity measurements
Prior to conductivity measurements, ~100 mg of the SE powders was placed in a
polyetheretherketone (PEEK) mold (diameter 10 mm), and cold-pressed to form pellets under a
pressure of 3 tons for 1 min using a hydraulic press (Xinnuo Instruments, SYP-12B). The thickness
of the resulting pellets was controlled within the range of 0.06-0.07 cm. All pellet preparation steps
were conducted in an argon-filled glovebox with water and oxygen contents < 0.01 ppm. Ionic
conductivity of the SE pellets was evaluated via EIS using a symmetric cell configuration with
two stainless steel rods as blocking electrodes. EIS measurements were performed on a Bio-Logic
MTZ-35 analyzer over a frequency range of 7 MHz to 1Hz with alternating current (AC) amplitude
of 10 mV. A total of 10 data points were collected per frequency decade to ensure sufficient data
resolution. Electronic conductivity was assessed through a direct current (DC) polarization
measurement. In this measurement, a voltage of 1.0 V was applied to the same symmetric cells.
All these measurements temperature control was realized in an environment test chamber.

Computational details
Reactions between anion clusters. The AIMD unit reactions of ZrCl and Li SO were calculated
4 2 4
in ORCA57-59 using the r2SCAN-3c60 composite electronic-structure method (r2SCAN functional61
with triple-ζ Gaussian atomic orbital basis set plus refitted Grimme’s D4 dispersion62 and
geometrical counter-poise corrections (gCP)63 for London-dispersion and basis set superposition
error). Def2/J acceleration64 is enabled for the formation of the Coulomb matrix in the calculation.
To investigate the dynamic behavior of these reactions, AIMD simulations were conducted in the
NVT (constant particle number N, constant volume V and a temperature fluctuating around an
equilibrium value T) ensemble at 350 K with a time step of 1.5 fs for 3000 steps, resulting in the
generation of two structural cluster units.

Construction of training dataset: The training dataset was constructed via fourth-iterative
(detailed in Supplementary Fig. 20) sampling AIMD trajectories from supercells containing 8
cluster units ([Li Zr Cl SO ] ) using the CP2K software65. Simulations employed the r2SCAN
2 2 8 4 8
functional61 with optimized norm-conserving Goedecker-Teter-Hutter pseudopotentials66 and the
matched double-zeta polarized basis set67 within the framework of Gaussian Plane Waves (GPW).
A plane-wave cutoff of 500 Ry was applied and each Gaussian function was mapped onto a grid
with a cutoff of 60 Ry. Orbital transformation (OT) method68 was used to facilitate convergence.
The self-consistent field energy was converged to a criterion of 1×10-5 Hartree. NVT AIMD
simulations were performed from 300 K to 1200 K using the Bussi-Donadio-Parrinello (BDP)
thermostat with a time step of 3 fs. All AIMD trajectory snapshots were recorded every two steps.

Machine learning force-filed finetuning. MLFF was constructed in the MACE framework69. We
employed MACE-matpes-r2scan-omat-ft as base model and subsequently fine-tuned it for our
specific system (detailed in Supplementary Fig. 20). The training dataset were randomly
partitioned into 75% training, 5% validating and 20% testing sets. During fine-tuning, all model
parameters remained frozen, with loss weights set to 1 for energy, 10 for forces, and 100 for stress.
As shown in Supplementary Fig. 21, the root mean square error (RMSE) of the MLFF model for
energy, forces and stresses is 0.4 meV/atom, 20.8 meV/Å and 254.5 bar, respectively, with R-
squared values of 0.9838, 0.9990, and 0.9929.

Construction of amorphous structure: The multiscale modeling from the first-principles
calculations to MLFF-MD were adopted for construction of amorphous structure. The workflow
was shown in Supplementary Fig. 15. The initial amorphous structure (unrelaxed, [Li Zr Cl SO ] )
2 2 8 4 64
was formed by stacking cluster units generated through AIMD. Subsequently, MLFF-MD was
responsible for performing structural relaxation and amorphous reconstruction. MLFF-MD
operates at the interface between the MACE and the Atomic Simulation Environment (ASE)
software70.

The primary objective of the MLFF-MD simulations is to equilibrate the density of the system,
a critical step in obtaining an accurate amorphous structure. This was accomplished using the NPT
ensemble at room temperature, with gradual adjustments to the external pressure until the density
difference between the simulation and experimental values was less than 0.1 g/cm3. A secondary
goal was to equilibrate the structural configuration. To achieve this efficiently, we employed a
method that involves initial heating followed by annealing, which could significantly reduce
computational costs compared to traditional prolonged relaxation at room temperature. These
processes resulted in a fully ab-initio-derived amorphous structure. Subsequent analyses, including
structure, PDF and ionic conductivity, were conducted based on this amorphous structure. The
optimized computational structures and the atomic configurations in AIMD and MLFF-MD
simulations used in this study are available as Supplementary Data 1-3.

Ionic conductivity calculation: The ionic conductivity of the material was evaluated using
MLFF-MD through the MACE-ASE interface, employing a time step of 3 femtoseconds for a total
of 100,000 steps across a range of temperatures. Specifically, for room-temperature ionic
conductivity, we performed MD simulations for a duration of 3 nanoseconds under the same
timestep. Following the simulations, trajectory analysis (e.g. MSD) was conducted with the ASE,
Kinisi71 and Pymatgen72, 73 packages. The ionic conductivity () at a specific temperature (T) is
calculated from the self-diffusion coefficient (D) using the Einstein-Nernst relation  = Dnq2/k T,
B
where k is the Boltzmann constant, q is the charge of the diffusing ion, and n is the concentration
B
of diffusing ion. The three-dimensional self-diffusion coefficient D is defined as 𝐷 = 1 lim 𝑑𝑀𝑆𝐷(𝑡) ,
here 𝑀𝑆𝐷(𝑡) = 1 〈∑𝑁 [𝑟(𝑡 +𝑡)]−𝑟(𝑡 )2〉, N is the number of diffusing ion. 𝑟(𝑡 +𝑡) and 𝑟(𝑡 ) denotes
6t→∞ 𝑑𝑡 𝑁 𝑖=1 𝑖 0 𝑖 0 𝑖 0 𝑖 0
the atom position at time t + t and t , respectively. 〈 〉 denotes an average over time origins.
0 0

PDF calculations of amorphous structures: The neutron and X-ray PDFs were computed using
the DebyeCalculator package74, with an isotropic atomic displacement factor (Debye-Waller factor)
of 0.02. The final PDF was obtained by averaging the configurations from 100 structural snapshots
sampled from the MLFF-MD simulations performed at a temperature of 300 K.

Electrochemical window and interfacial stability calculations
The chemical stability between the SE and electrode materials was assessed using a
thermodynamic approximation method based on first-principles calculations14, 75. The phase
equilibria along the energy minimum path were determined by constructing and comparing the
energies of all relevant phases within the corresponding compositional space. For a given phase
with composition 𝐶 , its equilibrium phase composition Eunder the chemical potential 𝜇
𝑀
of
element M is given by the grand potential phase diagram, denoted as 𝐶 (𝐶, 𝜇 ) . The
𝑒𝑞 𝑀
decomposition reaction energy ∆𝐸 of the SE in a n open system at the lithium chemical potential
𝐷
𝜇 (referenced to Li metal) is calculated as:
𝐿𝑖
∆𝐸 𝑜𝑝𝑒𝑛(𝑆𝐸,𝜇 ) = E 𝐸 (𝐶 (𝐶,𝜇 ))−𝐸(𝑆𝐸)− ∆𝑛 ∙𝜇 (2)
𝐷 𝐿𝑖 𝑒𝑞 𝑒𝑞 𝐿𝑖 𝐿𝑖 𝐿𝑖
where 𝐸(𝑆𝐸) is the ground-state energy of the SE, and ∆𝑛 is the stoichiometric change in the
𝐿𝑖
number of Li atoms during the reaction.

The SE-electrode interface was treated as a pseudo-binary system76. The composition of the
interfacial region, 𝐶 (𝑐 ,𝑐 ), is expressed as:
𝑖𝑛𝑡𝑒𝑟𝑓𝑎𝑐𝑒 𝐴 𝐵
𝐶 (𝑐 ,𝑐 ) = 𝑥𝑐 +(1−𝑥)𝑐 (3)
𝑖𝑛𝑡𝑒𝑟𝑓𝑎𝑐𝑒 𝐴 𝐵 𝐴 𝐵
Where 𝑥 is the molar fraction of the SE, and 𝑐 and 𝑐 represent the specific compositions of
𝐴 𝐵
the SE and the electrode, respectively. The total energy of this pseudo-binary interface system is
approximated as a linear combination of the energies of its components:
∆𝐸 (𝑐 ,𝑐 ,𝑥) = 𝑥𝐸(𝑐 )+(1−𝑥)𝐸(𝑐 ) (4)
𝑖𝑛𝑡𝑒𝑟𝑓𝑎𝑐𝑒 𝐴 𝐵 𝐴 𝐵
Here, 𝐸(𝑐 ) and 𝐸(𝑐 ) are the ground state energy of the two components.
𝐴 𝐵
The mutual reaction energy ∆𝐸 (𝑐 ,𝑐 ,𝑥) between the SE and the electrode was
𝐷,𝑚𝑢𝑡𝑢𝑎𝑙 𝐴 𝐵
evaluated by constructing a pseudo-binary phase diagram. This involves identifying the
composition ratio (𝑥) that yields the most negative reaction energy:
∆𝐸 (𝑐 ,𝑐 ,𝑥) = 1 min [𝐸 (𝑥𝑐 +(1−𝑥)𝑐 )−𝑥𝐸 (𝑐 )−(1−𝑥)𝐸 (𝑐 )] (5)
𝐷,𝑚𝑢𝑡𝑢𝑎𝑙 𝐴 𝐵 𝑁 𝑥∈[0,1] 𝑒𝑞, 𝑖𝑛𝑡𝑒𝑟𝑓𝑎𝑐𝑒 𝐴 𝐵 𝐷 𝐴 𝐷 𝐵
In this equation, 𝐸 is the reaction energy of the equilibrium phase in the pseudo-
𝑒𝑞, 𝑖𝑛𝑡𝑒𝑟𝑓𝑎𝑐𝑒
binary system, 𝐸 (𝑐 ) and 𝐸 (𝑐 ) are the decomposition energies of the SE and the electrode,
𝐷 𝐴 𝐷 𝐵
respectively, and N is the total number of atoms involved in the phase equilibrium, used for
normalization.

Electrochemical measurements
Linear sweep voltammetry (LSV) test. The electrochemical stability of 0.5Li SO -ZrCl SE was
2 4 4
assessed by the cell configuration: Li|Li PS Cl|SE|SE+CNT. Carbon nanotube (CNT) with a purity
6 5
of 99.9% and a diameter of 20 nm was used as the conductive agent. Cell assembly was
performed in an argon-filled glovebox (H O, O < 0.01 ppm) as follow. Initially, 70 mg of SE
2 2
powder was placed in a PEEK mold. A pressure of 2.0 tons was then applied to press the powder.
Subsequently, the working electrode was prepared. For this, SE was manually mixed with CNT at
a weight ratio of 7:3. A total of 10 mg of the resulting mixture was uniformly spread on one side
of the SE pellet, and then pressed under a pressure of 3.0 tons. On the opposite side of the SE pellet,
30 mg of Li PS Cl powder was pressed at 2.0 tons. After that, a piece of Li foil (Tianjin Energy
6 5
Lithium Co. LTD, 99.95%, 6 mm diameter, 20 μm thickness). was attached to serve as the counter
electrode. LSV measurements were performed on a MTZ-35 analyzer (Bio-Logic) at a scan rate
of 0.1 mV s-1 and 25 ℃, scanning from open circuit to 6 V and 0 V, respectively. The Li PS Cl and
6 5
CNT were provided by Heifei Kejing Star Technology Company Co., Ltd.

Assembly of ASSLBs. The composite LiCoO positive electrode material was prepared by hand
2
mixing LiCoO with SE at a weight ratio of 70:30 for 10 min, without further additional processing.
2
For the preparation of the composite NCM811 positive electrode material, the following steps were
taken. Single-crystal NCM811, Vapor-Grown Carbon Fiber (VGCF) and 0.5Li SO -ZrCl were
2 4 4
mixed at a weight ratio of 70:30:3. Subsequently, the composite NCM811 positive electrode
material underwent ball-milling at a rotational speed of 300 rpm for 5 h. LiCoO , NCM811 and
2
VGCF were purchased from Guangdong Canrd New Energy Technology Co., Ltd.
The assembly process of ASSLBs initiated with the compression of 70 mg of SE powder in a
PEEK mold. A pressure of 2.0 tons was applied for 1 min. After that, approximately 30 mg of
Li PS Cl powder was uniformly dispersed on the other side of the SE layer, and then it was pressed
6 5
at 2.0 tons for 1 min. Next, the composite positive electrode was evenly spread on one side of the
cold-pressed SE pellet and then compressed under a pressure of 3.0 tons for 3 min. Finally, a piece
of indium foil (diameter 10 mm, thickness 0.1 mm) was placed on top of the Li PS Cl layer. This
6 5
was followed by the attachment of a piece of Li foil (diameter 6 mm, thickness 20 μm). Li-In alloy
served as the negative electrode, verified to maintain a stable potential of +0.62 V versus Li/Li+18.
The assembled cells were cycled using LAND CT2001A and Neware-CT-4008T battery testing
systems. The LiCoO -based cells were conducted at 25 ℃ and the NCM811-based cells were
2
measured at 30 ℃. The temperature control was achieved by placing the cells in an incubator
(Tianjin Hongnuo Instrument, SPX-250B, with a temperature accuracy of  1 ℃) during the
cycling process. All the assembly operations were carried out in an Ar-filled glovebox (H O, O <
2 2
0.01 ppm). To ensure the reproducibility of experimental results, at least 5 replicate cells were
tested for each electrochemical experiment. Specific current and specific capacity values reported
in this work are normalized to the mass of the active material in the positive electrode.

Air stability tests
Air stability evolution were systematically performed by exposing 500 mg powder samples of both
0.5Li SO -ZrCl and 2LiCl-ZrCl powder samples in open quartz crucibles to two controlled
2 4 4 4
environments: (1) a dry room at -40 ℃ dew point (~1% RH) for 6 h; (2) Ar atmosphere of 1% RH
for 24 h and humid air of 30% RH at 25℃ for 6 h, respectively. EIS measurements were performed
on the samples both before and after each exposure to characterize changes in ionic conductivity
at 30℃. The moisture absorption characteristics were evaluated by exposing 500 mg powder
samples in open Petri dishes (100 mm diameter) under 30% RH, with precise mass measurements
taken before and after exposure using an analytical balance (±0.1 mg resolution). In-situ X-ray
diffraction (XRD) analysis was carried out to monitor structural evolution: 100 mg unexposed
electrolyte pellets were subjected to 30% RH after ~2 min, and then XRD measurements were
performed at a scan rate of 20°/min to track changes in crystal structure during exposure.
The measurement of gas evolution was conducted by Differential Electrochemical Mass
Spectrometry (DEMS) using a QAS 100 mass spectrometer (Linglu Instruments, Shanghai). Three
hundred milligrams of the powder samples were weighed, and compressed into round tablets with
a diameter of 13 mm using a tablet press. The disc was then sealed in the battery mold. After that,
the carrier gas system was connected, and a leak check was performed. The gas line was purged
with a 2 mL/min Ar gas flow to expel all the air. The background argon gas flow rate was set to 2
mL/min, and the molecular masses of the gases to be measured were set as H S, HCl, SO , and
2 2
Cl . After the signal stabilizes, the formal collection of gas signals was initiated, with a collection
2
time of 30 min.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "We thank the authors for their detailed response to our review. After reading the revised manuscript, our original concerns\nhave been sufficiently addressed. The updated version provides clearer justification of the novelty and offers stronger\nsupport for their discussions and conclusions. We appreciate the additional effort taken to update figures, explain\nmethodology choices, and strengthen the experimental validation. Overall, we believe the manuscript has been significantly\nimproved and we recommend its publication in Nature Communication.", "reviewer_2": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature\nCommunications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career\nResearchers who co-review manuscripts.", "reviewer_3": "The authors have addressed my concerns. I have no further comments.", "reviewer_4": "The author provided lots of extra information in the revised manuscript, which is greatly appreciated. However, some\ncomments are not addressed well, and a few puzzles need to be addressed further.\n1. The response to comment 4 of reviewer 3 is not relevant. The comment is “Please clarify the designing principle involving\nSO42- instead of other polyanions.”, but the authors’ response is not relevant to other polyanions at all.\n2. The high energy of mechanochemical ball milling could plausibly induce the decomposition of the sulfate group, leading to\nthe evolution of SO2 gas and the formation of an oxychloride rather than the proposed polyanion-stabilized structure.\nTherefore, it is highly necessary to perform gas mass spectrometry analysis of the atmosphere inside the milling jar after ball\nmilling. In addition, it is strongly recommended that the authors conduct a quantitative elemental analysis on the final product\nto confirm that the molar ratio of Sulfur (S) to Zirconium (Zr) aligns with the stoichiometric input, which would directly refute\nthe possibility of sulfur loss and provide more robust support for the manuscript's central claim.\n3. The authors used LiCl–ZrCl4 as a reference for comparison with Li2SO4-ZrCl4. It should be noted that Li2SO4-ZrCl4\nexhibited an amorphous structure, whereas LiCl-ZrCl4 remained crystalline. Directly comparing their ionic conductivities is\ntherefore not entirely fair. According to most of the relevant literature, amorphous halide solid electrolytes typically exhibit\nmuch higher ionic conductivities than their crystalline counterparts (e.g., amorphous LiTaCl6: 10.95 mS/cm (Chem. Lett. 2023,\n52, 237–241) vs. crystallized LiTaCl6: ≈10 -5 mS/cm (Adv. Energy Mater. 2024, 14, 2402555); crystallized NaTaCl6:\n≈10 -3 mS/cm vs. amorphous NaTaCl6: ≈4 mS/cm (Matter 2024, 7, 1018–1034); crystallized NaNbCl4O: 0.1 mS/cm (Angew.\nChem. Int. Ed. 2024, e202416979) vs. amorphous NaNbCl4O: ≈1 mS/cm (Small 2025, 21, 2504677)). Therefore, it is\nrecommended to extend the ball-milling duration to synthesize amorphous LiCl-ZrCl4 and then perform a direct comparison\nwith Li₂SO4-ZrCl4 under similar structural states.", "reviewer_5": "The authors have provided extensive experimental data and analyses in response to the reviewers’ comments. While the\nresponses are extensive, the explanations related to moisture stability and the Li+ transport mechanism still leave certain\npoints insufficiently addressed.\n1. The authors experimentally demonstrate that 0.5Li₂SO₄-ZrCl₄ exhibits enhanced moisture stability through in situ XRD\nand DEMS. However, the computational results do not appear to sufficiently support this conclusion. (1) How the initial\nstructures were defined for calculating the reaction energies of 2LiCl-ZrCl4 and 0.5Li2SO4-ZrCl4 is not clearly explained\nand requires clarification. This is because the calculated DFT energies show large deviations depending on the structural\nconfiguration. It is also necessary to clarify whether the energies of various compounds were calculated using structures\ntaken from the MP Database. (2) Moreover, the reported positive hydrolysis reaction energies would generally indicate that\nreactions with H2O are unfavorable (ΔG > 0). An energetically non-spontaneous state implies that decomposition does not\noccur in either case. (3) In addition, since the differences in calculated ΔG values for ZrOCl2 and ZrO2 formation between\n0.5Li2SO4-ZrCl4 and 2LiCl-ZrCl4 are only ~0.03 eV, it is unclear whether such a small variation can be regarded as\nsignificant within the accuracy limits of DFT calculations.\n2. The authors suggest that Li⁺ ions with fewer initial oxygen coordinations are more likely to undergo long-range diffusion in\n0.5Li₂SO₄–ZrCl₄, and they suggest that the beneficial effect of amorphization outweighs the negative impact of oxygen.\nHowever, considering that the Li/O ratio scales directly with composition in xLi2SO4–ZrCl4, it remains unclear how a\nreduced Li concentration in a low-Li compound such as 0.5LiSO4–ZrCl4 could still promote ionic conduction despite the\ntrapping effect of oxygen and fewer carrier concentration. The explanation for why 0.5LiSO4–ZrCl4, in the presence of\noxygen potential wells, nonetheless exhibits enhanced Li⁺ transport therefore remains unconvincing. To rigorously verify the\nrole of oxygen, a direct comparative model (e.g., 0.5LiSO4–ZrCl4 vs. LiSO4–ZrCl4) should be considered. Furthermore, if\noxygen incorporation is intrinsically detrimental to conduction, one might expect that in oxygen-free systems such as 1LiCl–\nZrCl4, extended ball-milling could induce amorphization that achieves superior transport properties, highlighting the need for\na more systematic comparison with compositions like xLiCl–ZrCl with xLiSO4–ZrCl4 to substantiate the authors’ claim.\n3. It is impressive that the authors were able to investigate the detailed mechanisms of amorphous structures at a broader\nscale using MLFF, which would be inaccessible by AIMD. However, the explanation of these mechanisms fundamentally\nrests on the assumption that the MLFF model achieves very high accuracy. In this regard, the present work only discloses\nvalidation-set results for MLFF, and the study would become more complete if additional verifications of the MLFF were\nprovided. Specifically, in Supplementary Fig. 21 the accuracy of MLFF is evaluated using the validation set. Yet, since the\nvalidation set is already used during training for loss comparison, it is important to also report how well the model performs\non an independent test set. According to Liu et al. (npj Comput. Mater. 9, 174 (2023)), even if the MAE or RMSE of forces\nappear small, these averaged metrics often fail to capture the occasional large errors in specific trajectories during MD\nsimulations. Therefore, a more detailed comparison of errors in both the magnitude and direction of forces would be\nvaluable. Furthermore, it would be helpful to include additional validations such as the force softening scale recently\nproposed (Deng et al., npj Comput. Mater. 11, 9 (2025)), or out-of-distribution and phonon DOS tests as in recent MLFF\nstudies (Ding et al., Nat. Phys. 21, 118–125 (2025)). Incorporating such validations would significantly strengthen\nconfidence in the MLFF model and, consequently, the proposed mechanistic insights."}, "author_response": "Response to the Reviewer’s Comments (NCOMMS-25-21110A)\nDear Reviewers"}]

Decision letter