A highly utilized and practical lithium-sulfur positive electrode enabled in all-solid-state batteries

Domain: chemistry
Published: 2026-02-27
Editorial outcome: major_revision

Abstract
All-solid-state batteries using sulfur-based positive electrodes (cathodes) offer a cost-effective route to achieve high specific energy. However, low active material utilization and cycle life hinder performance. Here we demonstrate a positive electrode design that employs sulfide solid-state electrolytes, where a high energy synthesis approach forms a metastable and ionically conductive interphase on the active material surface. This interphase facilitates high active material utilization and contributes capacity with cycling. We also show that tailoring active material particle sizes to the micron-scale improves rate performance and cycling stability. Structural analysis reveals that the substantial volume change of sulfur-based positive electrodes during operation can partially offset that of the negative electrodes, thereby mitigating internal mechanical stress. The combined design principles enable sulfur areal capacities up to 11 mAh cm while maintaining stable cycling
-2
at 25C. We further demonstrate several specific-energy-focused cell architectures, particularly a
LiS anode-free pouch cell that operates under “low stack pressure” of 10 MPa. This work outlines
2
practical design strategies for constructing high-specific-energy all-solid-state batteries for a broad
range of emerging applications.

Introduction
In the last decade, the need for safe and cost-effective energy storage systems has grown
significantly. By 2030, the global demand for lithium-ion batteries is projected to double from 2.8
to 6 TWh1, exceeding the projected supply. Due to the increasing adoption of electric vehicles and
electrified aviation, much of this demand is driven by the transportation sector. While lithium-ion
batteries using insertion-type positive electrodes have made substantial progress in terms of cost
and energy density, these electrodes are reaching capacity and performance limitations,
necessitating the development of alternatives that are safer, lightweight, with lower cost to further
advance electrification technologies.2,3 All-solid-state batteries (ASSBs) using conversion-type
positive electrodes, such as lithium-sulfur (Li-S), can overcome the shortcomings of current
lithium-ion battery technology. Sulfur’s high specific capacity (1675 mAh g-1)4 and abundance5
make it a promising low-cost energy storage solution. ASSB architecture also improves
operational safety by using non-flammable solid-state electrolytes (SSE) and eliminates the
polysulfide dissolution and shuttling effect, a major challenge hindering the commercialization of
Li-S in liquid electrolytes6,7.

While ASSBs eliminate the polysulfide dissolution and shuttling effect, additional
interfacial challenges inherent of SSEs are introduced. Efforts have shifted to addressing the slow
kinetics and insulating properties of sulfur and Li S, in addition to their chemo-mechanical
2
degradation from expansion and contraction8. The conversion from S to Li S result in a 79%
8 2
volume change9 – ten times greater than that of conventional positive electrodes. This significant
volume change generates high internal pressure on the surrounding SSE matrix, leading to void
formation and poor interfacial contact10. Consequently, most Li-S solid-state studies have adapted
electrode fabrication methods from liquid systems. Typically, active materials are incorporated
within high surface area carbon hosts through ball-milling11–14, solution processes15–17, or vapor
deposition18,19 as a means to increase conductive interfaces and constrain volume changes.
Nevertheless, these methods have resulted in inconsistent utilization and cycle life. Given the large
amount of carbon typically used, inadequate ionic networks and insufficient contact to sustain
conversion are likely responsible. Strategies like heat treatments20 or creating 3D solid-electrolyte
structures21 demonstrate proof-of-concepts to enhance interfacial contact. Incorporating catalysts
has been found to overcome poor contact by improving conversion kinetics22,23, although using
critical elements like cobalt compromises the low-cost novelty of the Li-S system. While some of
these approaches have shown improvement compared to the liquid system, high areal loadings
with long cycle life necessary for practicality, have yet to be demonstrated.

Li-S conversion requires “triple-phase” contact between the active materials, ionic and
electronic network.8,24 This is easier to achieve in liquid systems, as liquid electrolyte can flow
through electrode pores. In ASSBs, intimate solid-solid contact can be limited and dependent on
electrode architecture, including optimal particle sizes and their distribution. Ideally, a uniform
distribution of active materials, SSE, and carbon should be achieved. Common synthesis
approaches to achieve this can occasionally promote decomposition or redox activity when using
sulfide-based SSEs16,25,26, formally associated with generating irreversible decomposition products
at high voltages27, especially with high surface area carbon28. However, the lower operating voltage
of sulfur conversion is more compatible with the lower oxidation stability of sulfide-based
electrolytes, where incomplete redox can be reversible29,30. Leveraging the redox activity of sulfide
SSE can enhance the reaction kinetics of sulfur and Li S31 and ultimately improve cycling
2
stability.

While SSE redox can potentially enhance reaction kinetics and cycle life, sustaining
interfacial contact between Li-S active materials and SSE remains a challenge. Interfacial reactions
between sulfur or Li S and sulfide SSEs can result in the formation of ionically conductive phases.
2
Previous studies by Lin et al., found that Li PS (LPS) and sulfur can react to form a sulfur rich
3 4
thiophosphate phase (Li PS )32, creating bridged sulfur chains no longer in cyclo-S form.
3 4+n 8
Similar strategies have also been carried out with Li PS Cl (LPSCl) and sulfur15 for its higher ionic
6 5
conductivity. However, previous synthesis routes employed solvents to facilitate these reactions,
which can introduce complexities to the fabrication process. Ideally, a solvent-free synthesis could
be employed. Tanibata et al., showed that sulfur can bond to P S units using a mechanochemical
2 5+n
milling process, which delivered a high initial discharge capacity.33 Leveraging both SSE redox
activity and forming ionically conductive interphases can potentially solve the utilization and
stability challenges still observed with Li-S positive electrodes. While mechanochemical synthesis
has been widely used, standards have yet to be established, likely due to the many synthesis
permutations that can be employed. Therefore, the impact of mechanochemical synthesis on Li-S
electrode architecture is still not fully understood.

Here, we present a Li-S positive electrode design that overcomes interfacial, kinetic, and
chemomechanical challenges when implemented in ASSBs (Fig.1). A high energy single step short
duration mechanochemical process creates ionically conductive interphases on the sulfur particle
surface through sulfur bonding between the solid electrolyte and sulfur particle. To enhance cycle
life, three features were enabled. First, activation of the sulfide SSE redox activity was
accomplished using the synthesis process and was confirmed to be reversible within the sulfur and
Li S voltage windows. Using differential capacity analysis and X-ray absorption spectroscopy
2
(XAS), we deconvolute these capacity contributions for both sulfur and Li
2
S positive electrodes.
Second, tailoring the active material particle sizes to the micron scale reduces ionic tortuosity,
enabling high-rate cycling. Third, structural analysis confirms that both sulfur and Li S positive
2
electrodes undergo significant volume change. We also reveal the mechanical disadvantage of Li
2
S
compared to sulfur due to its volume contraction behavior. Nevertheless, the volume change of
the positive electrodes was found to balance the volume change of negative electrodes, alleviating
internal pressures within the cell, especially when paired with high-capacity negative electrodes
such as silicon. As a result, stable cycling was achieved with various solid-state cell configurations
utilizing both sulfur and Li S positive electrodes. Sulfur positive electrodes at 10 mAh cm-2
2
delivered stable cycling performance with 87 % retention for over 140 cycles at 25C. A ‘proof of
concept’ Li S||anode-free solid-state pouch cell delivers a high reversible capacity of 900 mAh g-
2
1 under practical operating conditions, without catalysts or liquid electrolytes. This study provides
an in-depth investigation into Li-S positive electrode and cell level design, supporting the
development of practical and high specific energy ASSBs using Li-S conversion chemistry.

Results
Creating conductive interfaces. Positive electrode architecture including conductive interfaces
are one of the critical components to enable high performing Li-S ASSBs. The reactivity
between the active materials and the sulfide SSE was investigated using a one-step
mechanochemical milling process, where high milling intensities were used for a short duration.
Milling intensity influences the reaction kinetics by tuning the collision frequency between
particles. High intensities have been found to accelerate reaction products often leading to
shorter process times.34,35 These two parameters were found to be critical in creating conductive
interfaces between sulfur and the sulfide SSE without degrading ion conduction pathways within
the composite. To prove the effectiveness of this approach, two common fabrication methods
were re-visited: hand-mixing and a multi-step milling process where the sulfur, carbon, and SSE
are milled separately. Schematics illustrating each method and their expected distribution are
shown in Fig. S1. A commonly used sulfide electrolyte, carbon type, and weight composition
was employed for our study, where unmodified elemental sulfur, argyrodite SSE (LPSCl), and
acetylene black (AB) carbon were used to fabricate the positive electrode with a weight ratio of
30:50:20 wt.%. This composition was used in all electrochemical tests unless otherwise stated.
Prior to positive electrode composite synthesis, LPSCl was pre-milled to the micron scale to
induce more interfacial contact to sulfur (Fig.S2).

To compare the different positive electrode synthesis methods, cells were fabricated with
LPSCl as the separator layer and a Li In alloy as the negative electrode. Figure 2a shows their first
cycle voltage profiles where increases in utilization and capacity are attained by introducing high
energy milling steps to the fabrication process. The multi-step process however suffered from low
utilization, implying insufficient sulfur-SSE contact. It was also irreversible, marked by a low
conversion efficiency (CE) of 17%. Commonly observed in prior works13,18,25,36,37, a low CE
suggests insufficient mass transport to reconvert Li S back to sulfur, usually requiring high
2
activation potentials to do so38. The low CE can also be from the isolation of sulfur particles and
active surface areas after volume expansion. A comparison between hand-mixing and ball-milling
has been previously investigated by Ohno et al., where ball-milling achieved higher capacities,
resulting in a CE of 100%.39 However, the single-step method used here delivered a discharge
capacity near the theoretical (~1500 mAh g -1 at 25C) and a CE of 129%. This means additional
capacity from the SSE has been obtained. A high discharge capacity coupled with additional
capacity from the SSE, suggests that this fabrication strategy produced an architecture to facilitate
improved ionic/electronic transport and “triple phase” sites for conversion. Mechanochemical
synthesis has been a common method to incorporate solid state materials. Usually, this is done
using low milling intensities, long milling times, and in multiple steps to mitigate SSE
decomposition, recreated above. The one-step approach introduced here subjects all components
to the high energy milling process, likely altering the physiochemical properties of sulfur and
LPSCl.

The diffraction patterns of the positive electrode composites are shown in Figure 2b. With
all three methods, LPSCl and sulfur are detectable. Only the one-step method facilitated
amorphization. Since the diffuse scattering of AB carbon can mask diffraction peaks, the milling
procedure was done without carbon (i.e., S and LPSCl) (fig. S4a), where the full width at half
maximum (FWHM) increases for both characteristic peaks of sulfur and LPSCl. Amorphization,
evidenced by peak broadening can be explained from the high energy milling process, which
introduces heat, breaking down sulfur rings, and has been found to distort PS 3- units, creating
4
loose P-S bonds40. This is advantageous since crystalline sulfur (cyclo-S ) requires large activation
8
energies to break covalent bonds between sulfur atoms41, lowering the energy barrier for
conversion. Thermogravimetric analysis (TGA) was used to quantify elemental sulfur within the
positive electrode composite (fig. S5), which should sublime at 350C42. However, 6.5 wt.% of
sulfur is unaccounted for, indicating an alteration of the sulfur bonding environments and possible
reaction with LPSCl. To verify this, Raman spectroscopy was conducted focusing on the S-S
bending region centered at 152 cm -1, a signature peak of
2
symmetrical bending43. In this region,
a redshift is observed with the milled S and LPSCl and one-step milled composite which includes
carbon (Fig. 2c). The S-S bending at lower wavelengths suggests an increase in bond lengths from
P-S stretching, confirming the formation of sulfur rich thiophosphates (Li PS ), where elemental
3 4+n
sulfur bonds with the sulfur at the PS
4
3- terminals of LPSCl.15,32 This observation is complimented
by both a reduction in intensity and wavenumber for the symmetric P-S stretching of the
thiophosphate unit (PS 3-) in LPSCCl at 425 cm -1 (fig. S4b), suggesting polymerization with
4
bridging S-S bonds44. The bonding between the thiophosphate units of LPSCl and sulfur,
facilitated by the high energy synthesis used in this work, ensures intimate contact of the sulfur
and SSE interface, leading to high sulfur utilization shown in Fig. 2a. A high milling intensity is
required to induce the interfacial reaction between sulfur and LPSCl. Sulfur utilization was found
to decrease when using lower milling intensities (fig. S3), suggesting incomplete formation of the
Li PS phase.
3 4+n

The X-ray absorption near-edge structure (XANES) region is useful in examining
oxidation states and local bonding environments of sulfur. Synchrotron radiation at the Sulfur K-
edge was measured for the sulfur positive electrode and reference samples (Fig. 2d). Two main
peaks can be assigned to elemental sulfur at 2473.6 eV and LPSCl at 2472.4 eV. A pre-edge feature
at 2470.1 eV is also evident from 1st derivative of the spectra. Pre-edge features indicate a reduction
of the sulfur oxidation state and have been observed for long chain polysulfides (Li
2
S
y
)45. The pre-
edge observed here is likely from chains of sulfur in Li PS . In addition to milling intensity,
3 4+n
milling times are another important synthesis parameter. The ionic conductivity of composites
consisting of only sulfur and LPSCl under high milling intensities (500 rpm) and increasing milling
durations were measured, where an increase from 6×10-6 S cm -1 to 2×10-5 S cm -1 is observed after
1 hour (fig. S4c). The increase in ionic conductivity after 1 hour of milling supports the formation
of the ionically conductive sulfur rich phase. Standalone XRD, Raman, and EIS was measured for
the LPSCl catholyte at 1- and 10-hour milling intervals (Fig. S4d-f). Contrary to what was
observed for the S and LPSCl composites, continued milling of LPSCl alone decreases its ionic
conductivity from 1.3 ×10-3 S cm -1 to 4×10-5 S cm -1, attributed to the formation of insulative Li S
2
and LPS phases. These results highlight the importance of limiting milling durations, where high
milling intensities used for short durations were found to be sufficient to complete the interfacial
reaction between sulfur and LPSCl while preserving the ionic conductivity of the SSE and positive
electrode.

Low-dose cryo-TEM, high-angle annular dark-field scanning transmission electron
microscopy (HAADF-STEM) imaging, and elemental mapping was conducted on multiple
particles of the positive electrode composite (fig. S6 and S7). Hyperspectral mapping shows a
homogenous distribution of all components. Line scan results show only Cl signal on the particle
edges, attributed to LPSCl. At 50 nm from the particle edge, the Cl signal vanishes while only
sulfur is present, suggesting that the bulk of the particle is sulfur. As the scan moves toward the
center, the sulfur atomic fraction stabilizes to double the expected stoichiometry for LPSCl,
revealing the sulfur-rich phase on the surface (table S1). These results corroborate the interfacial
reaction between sulfur and LPSCl, formed on the particle surface by the high energy synthesis
process. This sulfur rich interphase lowers the energy barrier for lithiation, facilitating fast lithium
transport from the SSE matrix into the sulfur bulk as illustrated in Fig. 2e. Sulfur can also bond to
the thiophosphate units in Li PS (LPS). However, the lower ionic conductivity of LPS compared
3 4
to LPSCl (0.04 mS cm-1 vs. 2 mS cm-1) shown in Figure S8a, will likely result in lower sulfur
utilization when used as a catholyte. Cells were fabricated using LPS, where lower sulfur
utilization was obtained (Fig. S8b). Although, reasonable capacity and cell performance is still
achievable.

The interaction between LPSCl and Li S was also investigated. A comparison between
2
positive electrode preparation methods was conducted as done so for sulfur, where half cells were
fabricated with a Li In negative electrode. Comparable trends seen with sulfur are also observed
0.5
with Li S (Fig. S9). Hand-mixing the composite fails to cycle, and the multi-step process delivers
2
low utilization coupled with large polarization. The one-step method, however, delivers a high
specific capacity of 723 mAh g-1 with a high coulombic efficiency of 99.3%, indicating good
reversibility. With the Li
2
S cathode, the synthesis method resulted in the decomposition and
amorphization of LPSCl, evidenced by undetectable peaks in the diffraction pattern and a shift of
P-S stretching in PS
4
3- from 425 cm -1 to 418 cm -1, assigned to Li
3
PS
4
(LPS)46 (fig. S10a and b).
This indicates that Li
2
S reduces LPSCl to LPS. Despite the catholyte decomposition and reduction
in composite ionic conductivity (fig. S10c), the Li S cell exhibits stable cycling, where additional
2
plateaus present in the voltage profile are attributed to SSE redox after the 1st cycle (fig. S11). This
confirms the redox activity from LPSCl decomposition products. Usually formed
electrochemically, here the redox active products are formed during synthesis.

Evaluating LPSCl redox activity with the influence of sulfur. The redox activity of sulfide
SSEs has been responsible for delivering capacity beyond the theoretical when paired with Li-S
positive electrodes15,16,19 and associated with reversible29 and irreversible30 electrochemical
behavior depending on the voltage cutoff and formed decomposition products. The impact of this
behavior on Li-S ASSBs performance is still debated. Often overlooked in literature,
deconvoluting the capacity contribution between sulfur and the solid electrolyte is essential in
quantifying the effectiveness of the positive electrode architecture.

In this system, three active materials exist: bulk LPSCl, the formed interlayer phase
(Li PS ) consisting of sulfur that has reacted with LPSCl, and unreacted elemental sulfur. TGA
3 4+n
was used to quantify the reacted and unreacted sulfur mass, where 23.5 out of the 30 wt.% of
elemental sulfur remains unaltered after synthesis. Using the known sulfur and LPSCl mass, we
can estimate the capacity contributions of the active materials and compare to what was obtained
experimentally. For a 1.6 mAh cm-2 cell using a total composite mass of 2.5 mg, 23.5 wt.%
unreacted sulfur is expected to deliver 1 mAh (using a specific capacity of 1675 mAh g-1).
Therefore, during the discharge, unreacted bulk sulfur is estimated to contribute 83% of the
capacity with the remainder attributed to LPSCl and Li PS reduction.
3 4+n
To isolate the capacity
expected from just LPSCl, cells with just LPSCl and carbon were fabricated using a Li-In alloy
negative electrode as a lithium source. Constant current discharge and charge within the sulfur
voltage limits delivered a reduction and oxidation capacity of 115 mAh g-1 and 355 mAh g-1 (Fig.
3a). Therefore, a much higher charge capacity from LPSCl is expected. Subtracting the expected
LPSCl and unreacted bulk sulfur capacity from the total cell capacity obtained after the charge
leaves the remaining capacity to be assigned to reacted sulfur in Li PS . This equates to a
3 4+n
capacity contribution of 10% for the reacted sulfur and 27.5% for LPSCl after the charge,
suggesting that the reacted sulfur in Li PS and LPSCl are both redox-active and actively
3 4+n
participates during cycling. This coupled redox behavior is observed between the sulfur positive
electrode and LPSCl cells in the cyclic voltammetry (CV) scans (Fig. 3b).

The addition of sulfur will alter the redox behavior of LPSCl, especially after the high
energy synthesis process. Additionally, the capacity contributions of LPSCl and sulfur in Li PS
3 4+n
during the discharge remains unknown. Therefore, cells were fabricated with lower sulfur weight
percentages (10 wt.% and 20 wt.%) to highlight redox activity from the SSE and formed sulfur
rich interlayer phase. Their capacities were normalized by their total active mass (i.e., sulfur and
LPSCl) where differential capacity (dQ/dV) analysis was conducted to differentiate overlapping
redox potentials and capacity contributions (Fig. 3c).

Four main regions are labeled in the dQ/dV curves. Region I highlights the dominant
reaction during the discharge, attributed to the reduction of sulfur. Cells with higher sulfur wt.%
deliver most of their discharge capacity at this potential, meaning this is from unreacted bulk
sulfur. Region II highlights the beginning of SSE reduction. Starting at 1.3V, the reacted sulfur in
Li PS begins to reduce. This reduction peak is identifiable in the dQ/dV curves for all cells, but
3 4+n
is mostly prominent for the cell with the lowest sulfur wt.% due to a higher capacity contribution
from this phase. The Li
3
PS
4+n
phase is expected to follow similar redox as LPS, which reduces at
a higher potential than LPSCl (fig. S12). Therefore, we would expect the reduction of this phase
to dominate until 1.15V, where LPSCl begins to reduce. We can estimate their individual discharge
capacity contributions based on the reduction potentials obtained from dQ/dV. From OCV to 1.3V,
1 mAh is obtained, as estimated above. The remaining capacity from 1.3 to 1.15V can be attributed
to Li PS , which is responsible for 7.5% of the overall discharge capacity. The remaining
3 4+n
capacity obtained from 1.15 to the lower voltage cutoff of 1V is from LPSCl, which contributes
9.2%. The specific capacity of the reacted sulfur is estimated to be 553 mAh g-1 during the
discharge, likely existing in the Li PS phase after synthesis. These calculations are presented in
3 4+3
table S2. The additional sulfur in the PS 3- framework increases its functional capacity and reduces
4
the electrochemical stability of the SSE. Region III is attributed to Li S oxidation. Higher wt.% of
2
sulfur will contribute higher capacity at this potential since more Li S was formed during the
2
discharge and is now being oxidized. The cell with just LPSCl and carbon will contribute capacity
at this potential, attributed to Li S oxidation as investigated from previous work by Tan et al.26.
2
Region IV starts at 2.7V where SSE oxidation occurs. This additional oxidation capacity is
recoverable upon the subsequent discharge, supporting the formation of a reversible redox-active
LPS like phase. This is highlighted in the dQ/dV curves for the 3rd cycle (Fig. 3d), where reversible
redox occurs at similar potentials as sulfur conversion and extends the obtained capacity at the
nominal voltage. Despite a portion of sulfur reacting with LPSCl, the reacted sulfur is still utilized
and can be estimated. Positive electrodes with 30 wt.% sulfur is 84% utilized during the 1st
discharge. Sulfur utilization significantly increases for lower sulfur weight percentages, increasing
to above 90% for the positive electrode with 10 wt.% sulfur. These estimations are listed in table
S3.

From the CV scans and dQ/dV curves, LPSCl and sulfur positive electrodes exhibit
reversible electrochemical behavior. Reversible electrochemical behavior is also observed while
operating within the Li S voltage limits (Fig. S13). Assuming this behavior is consistent when
2
paired with Li-S positive electrodes, redox products of the LPSCl electrolyte can be predicted.
Complete oxidation of LPSCl forms sulfur, LiCl, and P S with reduction products being Li S,
2 5 2
LiCl, and Li P.26,29 In both redox pathways, LPS is formed intermediately. Using the reduction and
3
oxidation capacity of the LPSCl cells, the expected redox products can be hypothesized where
upon discharge and charge, a lithium rich Li PS and lithium deficient Li PS can be formed.
4 4 2.5 4
This is consistent with previous reports discussing the reversibility of the LPS phase30,47.
Li PS Cl + 1 Li+/e- ⟶ Li S + Li PS + LiCl
6 5 2 4 4
Q / Capacity ≈ 100 mAh g -1
red
Li S + Li PS + LiCl ⟶ Li PS + S + 3.5 Li+/e- + LiCl
2 4 4 2.5 4
Q / Capacity ≈ 350 mAh g -1
ox
Li S + Li PS  Li PS + S + 3.5 Li+/e-
2 4 4 2.5 4
Q
reversible
/ Capacity ≈ 350 mAh g -1

It is expected that the formation of LiCl will reduce the effective ionic conductivity of the
composite. However, its’ formation does not seem to impact redox reversibility within this voltage
window. However, the degree of LPSCl decomposition, like most metastable SSEs, is highly
dependent on carbon contact and surface area, where a mixture of metastable (Li PS Cl)
6x 5
and
stable phases likely exists in parallel29. LiCl interaction with other redox products formed within
the cathode may help retain ionic conductivity and ionic conduction pathways. The upper voltage
cutoff below 4 V limits the complete oxidation of LPSCl, preserving ionic conductivity with
continued cycles. Meanwhile, the high oxidative tendency of LPSCl at 2.3 V vs. Li/Li+ is also
effective at reducing the activation potential of Li S oxidation. In this work, the activation potential
2
was found to be 2.4 V, without requiring the use of catalysts or kinetic promoters beyond the SSE
itself.

To investigate the reversibility of the sulfur positive electrodes, redox products during the
1st formation cycle were probed. In-situ electrochemical impedance spectroscopy (EIS) was
performed on the sulfur positive electrode at each state of charge (Fig. 4a). After discharge, an
impedance growth of 85.2  is observed, attributed to charge transfer resistance from Li S
2
formation. After charging, the charge transfer resistance reduces and returns near the pristine state,
indicating good electrochemical reversibility (Fig. S15). Post-mortem analysis was conducted to
validate Li S formation and conversion, where diffraction peaks attributed to nanocrystalline Li S
2 2
are detected in the XRD spectra of the discharged sample (Fig. 4b), reinforcing the high utilization
of sulfur. We also observe that LPSCl regains crystallinity during electrochemical Li S formation
2
as compared to the pristine sulfur positive electrode. In the charged sample, Li S is undetectable,
2
indicating its complete oxidation, while the formation of sulfur at 10.5 2 is observed. The high
charge capacity yet low peak intensity and amorphous background indicate that the reformed sulfur
is amorphous as in the pristine state and suggests the sulfur rich interlayer is preserved with
cycling. This was confirmed by conducting Raman spectroscopy on the cycled sulfur positive
electrode, where a redshift to lower wavenumbers at the S-S bending region is observed (fig. S14),
as previously shown for the pristine composite. A new diffraction peak at 15 2 is also identified
after the 1st cycle. Raman spectra collected at a region within the SSE bulk shows a convoluted
PS 3- peak, attributed to a mixture of LPSCl and LPS (fig. S16a and b). This new phase after the
4
1st cycle can be attributed to LPS, validating active participation and reversibility of LPS, as
hypothesized from the electrochemical and dQ/dV results. Sulfur quantification using TGA also
demonstrates good conversion efficiency (Fig. 4c). The mass loss differences between the pristine
and charged composite originate from the sulfur rich interlayer, which increases the thermal
stability of the reacted sulfur. The preservation of ionically conductive interfaces along with SSE
redox behavior helps sustain sulfur utilization and cycle life.

Deconvoluting the redox products between LPSCl, sulfur, and Li S is challenging due to
2
their overlapping redox potentials. This can be challenging using XRD, which is unable to resolve
amorphous phases, a prominent feature after electrode synthesis. Therefore, XANES spectra at the
Sulfur K-edge was measured for Li-S positive electrodes at various states of charge. The reference
spectra are reported in Fig. S17. Using the reference spectra, linear combination fit (LCF) analysis
can be used to quantify species.48,49 The fitted spectra for each electrode system, namely LPSCl
(Fig. S18a-c), sulfur (Fig. S19a-d), and Li S (Fig. S20a-c) positive electrodes, are reported in the
2
supplementary information.

Figure 4d-f shows the weight distributions of the quantified redox products. For cells
consisting of only LPSCl and carbon, half of LPSCl was found to decompose to LPS during the
initial discharge, as hypothesized from the electrochemical results. After charging, 9.7 wt.% of the
composite is estimated to be sulfur with 22 wt.% of LPS retained after the 1st cycle. LPS formation
and retainment reinforces its role as a redox intermediate. The redox activity of LPSCl likely
improves Li S oxidation kinetics and contributes active mass from sulfur formation after charging.
2
The LCF results for the sulfur positive electrodes are shown in Fig. 4e. After the initial discharge,
20.2 wt.% of Li S is estimated. This supports the high conversion efficiency of the positive
2
electrode architecture. LPSCl also partially decomposes to LPS, as observed for the LPSCl system.
A higher amount of LPS is also present than expected but is likely from the Li PS phase from
3 4+n
reacted sulfur. It must be noted that XANES quantification of the entire electrode can be
challenging due to the shallow probe depth when using tender X-rays. Discussion regarding
limitations of XANES quantification is included in the Supplementary Information. After
charging, Li S is no longer detected, with 32 wt.% assigned to sulfur. The additional capacity
2
observed after the formation cycle of the sulfur system can be attributed to the formation of LPS
and sulfur from LPSCl oxidation. These results coupled with EIS, XRD, and TGA support the
reversibility of the Li-S system, attributed to the formation of ionically conductive interphases and
redox activity of the LPS phase from LPSCl decomposition. Similar conclusions can be drawn for
the Li S system. Complimenting findings revealed by XRD and Raman, half of LPSCl decomposes
2
into LPS after synthesis (Fig.4f). Similarly for the sulfur system, the amorphous LPS/LPSCl
mixture retains its ionic conductivity, supported by the Li S electrochemical performance. A
2
complete summary of the fitting results can be found in table S5.

Enhancing rate and cycling stability. The previous electrochemical performances were obtained
using as received sulfur and Li S. Their particle sizes were found to range drastically, with some
2
particles on the order of 100 microns for sulfur (fig. S23a-c) and 30 microns for Li S (Fig. S24a-
2
d). After composite synthesis, which subjects sulfur to high energy milling with LPSCl and carbon,
it is possible that particle sizes can reduce, although, sulfur particles were found to retain close to
their original size as observed from the SEM-EDS mapping results (fig. S25). These particle sizes,
while sufficient for low-rate cycling, need to be reduced to improve contact and Li+ transport for
higher rate operation. Sulfur particle sizes were produced on the micron and sub-micron scale (fig.
S26a-f). XRD was conducted after the size reduction process to ensure sulfur’s crystallinity was
preserved (fig. S27). Given the many possible compositions, geometrical modeling was conducted
to aid in experimental design, connecting sulfur particle size and mass fraction to active surface
area and transport properties. Sulfur positive electrode geometries were stochastically generated
with the sulfur as spherical particles and the carbon additives as aggregates (Fig. 5a). To compare
with experimental capabilities, particle sizes for bulk, micron, and sub-micron sulfur ranged from
25 to 50 µm, 0.5 to 5 µm, and 0.25 to 0.5 µm, respectively. Figure 5b shows the evolution of the
active surface area (the percentage of sulfur in contact with the SSE) as a function of sulfur content,
represented as the active material (AM) mass fraction within the positive electrode. The reported
values are the average over three repetitions.

Intuitively, the electrode with the smallest sulfur particles achieved the highest active
surface area for all AM compositions, followed by the micron sulfur, and bulk sulfur electrodes.
Therefore, sub-micron sulfur particles are expected to achieve the highest utilization due to more
SSE contact. The tortuosity of the LPSCl phase was explored with the reported value being the
average tortuosity in all directions. Surprisingly, the micron sulfur electrode exhibits the lowest
ionic transport tortuosity, with all cases obtaining similar results until higher AM mass fractions
(Fig. 5c). The geometrical modeling results suggest that sub-micron particles may be the best
choice to achieve high utilization, yet high active wt.% may be challenging to implement since the
ionic transport tortuosity increases drastically after 50 wt.%. Despite the expected high utilization
with sub-micron particles from higher surface areas, it must be noted that the geometrical modeling
does not consider the chemomechanical effects from lithiation.

To verify the modeling results, room temperature electrochemical performance was
performed evaluating each sulfur particle class in Li-In half cells. All particle sizes deliver
comparable discharge capacities during the 1st formation cycle at low rates of 80 mA g-1 (C/20)
(Fig. 5d). The sub-micron sulfur cell delivers the highest discharge capacity of 1694 mAh g-1,
beyond the theoretical. This was followed by micron and bulk sulfur which delivered 1615 mAh
g-1 and 1500 mAh g-1, respectively. This means that more LPSCl redox activity can be activated
with higher surface area particles. Higher charge capacities are also observed for all cells. The
additional charge capacity results in a higher discharge capacity from the 2nd formation cycle with
a reduction in cell polarization from 577 mV to 452 mV (fig. 28a). The reduced cell polarization
is from the LPSCl and Li PS redox products and potentially beneficial chemomechanical
3 4+n
behavior that improves contact during volume expansion. Rate capability was conducted up to 1C
(1.6 A g -1 or 1.6 mA cm-2 ) (fig. S28b). The current density at 1C is beyond the critical current for
the Li-In negative electrode (~1 mA cm -2)50, however, all three cells deliver reasonable utilization
at all rates, with bulk sulfur delivering the lowest due to kinetic limitations within the positive
electrode. Since LPSCl was found to contribute capacity, the specific capacity considering both
active masses (i.e., sulfur and LPSCl) show reasonable utilization (700 mAh g -1) for both
active
micron and sub-micron at these rates. This is critical for solid-state, as the catholyte is usually
considered inactive, responsible for “dead weight” within the cell. In this work, the catholyte
contributes electrochemically to the overall cell capacity.

Long term high-rate cycling was evaluated to investigate the particle size effect on cycle
life (Fig. 5e). The bulk sulfur composite exhibits fast decay, due to longer Li+ transport lengths
from large sulfur particles and poor contact observed after densification (fig. S23c). Despite
achieving the highest utilization, the sub-micron sulfur composite delivers 61% retention, with
micron being the most stable system, achieving 81% retention after 500 cycles. A 20% increase in
retention with the larger micron scale particles is a significant improvement. The capacity fade
observed with the cell using the sub-micron sulfur particles may be due to excessive decomposition
from higher surface area between the carbon and SSE. To test this, cells were constructed with the
micron sulfur composite using carbons that possess different specific surface areas and
morphologies (fig. S29a-e), each chosen to intentionally facilitate more SSE redox activity.
Increasing the carbon surface area increased utilization, mainly from the SSE, but showed no
impact on cycling stability. This means the decay observed with the sub-micron sulfur composite
is likely of (chemo)mechanical origin, causing degradation at the positive electrode and SSE
particle interfaces. This interfacial degradation should result in impedance growth. Therefore, EIS
of the micron and sub-micron sulfur cells was measured after 100, 300, and 500 cycles (Fig. 5f).
The intermediate frequency range was assigned to the interfaces within the positive electrode
(cathode electrolyte interface).51 The sub-micron sulfur cells were found to possess a higher
resistance after all cycles, attributed to more mechanical degradation with continued cycling.
However, these results differ marginally (Fig. S30a, b). Cross-section scanning electron
microscopy (SEM) imaging may be useful, although both electrodes were found to show
indistinguishable morphology (fig. S31a, b). Therefore, mechanical-based simulations using the
generated electrodes in Fig. 5a may shed insight on the accumulated stresses at these interfaces.
Sulfur undergoes large volume expansion upon lithiation52 resulting in internal stresses at
its interfaces. To simulate this mechanism, volume expansion was estimated based on % utilization
for each particle size using 1st cycle discharge capacities, subtracting the expected capacity from
LPSCl (fig. S32). Simulations were conducted using finite element method (FEM) where the sulfur
particles underwent the prescribed volume expansion derived from the electrochemical results.
Parameters and equations used in the simulations can be found in table S6 and S7. Figure 5g
displays the distribution of maximum von Mises stress on the SSE matrix, where for both micron
and sub-micron cases, stresses beyond 5 GPa were predicted. A bimodal stress distribution shown
in Fig. 5h shows that most of the SSE matrix does not experience stress since only 30 wt.% of
sulfur is used in the simulation. However, at the sulfur particle and SSE interface, the sub-micron
sulfur composite experiences a higher frequency of stress from more volume expansion and
increased tortuosity within the electrode. Stress accumulated at the boundaries of the sulfur
particles can propagate when these particles are in proximity, creating thinner SSE channels that
experience much higher stress versus the bulk. If SSE fracture occurs, pore formation is possible
after delithiation. These simulations only capture stress after the 1st lithiation. However, SSE
degradation is expected to accumulate with cycling, disrupting ion conduction pathways and
leading to more capacity fade as observed in Fig. 5e. Electrochemical evaluation, coupled with
FEM simulations uncover that sub-micron sulfur particles create a microstructure with high
electrode tortuosity, resulting in higher interfacial stresses and faster capacity decay. Nano-scale
sulfur and Li
2
S particles can be used to achieve high utilization53–55. However, in this work we find
that micron-scale particles balance both utilization and stable cycling, critical for practical Li-S
cathodes.

Morphological evolution of Li-S positive electrodes. In ASSBs, Li-S positive electrodes are
expected to experience chemomechanical degradation from their large volume changes, resulting
in stress on particle interfaces as simulated above. Electrochemical performances of sulfur, Li S,
2
and LPSCl were found to be highly reversible, but this impact on positive electrode morphology
remains unknown. To visualize this, cross-sectional SEM images were prepared for both micron
ARTICLE IN PRESS
sulfur and Li S cells at each state of charge (following the same protocol in Fig. 4). Both pristine
2
sulfur and Li S positive electrodes exhibit a dense structure, possessing a thickness of 26 µm and
2
45 µm, respectively, and good interfacial contact (Fig. 6a and 6b). Calculations were done to
estimate the volume change (%) as a function of sulfur wt.% assuming complete lithiation (fig.
S33). Here, 25.8 vol.% change is expected, resulting in a thickness increase to 32.7 µm, close to
the observed result. The lithiation capacity of this cell was 1.22 mAh. Therefore, the thickness
increase of the positive electrode equates to 4.9 µm mAh-1, which is also close to the expected
thickness growth of Li metal with cycling56. Therefore, lithiating the sulfur positive electrode can
compensate the volume reduction from stripping Li metal. After completing the 1st cycle, the sulfur
positive electrode thickness was retained but can be resolved by an increase in porosity. These
results reveal that during lithiation, the resulting stress on the SSE matrix causes it to plastically
deform to accommodate particle expansion. The deformed structure preserves intimate contact to
the separator layer and assists structurally during cycling. Subsequent volume expansion is likely
supported by the pre-deformed SSE structure. The plastic deformation of the SSE matrix is not
surprising, as the simulated stresses estimated above is near the shear modulus of LPSCl57.

Li S is expected to shrink after delithiation. Shown in Figure 6b, the Li S positive electrode
2 2
thickness decreases drastically (~ 40%), where columnar cracking is observed in the cross-section
as well as on the surface (fig. S34a, b). Surface cracking is likely strain induced from
inhomogeneous lithium removal. This morphological phenomenon has been observed in ASSBs
using silicon negative electrodes 58 and these results suggest that this chemomechanical behavior
occurs for conversion electrodes when constrained to 2D interfaces. After completing one cycle,
the electrode morphology is however, reversible, returning close to the pristine state and is a
similar thickness as the electrochemically formed Li S in Fig. 6a. These results provide two key
2
insights. First, conversion positive electrodes can alleviate internal pressure from negative
electrode volume changes. Second, Li S positive electrodes inherently face more mechanical
2
challenges compared to sulfur due to strain induced cracking after delithiation.

To demonstrate the pressure alleviation proof of concept, operando pressure monitoring
was conducted comparing LiCoO (LCO), sulfur, and Li S when paired with pure µSi and lithiated
2 2
Si negative electrodes. LCO will expand due to Jahn-Teller distortions59, although volume
expansion is 10 times less than Li-S positive electrodes. Si exhibits a high specific capacity and
undergoes near 200% volume expansion upon lithiation, however lower than expected due to the
constrained 2D interface in solid-state58. Nevertheless, pressure imbalance and resulting
chemomechanical degradation is one of the most challenging aspects for solid-state silicon
negative electrodes.60 In theory, a mechanically balanced system can be achieved using conversion
positive electrodes61. Since a lithium source is required for sulfur, cells were assembled using a
lithiated silicon negative electrode, where minimal pressure changes with sulfur are observed (Fig.
6c). However, the LCO cell pressure increases five times that observed for the sulfur case, with
pressure fluctuations consistent with previous studies62. The µSi||Li
2
S cell shows almost zero
pressure variation during the 1st formation cycle (Fig. 6d), partially responsible for the stable
cycling performance shown in Fig. S35, with 83% retention after 140 cycles. These results
demonstrate that pressure changes from high-capacity negative electrodes can be compensated
using conversion positive electrodes that are highly utilized. This improves cycle life and mitigates
cell ‘breathing’ during cycling, an important consideration for higher loading cells and when
integrating cells into pack level architecture.

High loading electrochemical evaluation. High loading electrodes are required to achieve high
cell specific energy. Therefore, cells with increasing sulfur loadings were constructed and first
paired with a LiIn alloy negative electrode (Fig. 7a). Areal capacities of up to 10 mAh cm -2 with
a discharge capacity of 1314 mAh g-1 was obtained (Fig. 7b), although with a slight increase in
polarization, attributed to the high sulfur loading of 7 mg cm -2 evaluated at 25C. Due to the
limited critical current density of LiIn alloys, cycling performance of this high loading system was
evaluated at relatively low current densities of 0.52 mA cm -2 (Fig. 7c), showing minimal decay,
with 86.8% retention after 140 cycles. The ability to achieve stable cycling at 10 mAh cm -2 at
room temperature (25C), demonstrates the effectiveness of the positive electrode microstructure.
To overcome the low current density of LiIn, free standing dry process positive electrode films
were fabricated and paired with a Li Si negative electrode (Fig. 7d). These cells were found to
2
deliver reasonable capacities up to 1C (5.5 mA cm -2) as shown in Fig. 7e, with full recovery at
C/20 (0.3 mA cm -2). This dry process Li Si||Sulfur system also demonstrates stable cycling
2
performance at reasonably high loadings 7.4 mAh cm -2, resulting in 77.4% retention after 150
cycles at 1.5 mAh cm-2 current densities (Fig. 7f).

Pouch cell form factors utilizing thick dry process positive electrodes, thin separator layers,
and high-capacity negative electrodes are necessary to achieve high specific energies in all-solid-
state.63 In addition, high areal capacity and active weight percentages are also required. Illustrated
in Fig. 7g, even lower weight percentages of sulfur at 10 mAh cm -2 can realize 500 Wh kg -1 and
is likely a more promising approach than increasing the weight percentage of sulfur. This is
because after lithiation, the volume of SSE compared to lithiated sulfur will reduce. Despite sulfur
possessing a higher specific capacity than its Li S counterpart, negative electrode selection is
2
limited to those with a lithium source, making lithium metal or pre-lithiated silicon a common
option, regardless of their high costs and manufacturing challenges. Using the positive electrode
composition reported in this work, Li S can achieve over 400 Wh kg-1 using silicon as the negative
2
electrode and go beyond 500 Wh kg-1 if combined with an anode-free architecture (Fig. 7h). To
demonstrate the capabilities of the positive electrode architecture using Li S as the active material,
2
a 15 mAh pouch cell was constructed without a negative electrode, utilizing dry process methods,
and reducing the separator layer thickness from 500 µm to 50 µm (Fig. 7i). Industry standard
formation rates and cycling protocols were used, where this configuration delivers high utilization
(1077 mAh g -1) and an initial coulombic efficiency of 83% (Fig. 7j). Stable cycling at C/3 (1.5
mA cm -2 or 0.34 A g-1) was also achieved at relatively low stack pressures of 10 MPa compared
to the pellet-type cells cycled at 75 MPa (Fig. 7k).

The electrochemical performances above showcase the versatility of the positive electrode
design methodology, where appropriate synthesis, SSE selection, and optimal electrode
microstructure are critical for Li-S chemistry to achieve high utilization and stable cycling.
Overall, this approach overcomes the interfacial, kinetic, and chemomechanical challenges
associated with Li-S positive electrodes in ASSBs. Our work provides a thorough electrochemical,
mechanical, and morphological analysis of the critical features required to enable high loading and
practical Li-S batteries.

Methods
Materials preparation and composite fabrication. All materials were dried under vacuum at
80ºC if not anhydrous and stored and prepared in an argon-filled glovebox (< 1ppm of H O and
2
O ). The solid-state electrolyte Li PS Cl (LPSCl, NEI Corp., USA) was used as the separator layer
2 6 5
and catholyte. When used as a catholyte, LPSCl was milled in an argon environment at 400 rpm
for 2 hours with 5 mm spherical yttria stabilized zirconia media using a high energy planetary ball
mill (Retsch, Germany). Similar procedures were done for -Li PS (LPS, NEI Corp, USA). For
3 4
the sulfur positive electrode, elemental sulfur (99.98%, Sigma Aldrich, USA) was either used as
received or milled at 400 rpm for 10 and 24 hours for micron and sub-micron particles with a
sample to milling media weight ratio of 1:10. For the Li S positive electrode, Li S was either used
2 2
as received (99.98%, Sigma Aldrich, USA) or milled following similar procedures as sulfur. For
electrochemical evaluation, optimal positive electrode composites were milled for 1 hour (unless
otherwise specified) at 500 rpm (unless otherwise stated) using a planetary ball mill with a sample
to mill media weight ratio of 1:30. Other composite trials were either first milled with carbon and
sulfur at 500 rpm for 1 hour, followed by hand mixing the SSE or hand mixing all components
using a mortar and pestle for 1 hour. The sulfur and Li S composites consist of 30 wt.% active
2
material, 50 wt.% LPSCl and 20 wt.% of a conductive agent. The LPSCl and LPS composites
consist of 80 wt.% SSE and 20 wt.% conductive agents. The primary conductive agent used in this
work was acetylene black (Sigma Aldrich, USA) for cell cycling and CV measurements. Other
conductive agents were also used such as vapor grown carbon fiber (Sigma Aldrich, USA) and
Ketjen black (EC-600JD, MSE Supplies, USA). For the LCO cells, niobium coated lithium cobalt
oxide (LCO, MSE Supplies, USA) was used as received. The LCO positive electrode composite
was prepared by hand-mixing with a mortar and pestle and fabricated with a composition of 70
wt.% LCO and 30 wt.% LPSCl. The Li-In negative electrodes were prepared by vortex mixing
stabilized Li metal powders (99.9%, FMC, USA) and indium metal powders (99.99%, Sigma
Aldrich, USA) with either a molar ratio of Li In or Li In for 5 min. The lithiated silicon negative
0.5 1
electrode with a molar ratio of Li Si was fabricated in a similar fashion, by vortex mixing Si
2
(99%, Sigma Aldrich, USA) and stabilized Li metal powders for 3 min, as described by Ham et al.
65 The silicon negative electrodes were fabricated by casting a 99.9 wt.% silicon slurry consisting
of Si particles, 99.5% purity N-methylpyrrolidone, and polyvinylidene fluoride (Sigma Aldrich,
USA) on a copper current collector using a doctor blade. The electrode films were dried under
vacuum at 80°C overnight, punched into 10 mm diameter discs and transferred into an argon-filled
glove box prior to use, as described by Tan et al. 58 All materials and electrodes were stored in an
argon-filled glovebox at a storage temperature of 22C  3C and used within 1-4 weeks of
preparation.

Dry process positive electrode fabrication. Positive electrode composite powders were mixed
with 1 wt.% PTFE (Chemours, USA) in a hot mortar and pestle (50-60°C) until a dough like
consistency is formed. The positive electrode composite was then hot rolled (MTI corp., USA)
under 60°C conditions with decreasing thickness until a 300 to 200-micron film was made,
correlating to an areal loading of approximately 4.5 – 6 mAh cm-2.

SSE film and anode-free layer fabrication. For preparing the materials used in the pouch cell,
the SSE film consisted of LPSCl (98 wt.%) and an acrylate binder (2 wt.%) that were mixed in p-
xylene (Sigma Aldrich, USA). The resulting mixture was casted on a polyethylene terephthalate
film and dried under vacuum at 40C overnight. The slurry for anode-free layer was prepared by
mixing carbon black (Imerys, USA), silver nanoparticles, and polyvinylidene fluoride (Solvay,
Belgium) in N-methylpyrrolidone (Sigma Aldrich, USA) at a weight ratio of 69.75:23.25:7.0 as
described by Lee et al.64 The slurry was coated onto a 10 μm thick stainless-steel foil using a doctor
blade and subsequently dried in a vacuum oven at 100°C overnight.

Cell fabrication and electrochemical measurements. Electrochemical measurements were
performed in custom 10 mm diameter pellet cells constructed out of Grade 5 titanium plungers and
polyether ether ketone (PEEK) dies. For electrochemical measurements in pellet cells, the
separator layers were first pressed at 3 tons (375 MPa) for 3 min, resulting in a thickness of 450 –
500 m. Positive electrodes and Si negative electrodes were pressed to 3 tons for 5 min, while
the LiIn and Li Si negative electrode was pressed to 1 ton (125 MPa) for 30 seconds. After
2
assembly, the pellet cells were inserted into custom cell holders and hand tightened to 75 MPa.
For each electrochemical measurement, three cells were fabricated and tested. EIS and CV
measurements were collected using a Biologic SP-300 potentiostat directly after cell assembly.
For EIS measurements a voltage amplitude of 30 mV and frequency range of 7 MHz to 100 mHz
was used, with 10 points per decade. Ionic conductivity for the solid-state electrolyte and
composites was calculated using the following equation 𝜎 = 𝐿/𝑅𝐴, where L is the pellet thickness,
R is the resistance, and A is the area of the pellet cell (0.785 cm2 for 10 mm cell diameter). The
cell configuration used to assess ionic conductivity was Ti|SSE|Ti or Ti|composite|Ti. Equivalent
circuit modeling was performed using ZView. For CV measurements, a sweeping rate of 0.1 mV/s
was used within the operating voltage range of sulfur (1-3 V vs. Li/Li+). For all electrochemical
performance evaluation, pellet cells were cycled under 75 MPa (unless otherwise stated).
Galvanostatic cycling measurements were evaluated using Neware Instrument cyclers (CT-
4008T). Constant current was applied to cut-off potentials of 1-3 V vs. Li/Li+ for Li In||Sulfur
1
cells, 0.7 – 3V vs. Li/Li+ Li Si||Sulfur cells, 3.4V – 1V vs. Li/Li+ for Si||Li S cells, and 3.4 – 1V
2 2
vs. Li/Li+ for anode-free||Li
2
S pouch cell. All cycling measurements were conducted in an inert
glovebox environment at 25C  1C unless otherwise stated. The anode-free||Li
2
S pouch cell was
cycled in a convection environmental chamber at 60C. The designed cell areal capacity was
determined by the active mass of the positive electrode multiplied by the corresponding specific
capacity (1600 mAh g-1 for sulfur and 1100 mAh g-1 for Li S) divided by the cell area (0.785 cm2
2
for pellet cells and 3.24 cm2 for pouch cell). For cells labeled with their areal capacity, areal
capacity was determined by dividing the charge capacity for Li S and average discharge capacity
2
for sulfur, by the electrode area. The reported specific currents are the applied currents divided by
the active mass of either sulfur or Li S. In some cases, the experimentally obtained specific
2
capacity is reported by dividing the cell capacity by the total active mass, considering both AM
and the SSE within the positive electrode. For the sulfur system, the coulombic efficiency is
calculated by dividing the charge capacity by the discharge capacity. For the Li S system,
2
coulombic efficiency is reported as the discharge capacity divided by the charge capacity.

Pouch cell fabrication. Al foil, dry processed Li S positive electrode, SSE film and anode-free
2
layer, were stacked, and packed into a pouch. The pouch was vacuum-sealed and pressed to 500
MPa at 80C using warm isostatic pressure (WIP).

Scanning electron microscopy. Scanning electron microscopy (SEM) was performed on a FEI
Apreo and/or FEI Scios DualBeam focused ion beam (FIB)/SEM with 5 kV accelerating voltage
and 0.1 nA beam current for powders and pellets. Powders, pellets, and cycled electrode samples
were prepared in an argon-filled glovebox. Pellets and cycled electrodes were carefully extracted
from plunger cells using a polished Ti plunger and mounted to an SEM stub using carbon tape. All
SEM samples were transferred using an air-tight transfer arm to avoid any air exposure. For FIB
cross-sectional images, milling was done under cryogenic conditions (-180C) where Ga was used
as an ion beam source. Parameters used for all milling conditions of 30 kV, 65 nA, with the
subsequent cross-section cleaning performed with 30 kV, 15 to 7 nA, if necessary.

X-ray Diffraction. XRD measurements were collected over a 5 – 50 2 range on a Bruker
ApexII-Ultra CCD microfocus Rotating Anode instrument with Mo K ( = 0.7107 Å) radiation

at the UCSD X-Ray Crystallography Department. Samples were prepared in an argon-filled
glovebox using 0.7 mm boron capillaries, and flame sealed to ensure air-tight measurements of
sensitive samples. For cycled electrodes, electrodes were extracted using a scalpel and lightly
ground to powder using a mortar and pestle for capillary preparation.

Raman Spectroscopy. Raman spectra were collected using a Renishaw inVia upright microscope
using 785 nm wavelength laser source. Samples were prepared in an argon-filled glovebox with
powders or extracted electrodes placed on 25x25mm glass slides and sealed with Kapton tape for
airtightness.

Thermogravimetric analysis. TGA measurements were conducted using a NETZSCH STA 449
F3 Jupiter Simultaneous Thermal Analyzer. Samples were prepared and loaded in aluminum pans
and crimped inside the glovebox for airtight analysis. Aluminum pans were sealed in 20 ml glass
vials for transport. During the measurements, nitrogen gas was used to reduce any air exposure.
All samples were measured from room temperature to 450C using a scan rate of 5C/min. Vendor
specified resolution is 0.1g.

Transmission electron microscopy. The composite powders were mounted to an airtight cooling
holder from Melbuild in an argon-filled glovebox to eliminate any contaminations to the sulfur
samples and transferred to the TEM column directly without any air or moisture exposure. The
sample was cooled down to cryogenic conditions (~180℃) and stabilized for additional 30 minutes
before electron beam exposure. (S)TEM results were obtained on ThermoFisher Talos X200
equipped with a Ceta camera operated at 200 kV with low dose capability. The energy dispersive
X-ray spectroscopy (EDS) characterization is installed with compositional mapping using 4 in-
column SDD Super-X windowless detectors. The data acquisition was operated at low dose
condition to minimize any beam damage to the sample.

X-ray Absorption Spectroscopy. Tender X-ray absorption spectroscopy (XAS) measurements
were conducted at Taiwan Light Source (TLS) beamline 16A1 of the National Synchrotron
Radiation Research Center (NSRRC) in Hsinchu, Taiwan. The beamline uses a double-crystal Si
(111) monochromator for the photon energy range from 2 to 8 keV. Pristine and cycled electrodes
were extracted in pellet form. All samples were sealed in a pouch made of 2.5 um thick Mylar®
film inside an argon-filled glovebox to prevent air exposure. Each sample was mounted onto the
holder and placed in the measuring chamber at an angle of 45 to the incident X-ray beam. The
chamber is constantly purged with He to reduce the X-ray attenuation for at least 45 minutes before
collecting the XAS data. The Sulfur K-edge X-ray absorption near-edge structure (XANES)
spectra were collected in the total fluorescence yield (TFY) mode using a Lytle detector with a
scanning step of 0.2 eV. The photon energy was calibrated to 2472 eV (maximum in the 1st
derivative) at the S K-edge using elemental sulfur. The XANES spectra background subtraction,
normalization, and the Linear Combination Fit (LCF) were performed on Athena software. 66

Modeling of sulfur electrode geometries. The electrodes structures were stochastically
generated using the MATLAB codes from Duquesnoy et al.67, with the S as spherical particles,
and the carbon additives as aggregates. A volume fraction of 10% was dedicated to pores, the S
amount ranged from 30 to 60 %, and the volume ratio between the LPSCl and carbon additives
was kept constant at 5:2. Three different cases were investigated, Bulk, Micro and Nano, with S
radii ranging respectively from 25 to 50 µm, 0.5 to 5 µm, and 0.25 to 0.5 µm. To have a
representative volume for each condition, the length of the cubic electrodes was 200 µm for the
Bulk, 50 µm for the Micro, and 15 µm for the Nano. For each set of S size and amount, 3
electrodes were generated to obtain statistically relevant observables. The evolution of the active
surface area was monitored as the specific surface area, i.e. the ratio between the number of
pixels of S in contact with LPSCl and the total number of S pixels. The tortuosity of the LPSCl
phase was investigated using TauFactor68 in MATLAB, and the reported value is the average
value of the tortuosity of the electrolyte phase in all directions.

Finite Elements Method (FEM) Simulations. The electrodes were meshed using the open-access
toolbox Iso2Mesh69 and later imported into COMSOL Multiphysics 6.1. There, using the Solids
Mechanics module, the set of parameters and equations in Table S3 and Table S4 were set. In the
model, the electrode was assumed to be fully compact (no porosity) and the S particles were
uniformly lithiated throughout the simulation, leading to a volume expansion made possible with
the “Hygroscopic Swelling” node which normally accounts for the volume expansion of solids due
to the amount of water. During the simulation, the external boundaries of the electrode were fixed.
To determine the analog hygroscopic coefficient of each type of S particles, a 2-D simulation
consisting of the exact same model for a single SP particle was performed. The hygroscopic
coefficient was deemed adequate when the S particle would reach the desired volumetric
expansion (controlled here by its radius) at full lithiation. The cases of sub-micro and micron S
were investigated through FEM simulations for an AM content of 30 wt.% where three electrodes
were used for each case, and the value reported in the manuscript are averaged over all three
electrodes.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_4": "The extensive data presented in this manuscript reflect meaningful technical progress in the development of all-solid-state lithium-sulfur batteries (ASSLSBs). However, given the number and complexity of the experiments, several aspects of the experimental design lack sufficient detail, and further clarification is necessary. Specific concerns are outlined below.\n1. Influence of mechanical milling on SSE properties and S-SSE Interface. Mechanical milling appears to play a critical role in altering the physicochemical properties of the SSE and the formation of the S-SSE interface. Milling intensity determines the amount of energy imparted to the material, and while hand mixing can be regarded as a low-energy baseline, directly comparing it with high-energy milling at 500 rpm offers only a limited perspective. Please include a discussion, supported by appropriate electrochemical tests and literature references, on how varying milling intensities affect interface formation and electrochemical performance.\n2. Electrochemical redox mechanism during cycling of 1-step milled ASSLSBs. What are the dominant redox reactions occurring at each potential during cycling in the 1-step milled ASSLSB? Furthermore, does Li3PS4+n contribute to the overall capacity? Since the reversible redox reaction of the SSE appears to occur from the second cycle (Fig. 3a), the voltage profile and dQ/dV analysis beyond the second cycle should be done to clarify the electrochemical behavior of the proposed ASSLSBs. It seems that the SSE’s redox activity occurs at higher potentials than that of sulfur. (Fig. 3a and 3b)\n3. Stability of Li3PS4+n during cycling. Beyond confirming the formation and functionality of Li3PS4+n in the initial stage, how does this material evolve during cycling? Is it retained or transformed in later cycles? Please clarify and provide evidence for these changes.\n4. Stability and structural evolution of milled sulfur. It is stated that micron-sized and sub-micron sulfur were obtained via 400 rpm milling, yet the final cathode composite is fabricated through 500 rpm milling. Is it reasonable to assume that the particle size from the pre-milling step is retained after subsequent composite fabrication? Also, does milling affect only the particle size, or are there other changes such as in crystallinity or bonding structure of sulfur? These aspects are critical to cathode design, and a more thorough explanation is required to convincingly support the role of milled sulfur.\n5. Clarification of cell compositions used in electrochemical tests. Due to the large number of cathode composite configurations and corresponding electrochemical data, it is difficult to follow the overall flow. Please clearly summarize and specify the cell configurations used in each data panel to improve readability and interpretation.", "reviewer_5": "In this work, the authors report the fabrication of high-performance all-solid-state sulfur/Li2S cathodes using a one-step high-energy ball milling method. This process induces the formation of sulfur-rich thiophosphate phases (Li3PS4+n), which facilitate enhanced interfacial ion transport. The study presents a comprehensive investigation into the structural evolution of cathode materials during cycling, the redox activity of cathode components, the influence of electrode microstructure (particularly sulfur particle size) on battery performance, volume-pressure evolution during electrochemical cycling, and the overall electrochemical performance of the all-solid-state batteries. The performance demonstrated in both half-cells and practical full pouch cells is promising and offers insights into the development of practical solid-state lithium-sulfur batteries. However, several critical issues need to be addressed before publication:\n1. The areal sulfur loading should be clearly specified for all electrochemical data, particularly in Fig. 2a, Fig. 3, and Fig. 5e.\n2. For improved clarity, the authors should provide the calculated ionic conductivity values of the cathode composites and Li6PS5Cl (LPSCl) prepared under different milling conditions, in addition to the EIS spectra shown in Fig. S3c,f.\n3. As X-ray diffraction (XRD) patterns only indicate the crystal structure of the materials, they are unable to identify the amorphous sulfur-rich interlayer (Li3PS4+n) after cycling (line 331). Instead, Raman spectroscopy should be conducted to validate the existence of the interlayer phase after cycling.\n4. Although sulfur with different particle sizes was used, the high-energy ball milling process shall further reduce their particle sizes. Therefore, SEM-EDS mapping of the pristine cathode surface is recommended to examine the actual sulfur particle size after fabrication.\n5. It is unclear in Fig. 6c and d which curves correspond to pressure changes, and which represent voltage during cycling. The figure should be redrawn to improve clarity. Additionally, due to the relatively small pressure changes in all-solid-state lithium-sulfur batteries, the authors are encouraged to reduce the y-axis range for pressure to better illustrate the pressure evolution trends during cycling.\n6. It is recommended to use the same color (yellow) for sulfur particles in fig. 7d as in Fig. 7a to maintain consistency.\n7. The current densities are reported in C-rate, but it is unclear how 1C is defined. Given the varying cathode compositions (e.g., sulfur, Li2S, and LPSCl), all current densities are encouraged to be reported in A g-1 based on the total weight of sulfur/Li2S and LPSCl.\n8. The authors claim the formation of both Li3PS4 and Li4P2S6 based on a single XRD crystalline peak (Fig. 4b). This is not sufficient evidence. Additional techniques such as solid-state NMR or Raman spectroscopy should be employed to verify the presence of Li4P2S6. Moreover, prior literature reports Li4P2S7 rather than Li4P2S6 as an oxidation product of LPSCl (Nature Materials 2020, 19, 428–435.) and this discrepancy should be addressed\n9. The error bars in the figures should be clearly defined. The number of replicates used to generate the error bars should be specified in the respective figure captions.\n10. The figure caption for Fig. 5f should explicitly state whether the EIS spectra were collected at the charged or discharged state.\n11. On line 354, what is L4.1PS4 phase? Is it Li4PS4 or Li4.1PS4 phase?\nRegarding the questions raised by previous Reviewer #2, the authors have addressed some of them but the following still needs further elucidation:\n1. Comment 3 and 8 by Reviewer #2\nThe authors claim the formation of Li4P2S6, but as previously noted, this has not been conclusively demonstrated. Additional characterization is needed to confirm its presence.\n2. Comment 6 by Reviewer #2 - Redox Reactions\nIn the response, the authors state that the cathode contains 43.5 wt% LPSCl, 6.5 wt% Li3PS4+n, and 23.5 wt% sulfur, and that the capacity contribution from Li3PS4+n is marginal.\nHowever, this analysis is flawed. First, if 6.5 wt% of sulfur has reacted, the corresponding Li3PS4+n must be more than 6.5 wt% due to the added mass of the thiophosphate framework. Second, the chain length of sulfur in Li3PS4+n is unknown, making accurate quantification via TGA alone infeasible. Third, the theoretical capacity of Li3PS4+n (depending on the sulfur chain length) can exceed 500 mAh g-1, making its capacity contribution significant (see Angew. Chem. Int. Ed., 2013, 52, 7608–7611).\n3. Comment 9 by Reviewer #2 - Discrepancy Between XANES and XRD Results\nThe quantification of the LPSCl using XANES is intriguing. However, as stated by reviewer #2 the XANES data in Fig. 4e suggest complete decomposition of LPSCl after discharge, whereas XRD (Fig. 4b) still shows strong LPSCl peaks. The authors attribute this to differences in sulfur loading and the shallow probe depth of XAS but provide insufficient supporting evidence. Given the high LPSCl content, full reduction to LPS is unlikely despite the low areal sulfur loading. To support this claim, the authors should include the corresponding electrochemical data (e.g., voltage profiles, dQ/dV curves) for the cells used in the XRD and XAS experiments in the Supplementary Information. Testing conditions and cell compositions should be clearly specified, and capacity contributions quantified. Additionally, the manuscript should include a concise explanation, with a more detailed discussion in the Supplementary Information, regarding the accuracy and limitations of the XANES quantification.\n4. Comment 11 by Reviewer #2 - XPS\nXPS should be removed from the Author Contribution section.", "reviewer_6": "I think the authors have adequately addressed the concerns from the previous reviewers for publication in this journal."}, "author_response": "Note: Reviewer’s comments are shown in black, and our responses to the comments are shown in blue.\nReviewer #4\nThe extensive data presented in this manuscript reflect meaningful technical progress in the development of all-solid-state lithium-sulfur batteries (ASSLSBs). However, given the number and complexity of the experiments, several aspects of the experimental design lack sufficient detail, and further clarification is necessary. Specific concerns are outlined below.\nComment 1: Influence of mechanical milling on SSE properties and S-SSE Interface.\nMechanical milling appears to play a critical role in altering the physicochemical properties of the SSE and the formation of the S-SSE interface. Milling intensity determines the amount of energy imparted to the material, and while hand mixing can be regarded as a low-energy baseline, directly comparing it with high-energy milling at 500 rpm offers only a limited perspective. Please include a discussion, supported by appropriate electrochemical tests and literature references, on how varying milling intensities affect interface formation and electrochemical performance.\nResponse: We thank the reviewer for the constructive feedback and opportunity to improve our synthesis investigation. We have increased the synthesis scope to include different milling intensities to demonstrate its effect on the S-SSE interface and electrochemical performance. With 500 rpm being the maximum for our planetary ball mill, this limits our upper milling intensity parameter. Additionally, with 1 hour being our ideal milling duration, it is not advantageous to increase milling times as this requires more energy and has already been demonstrated to cause LPSCl to degrade. Therefore, we have extended the investigation to include milling intensities of 300 rpm and 400 rpm. This is done to investigate the milling intensity dependance on S-SSE interface formation and sulfur utilization. Cells were fabricated in an identical fashion to the 500rpm case, using a Li-In anode. Their electrochemical performance, compared to our ideal synthesis protocol, is shown below.\nThe 1st formation cycle is normalized by the sulfur mass, highlighting the clear discrepancies in utilization when using the different milling intensities. The composites fabricated at 300 and 400 rpm deliver a specific capacity of 565 mAh g-1 and 900 mAh g-1, respectively. Reducing the milling intensity was found to directly impact sulfur utilization, suggesting that the lower milling intensity is likely insufficient to fully induce the reaction between sulfur and LPSCl, leading to a lower wt.% of reacted sulfur. The reacted sulfur in the Li PS phase was found to be ionically 3 4+n conductive. A higher percentage of unreacted bulk sulfur limits the ionic conductivity of the composite, evidenced by the increased cell polarization.\nIn all composites, the conversion efficiency was over 100%, indicating that LPSCl redox activity has been activated. The 2nd cycle is also shown below. Voltage profiles were normalized by the active mass (sulfur and LPSCl) to evaluate the impact milling intensity has on the capacity contribution of LPSCl. All composites also deliver a higher 2nd discharge capacity than the initial, indicating that LPSCl redox is reversible despite its lower capacity contribution for the 300 and 400 rpm cases.\nIn summary, high milling intensities are required to facilitate the interfacial reaction between sulfur and LPSCl. Lower rpms can be employed to activate LPSCl redox activity and enable reversible cycling at a consequence of lower sulfur utilization due to the incomplete formation of the S-SSE interface.\nAdditional discussion has been added to the manuscript with two new references.\n“Cathode composite architecture including conductive interfaces are one of the critical components to enable high performing Li-S ASSBs. The reactivity between the active materials and the sulfide SSE was investigated using a one-step mechanochemical milling procedure, where high milling intensities were used for a short duration. The milling intensity is a complex parameter that influences the reaction kinetics by tuning the collision frequency between particles. High intensities have been found to accelerate reaction products often leading to shorter process times.34,35 These two parameters were found to be critical in creating conductive interfaces between sulfur and the sulfide SSE without degrading ion conduction pathways within the composite.”\n34. C. Burmeister, L. Titscher, S. Breitung-Faes, A. Kwade, Dry grinding in planetary ball mills: Evaluation of a stressing model. Advanced Powder Technol. 29, 191-201 (2017).\n35. M. Hofer, M. Grube, C.F. Burmeister, P. Michalowski, S. Zellmer, A. Kwade, Effective mechanochemical synthesis of sulfide solid electrolyte Li PS in a high energy ball mill by 3 4 process investigation. Advanced Powder Technol. 34, 104004 (2023).\nThe following discussion was added to line 203 to accompany the new Supplementary Fig. S3.\n“The bonding between the thiophosphate units of LPSCl and sulfur, facilitated by the high energy synthesis used in this work, ensures intimate contact of the sulfur/SSE interface with high sulfur utilization shown in Fig. 2a. A high milling intensity is required to induce the interfacial reaction between sulfur and LPSCl. Sulfur utilization was found to decrease when using lower milling intensities (fig. S3), suggesting incomplete formation of the Li PS phase.” 3 4+n\nNew Fig. S3\nFig S3. Evaluation of lower milling intensities using the one-step milling process in Li-In half cells. a) 1st formation cycle normalized by sulfur mass. 1st and 2nd voltage profiles normalized by total active mass for b) 300 rpm, c) 400 rpm, and d) 500 rpm.\nAdditional emphases to the role of milling intensity were added to line 221 of the manuscript after the ionic conductivity measurements, where long milling durations were explored.\n“These results highlight the importance of limiting milling durations, where high milling intensities used for short durations were found to be sufficient to complete the interfacial reaction between sulfur and LPSCl while preserving the ionic conductivity of the SSE and cathode composite.”\nComment 2: Electrochemical redox mechanism during cycling of 1-step milled ASSLSBs. What are the dominant redox reactions occurring at each potential during cycling in the 1-step milled ASSLSB? Furthermore, does Li3PS4+n contribute to the overall capacity? Since the reversible redox reaction of the SSE appears to occur from the second cycle (Fig. 3a), the voltage profile and dQ/dV analysis beyond the second cycle should be done to clarify the electrochemical behavior of the proposed ASSLSBs. It seems that the SSE’s redox activity occurs at higher potentials than that of sulfur. (Fig. 3a and 3b).\nResponse: The reviewer has brought up some insightful points and great questions. We appreciate the opportunity to clarify and will address each question point by point.\n1. What are the dominant redox reactions occurring at each potential during cycling in the 1-step milled ASSLSB?\nFrom our work, we show that the 1-step high energy milling process facilitates sulfur bonding between sulfur and sulfur in LPSCl, where 6.5 wt.% of the composite (~ 20 % of sulfur) undergoes an interfacial reaction with LPSCl. The remaining 23.5 wt.% is unreacted sulfur. To better understand the dominant redox reactions, cells with decreasing sulfur wt.% were fabricated. Capacity was normalized by the total active mass (sulfur and LPSCl) to highlight individual capacity contributions.\nDuring the initial discharge (Region I), from OCV to 1.3V, the dominant reaction is the reduction of sulfur from the unreacted bulk sulfur in the composite. For cells with higher sulfur wt.%, most of the discharge capacity is obtained at this potential, meaning this is from unreacted bulk sulfur. This reduction process starts at 2.1V, commonly observed for Li-S solid-state cells attributed to the solid-solid reduction to Li S. Region II starts at 1.3V where the reacted sulfur in Li PS phase begins to reduce. This reduction peak is 3 4+n easily identifiable in the dQ/dV curves for all cells, but prominent for the cell with lower sulfur to LPSCl wt.% due to its higher capacity contribution compared to unreacted sulfur.\nThe early reduction of the sulfur rich LPS phase is expected and follows a similar redox behavior of LPS, which reduces at a higher potential than LPSCl. The Cl in LPSCl increases its reduction stability1, which is why the LPSCl/C cell reduces at a lower potential compared to LPS. This is also supported by the behavior of the SSE/C cells shown below, where LPS exhibits a narrower electrochemical stability window. The dQ/dV curves of LPSCl/C (no sulfur) and various S/LPSCl wt.% highlight the S and LPSCl interaction, which creates a less stable LPS phase that readily reduces before LPSCl in the 1st discharge. Since we observe that LPS reduces at a higher potential and delivers a higher discharge capacity, we would expect the reduction of this phase to dominate at this potential until near 1.13V, where LPSCl begins to reduce.\nNew Fig. S12\nWe hypothesize that the reacted sulfur in sulfur rich Li PS phase actively participates 3 4+n in reversible redox behavior following similar behavior to LPS. The additional sulfur in PS 3- increases its functional capacity contribution and reduces its stability. Both the 4 sulfur rich LPS phase and LPSCl are expected to reversibly form LPS.\nRegion III is from 2V till 2.7V where Li S oxidation is dominated by the conversion of 2 the formed Li S from the unreacted bulk sulfur. However, both LPSCl and LPS phases 2 will contribute capacity at this potential as shown in the dQ/dV curves. Higher wt.% of sulfur will contribute higher capacity at this potential (meaning there was more of the Li S phase due to the higher wt.% of sulfur within the composite that reduced). Region 2 IV starts at 2.7V where oxidation capacity is obtained from both the Li PS and LPSCl 3 4+n phases, forming elemental sulfur and LPS. The additional oxidation capacity is recoverable upon the subsequent discharge.\nLPSCl forms a reversible LPS phase after the first cycle. Li PS actively participates in 3 4+n the redox reactions. We are unable to quantify the sulfur chain length in the Li PS 3 4+n phase, but Li PS is estimated as the starting phase based on the calculations performed 3 4+3 below.\n1. Y. Liu, et al., Revealing the Impact of Chlorine Substitution on the Crystallization Behavior and Interfacial Stability of Superionic Lithium Argyrodites, Adv. Funct. Mater. 2207978 (2022).\nWe have clarified and elaborated the discussion on the dominant redox reactions within the manuscript.\n“Four main regions are labeled in the dQ/dV curves. Region I highlights the dominant reaction during the discharge, attributed to the reduction of sulfur. Cells with higher sulfur wt.% deliver most of their discharge capacity at this potential, meaning this is from unreacted bulk sulfur. Region II highlights the beginning of SSE reduction. Starting at 1.3V, the reacted sulfur in Li PS begins to reduce. This reduction peak is identifiable in the dQ/dV curves for all cells, but 3 4+n mostly prominent for the cell with the lowest sulfur wt.% due to a higher capacity contribution from this phase. The Li PS phase is expected to follow similar redox as LPS, which reduces at 3 4+n a higher potential than LPSCl (fig. S12). Therefore, we would expect the reduction of this phase to dominate until 1.15V, where LPSCl begins to reduce. We can estimate their individual discharge capacity contributions based on the reduction potentials obtained from dQ/dV. From OCV to 1.3V, 1 mAh is obtained, as estimated above. The remaining capacity from 1.3 to 1.15V can be attributed to Li PS , which is responsible for 3 4+n 7.5% of the overall discharge capacity. The remaining capacity obtained from 1.15 to the lower voltage cutoff of 1V is from LPSCl, which contributes 9.2%. The specific capacity of the reacted sulfur is estimated to be 553 mAh g-1 during the discharge, likely existing in the Li PS phase after synthesis. These 3 4+3 calculations are presented in table S2. The additional sulfur in the PS 3- framework increases its functional capacity and 4 reduces the electrochemical stability of the SSE. Region III is attributed to Li S oxidation. Higher 2 wt.% of sulfur will contribute higher capacity at this potential since more Li S was formed during 2 the discharge and is now being oxidized. The cell with just LPSCl and carbon will contribute capacity at this potential, attributed to LPS oxidation as investigated from previous work by Tan et al.26. Region IV starts at 2.7V where SSE oxidation occurs. This additional oxidation capacity is recoverable upon the subsequent discharge, supporting the formation of a reversible redox-active phase.”\n26. D. H. S. Tan, E. A. Wu, H. Nguyen, Z. Chen, M. A. T. Marple, J. M. Doux, X. Wang, H. Yang, A. Banerjee, Y. S. Meng, Elucidating Reversible Electrochemical Redox of Li PS Cl Solid 6 5 Electrolyte. ACS Energy Lett., 2418–2427 (2019).\n2. Furthermore, does Li PS contribute to the overall capacity?\n3 4+n\nYes, the sulfur rich conductive phase does contribute capacity. We calculated the estimated capacity based on known masses of each component within the cathode composite and compared it to results obtained experimentally.\nAs a starting point, we used a model cell composition. Our initial electrochemical results exhibited an areal capacity of 1.6 mAh cm-2 which used a total cathode composite mass of 2.5 mg. The cathode consists of 23.5 wt.% unreacted sulfur, 6.5% reacted sulfur, 50 wt.% LPSCl, and 20 wt.% carbon. A typical cell with this loading delivers 1.2 mAh and 1.60 mAh after the discharge and charge. The table below shows the expected capacity if every component fully contributes. The specific capacity used for unreacted and reacted sulfur is 1675 mAh g-1. The specific capacity used for LPSCl was obtained experimentally using cells with a LPSCl/C cathode, which was 115 mAh g-1 and 355 mAh g-1 during the discharge and charge. For the S/LPSCl/C cell, 1.38 mAh and 1.70 mAh is estimated after the discharge and charge. These capacities are close to what was obtained experimentally.\nAssuming full utilization of the unreacted sulfur mass, 1 mAh is estimated. Since the total discharge capacity is 1.2 mAh, the remaining 0.2 mAh is either from the reacted sulfur in Li PS , LPSCl, or a mixture of both. 3 4+n\nAfter the charge, a total cell capacity of 1.6 mAh is obtained. Assuming similar electrochemical behavior as for the LPSCl/C cells (80/20 wt.%), LPSCl is expected to deliver a charge capacity of 355 mAh g-1, which equates to 0.44 mAh using a known mass of 1.25 mg of LPSCl within the composite. To estimate the remaining capacity for Li PS the unreacted sulfur and LPSCl is subtracted from the total cell capacity, 3 4+n, equating to 0.16 mAh. Using the dQ/dV curves during the initial discharge, we can estimate the discharge capacity based on the potentials that Li PS and LPSCl reduces. 3 4+n\nLi PS phase reduces at 1.3V while LPSCl reduces at 1.15V. Using these reduction 3 4+n potentials from the dQ/dV curves, we can estimate the capacity obtained. From OCV to 1.3V, 1 mAh is obtained as estimated above. The remaining capacity from 1.3 to 1.15 results in the following estimation.\nThe specific capacity of sulfur in Li PS is estimated to be 553 mAh g-1 during the 3 4+n discharge and 984 mAh g-1 during the charge. Since we validated the existence of this phase after cycling, we can conclude that its redox is reversible although a higher capacity is expected during the charging process. Using these estimations, we can assign capacity contributions for each component.\nTable S2. Estimated capacity contributions between unreacted sulfur, reacted sulfur in Li3PS , and LPSCl.\n4+n\nThese estimations have been added to the supplementary information as Table S2, with additional discussion regarding the capacity contribution of the reacted sulfur in the Li PS 3 4+n phase.\n“In this system, three active materials exist: bulk LPSCl, the formed interlayer phase (Li PS ) 3 4+n consisting of sulfur that has reacted with LPSCl, and unreacted elemental sulfur. TGA was used to quantify the reacted and unreacted sulfur mass, where 23.5 out of the 30 wt.% of sulfur remains unaltered after synthesis. Using the known sulfur and LPSCl mass, we can estimate the capacity contributions of the active materials and compare to what was obtained experimentally. For a 1.6 mAh cm-2 cell using a total composite mass of 2.5 mg, 23.5 wt.% unreacted sulfur is expected to deliver 1 mAh (using 1675 mAh g-1). Therefore, during the discharge, unreacted bulk sulfur is estimated to contribute 83% of the capacity with the remainder attributed to LPSCl and Li PS reduction. To isolate the capacity expected from just LPSCl, cells with just LPSCl and 3 4+n carbon were fabricated using a Li-In alloy as a lithium source.”\n“Subtracting the expected LPSCl and unreacted bulk sulfur capacity from the total cell capacity obtained after the charge leaves the remaining capacity to be assigned to reacted sulfur in Li PS . This equates to a capacity contribution of 10% for the reacted sulfur and 27.5% for 3 4+n LPSCl after the charge, suggesting that the reacted sulfur in Li PS and LPSCl is redox-active 3 4+n and actively participates capacity during cycling.”\n“We can estimate their individual discharge capacity contributions based on the reduction potentials obtained from dQ/dV. From OCV to 1.3V, 1 mAh is obtained, as estimated above. The remaining capacity from 1.3 to 1.15V can be attributed to Li PS , which is responsible for 3 4+n 7.5% of the overall discharge capacity. The remaining capacity obtained from 1.15 to 1V is from LPSCl, which contributes 9.2%. The specific capacity of the reacted sulfur is estimated to be 553 mAh g-1 during the discharge, likely existing in the Li PS phase after synthesis. These 3 4+3 calculations are presented in table S2. The additional sulfur in the PS 3- framework increases the 4 functional capacity and reduces the electrochemical stability of the SSE.”\n3. Since the reversible redox reaction of the SSE appears to occur from the second cycle (Fig. 3a), the voltage profile and dQ/dV analysis beyond the second cycle should be done to clarify the electrochemical behavior of the proposed ASSLSBs. It seems that the SSE’s redox activity occurs at higher potentials than that of sulfur. (Fig. 3a and 3b).\nResponse: The reviewer has made an insightful observation regarding the SSE redox activity occurring at higher potentials than that of sulfur. We have addressed this discrepancy above, where LPSCl will reduce at lower potentials than that of sulfur and the reacted sulfur in Li PS . 3 4+n This is because sulfur reduces at 2.1V, with most of its capacity obtained at the 1.9V plateau. The Li PS phase is expected to follow similar redox behavior as LPS, which reduces at higher 3 4+n potentials than that of LPSCl. This is illustrated in the dQ/dV curves where cells with 10 wt.% of sulfur clearly highlight this reduction peak due to the larger capacity contribution of the Li PS 3 4+n phase. This reduction peak is still evident for cells with 20 and 30 wt.% sulfur, albeit with smaller capacity contribution.\nAs suggested by the reviewer, we also performed dQ/dV analysis on the 3rd cycle for all cells. This better highlights the reversible LPS phase formed by both LPSCl and the sulfur rich interlayer. The LPSCl/C cell exhibits reversible redox behavior centered near the reduction and oxidation potential of sulfur, which has previously made deconvoluting this behavior challenging. In the 3rd cycle, clear evidence of SSE redox and oxidation is present in all cells at 1.3 V and 2.7V.\nWe have updated the main Figure 3 by adding the dQ/dV curves of the 3rd cycle, which better highlights the reversible redox contribution from LPSCl and Li PS phase. 3 4+n\nFig. 3 | LPSCl redox activity and capacity contribution. Voltage profiles of LPSCl/C Li-In half cells evaluated under (a) sulfur voltage limits. b, Cyclic voltammetry of sulfur and LPSCl cathode composites with a Li-In anode. c, dQ/dV curves of the 1st and (d) 3rd cycle for cells with decreasing sulfur wt.% (S/LPSCl/C). The areal loading of sulfur was 0.3 mg cm -2, 0.84 mg cm -2, and 1 mg cm -2 corresponding to a total active mass of 2.8 mg cm -2, 2.9 mg cm -2, and 2.5 mg cm -2 for 10 wt.%, 20 wt.%, and 30 wt.% sulfur composites, respectively.\nThe following discussion has been added to manuscript.\n“This additional oxidation capacity is recoverable upon the subsequent discharge, supporting the formation of a reversible redox-active phase. This is highlighted in the dQ/dV curves for the 3rd cycle (Fig. 3d), where reversible SSE redox occurs at similar potentials as sulfur conversion. Despite a portion of the sulfur reacting with LPSCl, sulfur utilization within each of these cells can be estimated. Cathodes with 30 wt.% sulfur are 84% utilized during the 1st discharge. Sulfur utilization significantly increases for lower sulfur weight percentages, increasing to above 90% for the cathode with 10 wt.% sulfur. These estimations are listed in table S3.”\nComment 3: Stability of Li3PS4+n during cycling. Beyond confirming the formation and functionality of Li3PS4+n in the initial stage, how does this material evolve during cycling? Is it retained or transformed in later cycles? Please clarify and provide evidence for these changes.\nResponse: The reviewer has asked a great question. The sulfur rich Li PS phase was originally 3 4+n identified by observing a redshift of the E symmetric S-S bending at 152 cm-1 to shorter 2 wavelengths. The degree of the shift is influenced by the sulfur chain length in Li PS , with 3 4+n shorter sulfur chains corresponding to lower wavenumbers, as reported by Lin, et al., Angewandte Chemie - International Edition, 52, 7460–7463 (2013). To gain more insight on how the Li PS phase can evolve or be preserved with cycling, Raman spectroscopy was 3 4+n conducted on a cycled sulfur cathode after 5 cycles and 100 cycles.\nThe Li PS phase was confirmed to be retained after both short- and long-term cycling. A clear 3 4+n redshift is observed in the Raman spectra for all cathode samples, with the symmetric S-S bending region being identified at 150 cm-1. This is comparable to the peak position for the cathode composite directly after synthesis. The retainment of this conductive phase helps explain the long cycle life and sustained sulfur utilization obtained with this cathode system.\nThe following figure has been added to the Supplementary Information as Fig. S14.\nNew Fig. S14\nFig. S14. Raman spectra focusing on the symmetric S-S bending region of sulfur cathodes after cycling.\nThe following discussion has been added to the manuscript. “The high charge capacity yet low peak intensity and amorphous background indicate that the reformed sulfur is amorphous as in the pristine state and suggests the sulfur rich interlayeris preserved with cycling. This was confirmed by conducting Raman spectroscopy on the cycled sulfur cathode, where a redshift to lower wavenumbers at the S-S bending region is observed (fig. S14), as previously shown for the pristine composite.”\nComment 4: Stability and structural evolution of milled sulfur. It is stated that micron-sized and sub-micron sulfur were obtained via 400 rpm milling, yet the final cathode composite is fabricated through 500 rpm milling. Is it reasonable to assume that the particle size from the pre-milling step is retained after subsequent composite fabrication? Also, does milling affect only the particle size, or are there other changes such as in crystallinity or bonding structure of sulfur? These aspects are critical to cathode design, and a more thorough explanation is required to convincingly support the role of milled sulfur.\nResponse: The reviewer brings up a good point regarding the role of pre-milled sulfur. It should be clarified that two different milling parameters beyond milling intensity were used, which is the sample to milling media weight ratio. A higher milling media to sample weight ratio was used for composite synthesis (30:1), which induces more energy into the system due to a higher collision frequency between the milling media and sample. A lower milling media to sample weight ratio (10:1) was used to reduce the sulfur particle size prior to composite synthesis. This lower weight ratio induces more frictional stress from particle-to-particle interaction, leading to more breakage. This distinction has been added within the methods section and shown below.\n“For the sulfur cathode, elemental sulfur (99.98%, Sigma Aldrich) was either used as received or milled at 400 rpm for 10 and 24 hours for micron and sub-micron particles with a sample to milling media weight ratio of 1:10.”\n“For electrochemical evaluation, optimal cathode composites were milled for 1 hour (unless otherwise specified) at 500 rpm (unless otherwise stated) using a planetary ball mill with a sample to milling media weight ratio of 1:30.”\nPre-milling sulfur was found to induce more interfacial contact between the SSE and carbon, leading to higher utilization and better rate performance. We were also curious about the impact of high energy milling during synthesis had on particle morphology and size. SEM was first conducted on the bulk sulfur powders and composite after synthesis. We broke up the previous Fig. S16 to help with the flow for the addition of SEM-EDS mapping shown below. Fig. S22 shows that the Bulk sulfur exhibits a wide particle size distribution from 100 to 10 microns. After composite synthesis, a large sulfur particle is identified and surrounded by smaller LPSCl particles (fig. S22). This distinction was made due to their differences in contrast (differences in atomic number between sulfur and LPSCl) from the backscattered electron detector. This large sulfur particle is near 50 m in diameter, exhibiting a comparable size prior to composite synthesis (fig. S22a).\nFig. S22. Scanning electron microscopy (SEM) images of unmodified sulfur and composite.\nPowders of a) Bulk Sulfur as received and b) composite after synthesis. c) Top view after cathode composite fabrication.\nSEM-EDS mapping was also conducted on another batch of composite. We did find EDS mapping to be challenging on the composite since LPSCl will contribute to the sulfur signal. The Cl and P K mapping was also included but was still hard to discern the particles due to the good composite distribution and interfacial bonding between sulfur and the SSE. However, the P and Cl mapping show lower intensity at the location of the highlighted sulfur particle. These results are shown below.\nA large sulfur particle was identified and labeled in the above image, exhibiting a much darker contrast compared to the surrounding smaller LPSCl particles, and well within the expected range for these unmodified particle sizes. Smaller LPSCl particles are surrounding the sulfur particle, where their size is also retained. It is possible that some particles that experienced high enough frictional stress were reduced slightly after the 1-hour high energy milling process. However, this reduction if any, was not significant enough given the ability to still identify large sulfur particles, nor was it significant enough to improve the Bulk S composite’s rate performance, requiring further particle size reduction. For the micron and sub-micron case, particle sizes were not significantly reduced. We also conducted SEM-EDS mapping on the Micron and Sub-micron Sulfur composites, where particles were found to still be on the micron scale.\nWe included a new Supplementary Information figure to support the need and motivation of pre-milling sulfur for enabling higher rate performance, shown below.\nNew Fig. S24\nFig S24. SEM-EDS mapping of the cathode composite powders synthesized with Bulk Sulfur.\nTo address the potential changes of crystallinity of sulfur after the pre-milling step, we conducted XRD on the milled sulfur samples, comparing to the unmodified Bulk sulfur. The spectra are shown below. All sulfur samples exhibit diffraction patterns consistent with elemental sulfur, with no observable changes in crystallinity. The main diffraction peak at the (222) reflection is present at 10.5 2 for all samples, indicating no changes within the lattice. These results validate that pre-milling sulfur to reduce its particle size does not affect its crystallinity or structure and only reduces the size of their particles. The XRD spectra has also been added into the Supplementary Information.\nNew fig. S26.\nFig. S26. XRD spectra of sulfur particles after particle size reduction to the micron and sub-micron scale.\nThe following discussion has been added to the manuscript, referencing the new Supplementary figures. “After composite synthesis, which subjects sulfur to high energy milling with LPSCl and carbon, it is possible that particle sizes can reduce, although, sulfur particles were found to retain close to their original size as observed from the SEM-EDS mapping results (fig. S24). These particle sizes, while sufficient for low-rate cycling, need to be reduced to improve Li+ transport and enable higher rate operation. Therefore, sulfur particle sizes were produced on the micron and sub-micron scale (fig. S25). XRD was conducted after the size reduction process to ensure their crystallinity was preserved (fig. S26).”\nComment 5: Clarification of cell compositions used in electrochemical tests. Due to the large number of cathode composite configurations and corresponding electrochemical data, it is difficult to follow the overall flow. Please clearly summarize and specify the cell configurations used in each data panel to improve readability and interpretation.\nResponse: We appreciate the reviewer’s suggestion and provided extra clarification to follow the experimental workflow. We have made sure to specify each cell configuration within the presented electrochemical results. These changes have been reflected within the main manuscript and the Supplementary Information when necessary. We have also added the cell configurations in the figures where multiple cell configurations were used.\nWhen introducing Fig. 2a, we have included details of the cathode composition and cell configuration within the manuscript.\n“A commonly used sulfide electrolyte, carbon type, and weight composition was employed for our study, where unmodified elemental sulfur, argyrodite LPSCl, and acetylene black (AB) carbon were used to fabricate a cathode composite with a weight ratio of 30:50:20 wt.%. This composition was used in all electrochemical tests unless otherwise stated.”\n“To compare the different cathode synthesis methods, cells were fabricated with LPSCl as the separator layer and a Li-In anode. Figure 2a shows their first cycle voltage profiles where increases in utilization and capacity are attained by introducing high energy milling steps to the fabrication process.”\nFor Fig. S8, we have added the following within the manuscript.\n“The interaction between LPSCl and Li S was also investigated. A comparison between cathode 2 preparation methods was conducted, where half cells were fabricated with a Li-In anode.” We have also clarified the cell configuration within the figure.\nFor Fig. 3, we have added the cell composition within the figure caption.\n“Fig. 3 | LPSCl redox activity and capacity contribution. Voltage profiles of LPSCl/C Li-In half cells evaluated under (a) sulfur voltage limits. b, Cyclic voltammetry of sulfur and LPSCl cathode composites with a Li-In anode. c, dQ/dV curves of the 1st and (d) 3rd cycle for cells with decreasing sulfur wt.% (S/LPSCl/C). The areal loading of sulfur was 0.3 mg cm -2, 0.84 mg cm -2, and 1 mg cm -2 corresponding to a total active mass of 2.8 mg cm -2, 2.9 mg cm -2, and 2.5 mg cm -2 for 10 wt.%, 20 wt.%, and 30 wt.% sulfur composites, respectively.”\nFor Fig. 5d, e, and f, we introduced the experimental validation results by adding the following text within the manuscript.\n“To validate the modeling results, room temperature electrochemical performance was performed evaluating each sulfur particle class in Li-In half cells.”\nFor Fig. 6, we have added the cell configuration information within the Fig. 6c and 6d.\nFig. 6 | Quantifying cathode and cell level volume changes. Cryo-FIB images of the (a) micron sulfur and (b) Li S cathode composite at various states of charge. c, Operando pressure monitoring 2 of LCO and Sulfur using a lithiated silicon anode and of (d) LCO and Li S using a µSi anode 2 during the first formation cycle.\nReviewer #5\nIn this work, the authors report the fabrication of high-performance all-solid-state sulfur/Li2S cathodes using a one-step high-energy ball milling method. This process induces the formation of sulfur-rich thiophosphate phases (Li3PS4+n), which facilitate enhanced interfacial ion transport. The study presents a comprehensive investigation into the structural evolution of cathode materials during cycling, the redox activity of cathode components, the influence of electrode microstructure (particularly sulfur particle size) on battery performance, volume-pressure evolution during electrochemical cycling, and the overall electrochemical performance of the all-solid-state batteries. The performance demonstrated in both half-cells and practical full pouch cells is promising and offers insights into the development of practical solid-state lithium-sulfur batteries.\nComment 1: The areal sulfur loading should be clearly specified for all electrochemical data, particularly in Fig. 2a, Fig. 3, and Fig. 5e.\nResponse: The areal sulfur loading has been added to the electrochemical data presented in the manuscript. These changes are shown for each figure below.\nFig. 2 | Characterizing the sulfur cathode composites after synthesis. a, Voltage profiles of the investigated cathode composite preparation methods in Li-In half cells. b, X-ray diffraction (XRD) spectra of the cathode composite after fabrication. c, Raman spectroscopy and (d) X-ray absorption spectroscopy (XAS) of the one-step milled composite. e, Schematic illustrating the surface reaction with LPSCl bonded on the sulfur surface as a result of the one-step milling procedure.\nFig. 3 | LPSCl redox activity and capacity contribution. Voltage profiles of LPSCl/C Li-In half cells evaluated under (a) sulfur voltage limits. b, Cyclic voltammetry of sulfur and LPSCl cathode composites with a Li-In anode. c, dQ/dV curves of the 1st and (d) 3rd cycle for cells with decreasing sulfur wt.% (S/LPSCl/C). The areal loading of sulfur was 0.3 mg cm -2, 0.84 mg cm - 2, and 1 mg cm -2 corresponding to a total active mass of 2.8 mg cm -2, 2.9 mg cm -2, and 2.5 mg cm -2 for 10 wt.%, 20 wt.%, and 30 wt.% sulfur composites, respectively.\nFig. 5 | Modeling sulfur cathode microstructures with electrochemical validation. a, Geometrically modeled sulfur composite electrodes for bulk, micron, and sub-micron sulfur particles. b, Specific active surface area and (c) ionic transport tortuosity as a function of AM content (%) and particle size. d, First formation cycle voltage profiles of sulfur composites with bulk, micron, and sub-micron particles at C/20 (1C = 1.6 mA cm-2). e, Long-term cycling stability at C/2 with discharge capacity being normalized by the active mass (sulfur and LPSCl). f, Nyquist plots and equivalent circuit fitting results from EIS measurements of the micron (top) and sub-micron (bottom) sulfur cathodes at cycle 100, 300, and 500 in their de-lithiated states. g, Distribution of the von Mises stress on the SSE matrix after simulated lithiation. h, Predicted variation of maximum von Mises stress.\nComment 2: For improved clarity, the authors should provide the calculated ionic conductivity values of the cathode composites and Li6PS5Cl (LPSCl) prepared under different milling conditions, in addition to the EIS spectra shown in Fig. S3c,f.\nResponse: We appreciate the reviewer’s suggestion and have added the calculated ionic conductivity values within the manuscript for clarity.\n“Ionic conductivity of the S/LPSCl composites after increasing milling durations were also measured, where an increase from 6×10-6 S cm -1 to 2×10-5 S cm -1 is observed after 1 hour and saturates with continued milling to 10-hours (fig. S4c). Standalone XRD, Raman, and EIS was also measured for the LPSCl catholyte at 1- and 10-hour milling intervals (Fig. S4d-f). Contrary to what was observed for the S/LPSCl composites, continued milling of LPSCl results in a decrease in ionic conductivity from 1.3 ×10-3 S cm -1 to 4×10-5 S cm -1, attributed to the formation of insulative Li S and LPS phases.” 2\nThese calculated ionic conductivities have also been added to Fig. S3c and S3f.\nFig. S4. Characterization of sulfur and LPSCl catholyte after one-step synthesis without carbon. a, XRD spectra of S/LPSCl composite with increased milling durations. b, Raman spectra of composites shown in a. c, Corresponding Nyquist plots of S/LPSCl composites with increasing milling durations. Standalone d, XRD spectra e, Raman spectra, and f, Nyquist plots of LPSCl after 1 hour and 10 hour milling durations with their calculated ionic conductivity.\nComment 3: As X-ray diffraction (XRD) patterns only indicate the crystal structure of the materials, they are unable to identify the amorphous sulfur-rich interlayer (Li3PS4+n) after cycling (line 331). Instead, Raman spectroscopy should be conducted to validate the existence of the interlayer phase after cycling.\nResponse: The reviewer has provided an insightful suggestion. XRD provided structural information confirming that the poorly crystalline sulfur composite is retained. To support the assumption that the amorphous sulfur-rich interlayer is preserved, Raman was conducted on sulfur cathodes that were cycled for 5 and 100 cycles. This was done to validate the long-term existence of the formed conductive interfaces.\nThe Raman spectra of the cycled cathodes are shown below where the Li PS phase was 3 4+n identified after short and long-term cycling. A clear redshift is observed in the spectra at the symmetric S-S bending region, from 153 cm-2 to 150 cm-1. These results are comparable to the spectra of the cathode composite after synthesis, exhibiting a peak attributed to the S-S bending region at 149.5 cm-1. The retainment of this conductive phase helps explain the long cycle life and sustained sulfur utilization obtained with this system.\nThe following figure has been added to the Supplementary Information as Fig. S14.\nNew Fig. S14\nFig. S14. Raman spectra focusing on the symmetric S-S bending region of sulfur cathodes after cycling.\nThe following discussion has been added to the manuscript.\n“The high charge capacity yet low peak intensity and amorphous background indicate that the reformed sulfur is amorphous as in the pristine state and suggests the sulfur rich interlayeris preserved with cycling. This was confirmed by conducting Raman spectroscopy on the cycled sulfur cathode, where a redshift to lower wavenumbers at the S-S bending region is observed (fig. S14), as previously shown for the pristine composite.”\nComment 4: Although sulfur with different particle sizes was used, the high-energy ball milling process shall further reduce their particle sizes. Therefore, SEM-EDS mapping of the pristine cathode surface is recommended to examine the actual sulfur particle size after fabrication.\nResponse: The reviewer brings up a valid point as we were also curious about the possibility of reducing particle sizes with the planetary ball mill during the short (1 hour) high energy (500 rpm) milling protocol. SEM was first conducted on the unmodified sulfur composite after synthesis. We broke up the previous Fig. S16 to help with the flow for the addition of SEM-EDS mapping shown below. SEM of the unmodified Bulk sulfur powder is shown in Fig. S22a, where a wide particle size distribution from 100 microns to 10 microns is observed. After the high energy ball-milling process, a large sulfur particle is identified surrounded by smaller LPSCl particles shown in Fig. S22b. This distinction was made due to their differences in contrast (differences in atomic number between sulfur and LPSCl) from the backscattered electron detector. This large sulfur particle is near 50 m in diameter, exhibiting a comparable size prior to composite synthesis.\nFig. S22. Scanning electron microscopy (SEM) images of unmodified sulfur and composite.\nPowders of a) Bulk Sulfur as received and b) composite after synthesis. c) Top view after cathode composite fabrication.\nSEM-EDS mapping was also conducted on another batch of composite to verify. We did find EDS mapping to be challenging on the composite since LPSCl will contribute to the sulfur signal. The Cl and P K mapping was also included but was still hard to discern the particles due to the good composite distribution and interfacial bonding between sulfur and the SSE. However, the P and Cl mapping do show lower intensity at the location of the highlighted sulfur particle. These results are shown below.\nA large sulfur particle, exhibiting a much darker contrast compared to the surrounding smaller LPSCl particles was identified and labeled in the above image. Its particle size is 30 m in diameter, still on the expected range for precursor sulfur particle size and also indicating that particles still need to be further reduced. Smaller LPSCl particles surrounding the sulfur particle, shows that its size is also retained. It is possible that some particles that experienced high enough frictional stress were reduced slightly after the 1-hour high energy milling process. However, this reduction if any, was not significant enough given the ability to still identify large sulfur particles, nor was it significant enough to improve the Bulk S composite’s rate performance, requiring further particle size reduction.\nFor the micron and sub-micron case, particle sizes were found to not significantly reduced either.\nFig. S25. Scanning electron microscopy (SEM) images of milled sulfur particles and their composites. Powders of a) Micron scale sulfur, b) composite after synthesis, and c) top view after cathode composite fabrication. Powders of d) Sub-micron scale sulfur, e) composite after synthesis, and f) top view after cathode composite fabrication.\nWe also conducted SEM-EDS mapping on the Micron and Sub-micron Sulfur composites, where particles were found to still be on the micron scale.\nWe included a new Supplementary Information figure to support the need and motivation of pre-milling sulfur to enable high rate performance, shown below.\nNew Fig. S24\nFig S24. SEM-EDS mapping of the cathode composite powders synthesized with Bulk Sulfur.\nAs suggested by the reviewer, we also did SEM-EDS mapping of the cathode composite after fabrication. A slightly lower densification time was used for these samples in an effort to better differentiate sulfur and LPSCl. The higher contrast of these images compared to those in Fig. S22 are due to the higher accelerating voltage of 10kV used here for EDS acquisition. This higher accelerating voltage better highlights interfaces between the densified particles. While we are unable to fully differentiate the particles using EDS, the SEM images of the powder composites provide enough evidence to show that the sulfur and LPSCl particles are not drastically reduced after the short duration high energy milling. Pre-milling the sulfur particles prior to composite synthesis is helpful to improve interfacial contact between all components and create more surface area for sulfur and LPSCl to interracially react.\nThe following discussion has been added to the manuscript, referencing the new Supplementary figures to validate that particles are not significantly reduced after synthesis.\n“After composite synthesis, which subjects sulfur to high energy milling with LPSCl and carbon, it is possible that particle sizes can reduce, although, sulfur particles were found to retain close to their original size as observed from the SEM-EDS mapping results (fig. S24). These particle sizes, while sufficient for low rate cycling, need to be reduced to improve Li+ transport and enable higher rate operation. Therefore, sulfur particle sizes were produced on the micron and sub-micron scale (fig. S25). XRD was conducted after the size reduction process to ensure their crystallinity was preserved (fig. S26).”\nComment 5: It is unclear in Fig. 6c and d which curves correspond to pressure changes, and which represent voltage during cycling. The figure should be redrawn to improve clarity. Additionally, due to the relatively small pressure changes in all-solid-state lithium-sulfur batteries, the authors are encouraged to reduce the y-axis range for pressure to better illustrate the pressure evolution trends during cycling.\nResponse: We appreciate the reviewer’s suggestion and have made sure to differentiate between the pressure and voltage profile curves by adding their own y-axis in a stacked subplot format. The y-axis range for the pressure curves was also reduced to better illustrate the pressure changes with cycling. These changes to Fig. 6c and 6d are shown below.\nFig. 6 | Quantifying cathode and cell level volume changes. Cryo-FIB images of the (a) micron sulfur and (b) Li S cathode composite at various states of charge. c, Operando pressure monitoring 2 of LCO and Sulfur using a lithiated silicon anode and of (d) LCO and Li S using a µSi anode 2 during the first formation cycle.\nComment 6: It is recommended to use the same color (yellow) for sulfur particles in fig. 7d as in Fig. 7a to maintain consistency.\nResponse: We thank the reviewer for their suggestion and have changed the sulfur particle color to yellow in Fig.7d for consistency. This modification is shown below.\nFig. 7 | High loading sulfur and Li S electrochemical performance with energy density 2 outlook. a, Schematic of the sulfur cathode half-cell architecture in pellet cells. b, First cycle voltage profiles with increasing areal capacity. b, Long term cycling stability at the 11 mAh cm-2 level evaluated at room temperature. d, Schematic of the dry process sulfur cathode paired with a lithiated silicon anode in pellet cells. e, Rate performance at the 5.5 mAh cm-2 level. f, Cycling stability at the 7.4 mAh cm-2 level under C/5 current density. g, Theoretical gravimetric energy density as a function of wt.% of sulfur and areal capacity. h, Gravimetric energy density comparison between sulfur and Li S cathodes with various anodes. Values used for these 2 calculations can be found in table S7. i, Schematic of the Li S/anode-free pouch cell. 2 j, First formation cycle at C/10. k, Cycling performance at the 4.5 mAh cm-2 level under 10MPa and 60C conditions.\nComment 7: The current densities are reported in C-rate, but it is unclear how 1C is defined. Given the varying cathode compositions (e.g., sulfur, Li2S, and LPSCl), all current densities are encouraged to be reported in A g-1 based on the total weight of sulfur/Li2S and LPSCl.\nResponse: We thank the reviewer for their careful review and suggestion. Based on their feedback, we have reported current densities in A g-1 for all electrochemical data presented in the main figures of the manuscript. These changes are shown below for each modified figure.\nFig. 2 | Characterizing the sulfur cathode composites after synthesis. a, Voltage profiles and (b) X-ray diffraction (XRD) spectra of the investigated cathode preparation methods. Cells were constructed at the 1.6 mAh cm-2 level. c, Raman spectroscopy and (d) X-ray absorption spectroscopy (XAS) of the one-step milled composite. e, Schematic illustrating the surface reaction with LPSCl bonded on the sulfur surface as a result of the one-step milling procedure.\nFig. 3 | LPSCl redox activity and capacity contribution. Voltage profiles of LPSCl/C Li-In half cells evaluated under (a) sulfur voltage limits. b, Cyclic voltammetry of sulfur and LPSCl cathode composites with a Li-In anode. c, dQ/dV curves of the 1st and (d) 3rd cycle for cells with decreasing sulfur wt.% (S/LPSCl/C). The areal loading of sulfur was 0.3 mg cm -2, 0.84 mg cm -2, and 1 mg cm -2 corresponding to a total active mass of 2.8 mg cm -2, 2.9 mg cm -2, and 2.5 mg cm -2 for 10 wt.%, 20 wt.%, and 30 wt.% sulfur composites, respectively. Current densities used in the dQ/dV curves were 0.06 A g -1, 0.07 A g -1, and 0.08 A g -1 for 10 wt.%, 20 wt.%, and 30 wt.% sulfur S S S composites, respectively.\nFig. 5 | Modeling sulfur cathode microstructures with electrochemical validation. a, Geometrically modeled sulfur composite electrodes for bulk, micron, and sub-micron sulfur particles. b, Specific active surface area and (c) ionic transport tortuosity as a function of AM content (%) and particle size. d, First cycle voltage profiles at C/20 (1C = 1.6 mA cm -2) e, Long-term cycling stability at C/2 with discharge capacity being normalized by the active mass (sulfur and LPSCl). f, Nyquist plots and equivalent circuit fitting results from EIS measurements of the micron (top) and sub-micron (bottom) sulfur cathodes at cycle 100, 300, and 500 in their de-lithiated states. g, Distribution of the von Mises stress on the SSE matrix after simulated lithiation. h, Predicted variation of maximum von Mises stress.\nFig. 7 | High loading sulfur and Li S electrochemical performance with energy density 2 outlook. a, Schematic of the sulfur cathode half-cell architecture in pellet cells. b, First cycle voltage profiles with increasing areal capacity. b, Long term cycling stability at the 11 mAh cm-2 level evaluated at room temperature. d, Schematic of the dry process sulfur cathode paired with a lithiated silicon anode in pellet cells. e, Rate performance at the 5.5 mAh cm-2 level. f, Cycling stability at the 7.4 mAh cm-2 level under C/5 current density. g, Theoretical gravimetric energy density as a function of wt.% of sulfur and areal capacity. h, Gravimetric energy density comparison between sulfur and Li S cathodes with various anodes. Values used for these 2 calculations can be found in table S7. i, Schematic of the Li S/anode-free pouch cell. 2 j, First formation cycle at C/10. k, Cycling performance at the 4.5 mAh cm-2 level under 10MPa and 60C conditions.\nThe following Supplementary Information figures, Fig. S7, Fig. S8, Fig S10. Fig. S11, Fig. S18, Fig. S19, Fig. S24 have also been updated to report current densities in A g-1.\nComment 8: The authors claim the formation of both Li PS and Li P S based on a single XRD 3 4 4 2 6 crystalline peak (Fig. 4b). This is not sufficient evidence. Additional techniques such as solid-state NMR or Raman spectroscopy should be employed to verify the presence of Li P S . 4 2 6 Moreover, prior literature reports Li P S rather than Li P S as an oxidation product of LPSCl 4 2 7 4 2 6 (Nature Materials 2020, 19, 428–435.) and this discrepancy should be addressed.\nResponse: We thank the reviewer for giving us the opportunity to validate and clarify our findings. We also agree that additional characterization methods should be employed to verify the formation of Li PS and Li P S after cycling. Therefore, we conducted Raman spectroscopy 3 4 4 2 6 on the cycled sulfur cathodes. Raman spectroscopy is a local technique with a lateral resolution of 1 m. Given that the micron sulfur cathode composite consists of a distribution of elemental sulfur w/ sulfur rich Li PS , and LPSCl particles ranging from 1-10 microns, various locations 3 4+n need to be measured to elucidate the entire cathode composition.\nPreviously, spectra were collected at a region consisting of sulfur with evidence of the sulfur rich Li PS phase. This was shown in the new fig. S14 and addressed in Comment 3, providing 3 4+n evidence for the retainment of the conductive interlayer with cycling. From the XANES and XRD results, we expected LPSCl to be retained with some LPS formed after cycling. These results were validated by the newly collected Raman spectra shown below. At regions correlated to be within the SSE matrix of the cathode composite, we were able to identify both Li PS Cl 6 5 (LPSCl) and Li PS (LPS), with their signals convoluted in one peak. 3 4\nNew Fig. S15\nFig. S15. Raman spectra of the cycled sulfur cathode at a region correlating within the SSE bulk. a) Comparison to pristine LPSCl and b) fitted spectra highlighting two different PS 3- 4 phases attributed to LPSCl and LPS.\nCurve fitting was performed to deconvolute the measurement, revealing two peaks at 425 cm-1 and 418 cm-1 both attributed to the thiophosphate unit PS 3- present in LPSCl and LPS. We also 4 added the reference spectra of LPSCl, which highlights that LPSCl only exhibits one main peak at 425 cm -1.\nThe reviewer is correct that Li P S is a primary decomposition product (at the charged state) of 4 2 7 LPSCl and is discussed in multiple works. If Li P S or Li P S was present, a peak would be 4 2 7 4 2 6 identified at 400 cm-1 and 385 cm-1 in the Raman spectra (R. Xu et al., J. Mater. Chem. A, 2017,5, 2829-2834). However, we were also unable to confirm the presence of either Li P S or 4 2 7 Li P S using Raman. In the above spectra, there is a small hump observed near 385 cm-1, which 4 2 6 is the expected peak position of P S 4- but its intensity is very low and would be limited evidence 2 6 to claim P S 4- being formed. Therefore, we have decided to not claim Li P S formation in the 2 6 4 2 6 manuscript and attribute the peak present in the XRD results of the cycled cathode in Fig. 4b to -LPS, validated by Raman spectroscopy and XAS. Below shows the updated Fig. 4 in the main manuscript.\nFig. 4 | Probing electrochemical reversibility. a, Nyquist plots of in-situ EIS measurements of the sulfur cathode composite during the first formation cycle. b, XRD and (c) TGA of high loading (7.6 mg cm -2) sulfur cathodes at the pristine, discharged, and charged states. Quantified sulfur masses are listed in table S4. XANES fitting results as weight percentages of products at each state of charge for (d) LPSCl/C (e) sulfur/LPSCl/C, and (f) Li S/LPSCl/C composites. Error 2 bars represent the uncertainty of the quantified masses. The electrochemical performance of cells used for XRD and XANES are presented in Fig. S20 and S21.\nWe have also added the following discussion referencing the new Supplementary Information figure Fig. S15, confirming the presence of LPS after cycling.\n“A new diffraction peak at 15 2 is also identified after the 1st cycle. Raman spectra collected at a region within the SSE bulk shows a convoluted PS 3- peak, attributed to a mixture of LPSCl and 4 LPS (fig. S15). This new phase after the 1st cycle can be attributed to -LPS, validating the active participation and reversibility of LPS, as hypothesized from the electrochemical and dQ/dV results.”\nComment 9: The error bars in the figures should be clearly defined. The number of replicates used to generate the error bars should be specified in the respective figure captions.\nResponse: The error bars are shown in the LCF results in Fig. 4d-f. These represent the standard error and uncertainty for that specific component weight calculation. The uncertainty is computed during non-linear least squares minimization (performed with Athena Software) between the measured spectra and the linear combination of the reference spectra. The covariance matrix provides the statistical relationships of the fitted parameters as the variance for each computed weight. The standard deviation or error for each weight percentage is then computed by taking the square root of the variance. Weights with high uncertainty can be due to noise or if reference spectra possess similar features. We have defined the error bars in the Fig. 4d, e, f caption shown below.\n“Fig. 4 | Probing electrochemical reversibility. a, Nyquist plots of in-situ EIS measurements of the sulfur cathode composite during the first formation cycle. b, XRD and (c) TGA of high loading (7.6 mg cm -2) sulfur cathodes at the pristine, discharged, and charged states. Quantified sulfur masses are listed in table S4. XANES fitting results as weight percentages of products at each state of charge for (d) LPSCl/C (e) sulfur/LPSCl/C, and (f) Li S/LPSCl/C composites. Error 2 bars represent the standard error and uncertainty of the quantified masses. The electrochemical performance of cells used for XRD and XANES are presented in Fig. S20 and S21.”\nComment 10: The figure caption for Fig. 5f should explicitly state whether the EIS spectra were collected at the charged or discharged state.\nResponse: The reviewer has brought up a good point. We have added the lithiation state of the sulfur cathodes. All cells were in the charged (de-lithiated) state at their stated cycle life when EIS was conducted. This modification is shown in the Fig. 5f caption below.\n“Fig. 5 | Modeling sulfur cathode microstructures with electrochemical validation in Li-In half cells. a, Geometrically modeled sulfur composite electrodes for bulk, micron, and sub-micron sulfur particles. b, Specific active surface area and (c) ionic transport tortuosity as a function of AM content (%) and particle size. d, First cycle voltage profiles of cathode composites with various particles sizes e, Long-term cycling stability at C/2 with discharge capacity being normalized by the active mass (sulfur and LPSCl). f, Nyquist plots and equivalent circuit fitting results from EIS measurements of the micron (top) and sub-micron (bottom) sulfur cathodes at cycle 100, 300, and 500 in their de-lithiated states. g, Distribution of the von Mises stress on the SSE matrix after simulated lithiation. h, Predicted variation of maximum von Mises stress.”\nComment 11: On line 354, what is L4.1PS4 phase? Is it Li4PS4 or Li4.1PS4 phase?\nResponse: Within the sulfur voltage window of 1-3V, LPSCl will not fully reduce. We hypothesize the formation of a lithium rich Li PS phase, Li S, and LiCl. The incomplete 4 4 2 reduction of the SSE preserves the composite’s ionic conductivity for reversible electrochemical performance. We verify LPS formation within the LPSCl/C cathode cell using XAS. We also verify the presence of LPS within the sulfur system using XRD, Raman, and XAS.\nComment 12: Regarding the questions raised by previous Reviewer #2, the authors have addressed some of them but the following still needs further elucidation: Comment 3 and 8 by Reviewer #2 The authors claim the formation of Li P S6, but as previously noted, this has not 4 2 been conclusively demonstrated. Additional characterization is needed to confirm its presence.\nResponse: The reviewer brings up a good point which was addressed earlier in Comment 8. We conducted Raman spectroscopy on the cycled sulfur cathodes. At a region correlated to be within the SSE bulk of the composite with no observable signs of sulfur, we were able to identify both LPSCl and LPS at 425 cm-1 and 418 cm-1. These new results are presented in Fig. S15. A small hump was also observed near 385 cm-1, which is the expected location for P S 4-, but is limited 2 6 evidence to claim its formation. Therefore, we have removed the claim of forming Li P S within 4 2 6 the manuscript and attribute the XRD peak of the cycled cathode composite to be from -LPS only. This is consistent to where the main diffraction peak of -LPS should be located. The formation of LPS was also verified using Raman spectroscopy, clearly shown from the peak fitting results in Fig. S15b. The newly collected Raman spectra is shown below.\nNew Fig. S15.\nFig. S15. Raman spectra of the cycled sulfur cathode at a region correlating within the SSE bulk. a) Comparison to pristine LPSCl and b) fitted spectra highlighting two different PS 3- 4 phases attributed to LPSCl and LPS.\nThe discussion in the manuscript has been revised to the following.\n“A new diffraction peak at 15 2 is also identified after the 1st cycle. Raman spectra collected at a region within the SSE bulk shows a convoluted PS 3- peak, attributed to a mixture of LPSCl and 4 LPS (fig. S15). This new phase after the 1st cycle can be attributed to -LPS, validating the active participation and reversibility of LPS, as hypothesized from the electrochemical and dQ/dV results. The differences between the pristine and charged composite mass loss originate from the sulfur rich interlayer, which increases the thermal stability of the reacted sulfur. The preservation of ionically conductive interfaces along with SSE redox helps sustain high sulfur utilization and cycle life.”\nComment 13: Comment 6 by Reviewer #2 - Redox Reactions. In the response, the authors state that the cathode contains 43.5 wt% LPSCl, 6.5 wt% Li3PS4+n, and 23.5 wt% sulfur, and that the capacity contribution from Li3PS4+n is marginal. However, this analysis is flawed. First, if 6.5 wt% of sulfur has reacted, the corresponding Li3PS4+n must be more than 6.5 wt% due to the added mass of the thiophosphate framework. Second, the chain length of sulfur in Li3PS4+n is unknown, making accurate quantification via TGA alone infeasible. Third, the theoretical capacity of Li3PS4+n (depending on the sulfur chain length) can exceed 500 mAh g-1, making its capacity contribution significant (see Angew. Chem. Int. Ed., 2013, 52, 7608–7611).\nResponse: We thank the reviewer for bringing up this point. The reviewer is correct that we are unable to accurately quantify the sulfur chain length of this phase using TGA. Instead, we can use the TGA results to quantify the unreacted and reacted sulfur mass within the composite. This allows us to estimate the expected capacity we would obtain using a model cell composition and compare with our electrochemical results. We can also use the dQ/dV curves of cells that consist of different sulfur wt.% to determine the reduction potentials of the Li PS phase and LPSCl, 3 4+n where the capacity obtained at those reduction potentials can be correlated back from our electrochemical results.\nUsing a model cell composition of 1.6 mAh cm-2 which correlates to 2.5mg of composite, we can estimate the capacity of unreacted and reacted sulfur assuming full utilization (1675 mAh g-1) and LPSCl capacity based on the constant current discharge and charge results. A total discharge and charge capacity of 1.38 mAh and 1.70 mAh is expected. This is quite close to our electrochemical results, where a cell with this loading delivers near 1.2 mAh and 1.6 mAh during the discharge and charge. Since we know the reacted sulfur exists in the Li PS phase, its full 3 4+n utilization is not likely. Subtracting the unreacted sulfur mass capacity, estimated to be 1 mAh, from the total discharge capacity, we can assign the remaining 0.20 mAh to Li PS and LPSCl. 3 4+n\nFrom the dQ/dV curves, we expect the reacted sulfur in Li PS to reduce first and follow 3 4+n similar redox behavior as LPS. This reduction peak is observed for all cells but is prominent for the cell with the lowest sulfur wt.% since the reduction of Li PS dominates the reduction 3 4+n capacity.\nWe can see that the LPSCl/C cell clearly reduces after Li PS . Li PS phase reduces at 1.3V 3 4+n 3 4+n while LPSCl reduces at 1.15V. Using these reduction potentials from the dQ/dV curves, we can correlate the obtained capacity from our electrochemical results. From OCV to 1.3V, 1 mAh is obtained as estimated above. The remaining capacity from 1.3 to 1.15 is 0.09 mAh. From 1.15 to 1V (our lower cutoff), 0.11 mAh capacity is obtained, attributed to LPSCl.\nThe specific capacity of sulfur in Li PS is estimated to be 553 mAh g-1 during the discharge. 3 4+n This suggests Li PS exists as Li PS after synthesis. This aligns well with the large Raman 3 4+n 3 4+3 redshift, which correlates to shorter sulfur chains.\nLi PS + 6 Li+/e- ⟶ 3Li S + Li PS\n3 4+3 2 3 4\nnF\nQ = \n3.6M\nC g\nn = 6 ,F = 96485 [ ],M = molecular mass [ ]\nmol mol\nQ /Capacity = 582 mAh g-1\nred\nSince we validated the existence of this phase after cycling, we can conclude that its redox is reversible although a higher capacity is expected during the charging process. Using these estimations, we can assign capacity contributions for each component.\nTable S2. Estimated capacity contributions between unreacted sulfur, reacted sulfur in Li3PS , 4+n and LPSCl.\nThese estimations have been added to the supplementary information as Table S2, with additional discussion regarding the capacity contribution of the reacted sulfur in the Li PS 3 4+n phase.\n“In this system, three active materials exist: bulk LPSCl, the formed interlayer phase (Li PS ) 3 4+n consisting of sulfur that has reacted with LPSCl, and unreacted elemental sulfur. TGA was used to quantify the reacted and unreacted sulfur mass, where 23.5 out of the 30 wt.% of sulfur remains unaltered after synthesis. Using the known sulfur and LPSCl mass, we can estimate the capacity contributions of the active materials and compare to what was obtained experimentally. For a 1.6 mAh cm-2 cell using a total composite mass of 2.5 mg, 23.5 wt.% unreacted sulfur is expected to deliver 1 mAh (using 1675 mAh g-1). Therefore, during the discharge, unreacted bulk sulfur is estimated to contribute 83% of the capacity with the remainder attributed to LPSCl and Li PS reduction. To isolate the capacity expected from just LPSCl, cells with just LPSCl and 3 4+n carbon were fabricated using a Li-In alloy as a lithium source.”\n“Subtracting the expected LPSCl and unreacted bulk sulfur capacity from the total cell capacity obtained after the charge leaves the remaining capacity to be assigned to reacted sulfur in Li PS . This equates to a capacity contribution of 10% for the reacted sulfur and 27.5% for 3 4+n LPSCl after the charge, suggesting that the reacted sulfur in Li PS and LPSCl is redox-active 3 4+n and actively participates capacity during cycling.”\n“We can estimate their individual discharge capacity contributions based on the reduction potentials obtained from dQ/dV. From OCV to 1.3V, 1 mAh is obtained, as estimated above. The remaining capacity from 1.3 to 1.15V can be attributed to Li PS , which is responsible for 3 4+n 7.5% of the overall discharge capacity. The remaining capacity obtained from 1.15 to 1V is from LPSCl, which contributes 9.2%. The specific capacity of the reacted sulfur is estimated to be 553 mAh g-1 during the discharge, likely existing in the Li PS phase after synthesis. These 3 4+3 calculations are presented in table S2. The additional sulfur in the PS 3- framework increases the 4 functional capacity and reduces the electrochemical stability of the SSE.”\nComment 14: Comment 9 by Reviewer #2 - Discrepancy Between XANES and XRD Results\nThe quantification of the LPSCl using XANES is intriguing. However, as stated by reviewer #2 the XANES data in Fig. 4e suggest complete decomposition of LPSCl after discharge, whereas XRD (Fig. 4b) still shows strong LPSCl peaks. The authors attribute this to differences in sulfur loading and the shallow probe depth of XAS but provide insufficient supporting evidence. Given the high LPSCl content, full reduction to LPS is unlikely despite the low areal sulfur loading. To support this claim, the authors should include the corresponding electrochemical data (e.g., voltage profiles, dQ/dV curves) for the cells used in the XRD and XAS experiments in the Supplementary Information. Testing conditions and cell compositions should be clearly specified, and capacity contributions quantified. Additionally, the manuscript should include a concise explanation, with a more detailed discussion in the Supplementary Information, regarding the accuracy and limitations of the XANES quantification.\nResponse: We thank the reviewer for giving us an opportunity to clarify this discrepancy between the XRD and XANES results. We used linear combination fitting (LCF) of the XANES spectra to identify overlapping redox products between sulfur and LPSCl. More specifically, LCF deconvolutes the relative proportions of different compositions within the sample of interest. This assumes the composite spectra are a linear combination of their individual components. Quantification software like Athena performs non-linear least squares minimization to minimize the error between measured spectra and a linear combination of the reference spectra. The simulated or fitted spectra is a linear combination of the calculated weight corresponding to a particular reference spectrum. Good fits are highly dependent on accurate collection of reference spectra.\nThe discrepancy between the XRD and XANES results is only applicable for the discharged sulfur sample. The initial XANES analysis resulted in only Li S and LPS being reported, while 2 the XRD spectra of a discharged sulfur cathode showed evidence of LPSCl, with minor LPS formation. Given this discrepancy, we re-evaluated the XANES LCF results for this particular sample. The new LCF results were found to deliver a lower R-factor of 0.0017 compared to 0.057, meaning better statistical significance and a better overall fit. A weight distribution of 20.2 wt.% Li S, 34.4 wt.% LPS, and 45.3 wt.% LPSCl was determined. The new fit results are 2 shown below.\nThe new fitting results resolve the discrepancy between the XRD and XANES. The new fit does not significantly alter our conclusion of the SSE redox products and provides a more realistic cathode composition after discharge. If anything, it better supports our hypothesis of LPS to be formed as a redox intermediate with most of the SSE still existing as LPSCl. This helps retain the cathode composite’s ionic conductivity.\nThe new results have been added to the manuscript on line 384 with the new results updated in Fig. 4e. “The LCF results for the sulfur cathodes are shown in Fig. 4e. After the initial discharge, 20.2 wt.% of Li S is estimated. This supports the high conversion efficiency of the 2 cathode architecture. LPSCl also partially decomposes to LPS, as observed for the LPSCl system. A higher amount of LPS is also present than expected but is likely from the Li PS 3 4+n phase from reacted sulfur.”\nFig. 4 | Probing electrochemical reversibility. a, Nyquist plots of in-situ EIS measurements of the sulfur cathode composite during the first formation cycle. b, XRD and (c) TGA of high loading (7.6 mg cm -2) sulfur cathodes at the pristine, discharged, and charged states. Quantified sulfur masses are listed in table S4. XANES fitting results as weight percentages of products at each state of charge for (d) LPSCl/C (e) sulfur/LPSCl/C, and (f) Li S/LPSCl/C composites. Error 2 bars represent the uncertainty of the quantified masses. The electrochemical performance of cells used for XRD and XANES are presented in Fig. S20 and S21.\nTo address the quantification limitations, we also provided the electrochemical data for the cells used for XANES and XRD. The cell configurations have been included within the new Supplementary figures. All cells used for each system (sulfur, Li S, LPSCl) deliver comparable 2 electrochemical performance between the different discharged states providing confidence in the XANES quantification trends. The cells used for XRD were fabricated with a higher sulfur loading and areal capacity than for those prepared for XANES. This was done to ensure there was enough sample to conduct air-tight XRD, and ultimately enough sample for phase identification. The high loading sulfur cells deliver reasonable utilization at a very high areal capacity (~ 10 mAh cm -2) as shown below in Fig. S20. Therefore, the formed conversion and redox products should be comparable to the lower sulfur loading system. The phases identified with XRD are a bulk analysis of the entire composite while the XANES is a more local/interfacial analysis on near the cathode surface.\nFig. S20. Electrochemical performance of sulfur cathodes used for XRD analysis.\nFig. S21. Electrochemical performance of cells with a) LPSCl, b) Sulfur, and c) Li S cathodes 2 used for XANES quantification.\nTable S5. Sulfur K-edge XAS linear combination fitting results.\nWhile the new fitting results present a better fit to the discharged sulfur cathode spectra, the limitations of XANES, due to shallow probe depth and challenges with analysis of the entire composite originally discussed are still valid and should be mentioned. Clarification and discussion of XANES limitations has been added to the manuscript.\n“It must be noted that XANES quantification of the entire composite can be challenging due to the shallow probe depth when using tender X-rays. Discussion regarding limitations of XANES quantification is included in the Supplementary Information.”\nWe have added a more detailed discussion regarding quantification limitations of XANES within the Supplementary Informatiom.\nQuantification limitations of XANES at the Sulfur K-edge\nXANES was performed in total fluorescence yield (TFY) mode. In this mode, sulfur spectra can distort due to its strong self-absorption effect.1 X-ray energies for the Sulfur K-edge are low (2keV) as compared to hard X-rays (5-10 keV), resulting in a limited probe depth of 500 nm to several microns. This narrows the detection region to particle surfaces and their interfaces. The beam spot size for the beamline used in this work is 0.25 mm  0.16 mm, which covers 0.06% of the electrode area. Therefore, collecting information on the entire cathode composite can be challenging. This is why XRD was conducted in parallel.\nXANES quantification was performed using linear combination fitting (LCF) to identify and quantify overlapping redox products between sulfur, LPSCl, and Li S. LCF deconvolutes the 2 relative proportions of different references within the sample of interest, assuming the composite spectra are a linear combination of individual components. Quantification software such as Athena (used in this work) performs non-linear least squares minimization to minimize the error between measured and fitted spectra. This fitted spectra is a linear combination of the calculated weights, each corresponding to a specific reference spectrum. Weight uncertainties can be introduced from noise and when the reference spectra exhibit similar features (like LPSCl and LPS), where both possess the thiophosphate unit (PS 3-) within their structure. These factors 4 were considered in our analysis.\n1. F. Jalilehvand. Sulfur: not a “silent” element anymore. Chem. Soc. Rev., 2006, 35, 1256– 1268.\nComment 15: XPS should be removed from the Author Contribution.\nResponse: XPS has been removed from the Author Contribution section. We appreciate the careful review by the reviewer.\nReviewer #6\nI think the authors have adequately addressed the concerns from the previous reviewers for publication in this journal.\nResponse: We thank the reviewer for their review and satisfaction of our work for publication in Nature Communications.\nReviewer Comments\nNote: Reviewer’s comments are shown in black, and our responses to the comments are shown in blue.\nReviewer #4\nAll of my comments have now been addressed.\nResponse: We appreciate the reviewer’s satisfaction with our work.\nReviewer #5\nI would like to thank the authors for their extensive efforts in addressing all the previous questions. I am in favor of its publication. In addition, please correct the typo, “LiCO2”, in Fig. 6c and d to “LiCoO2.”\nResponse: We are grateful for the reviewer’s feedback and for their support in publishing our work. We have corrected the typo in Fig. 6c and d."}, {"round": 2, "reviewer_comments": {"reviewer_4": "All of my comments have now been addressed.", "reviewer_5": "I would like to thank the authors for their extensive efforts in addressing all the previous questions. I am in favor of its publication. In addition, please correct the typo, “LiCO2”, in Fig. 6c and d to “LiCoO2.”"}, "author_response": "Response: We are grateful for the reviewer’s feedback and for their support in publishing our work. We have corrected the typo in Fig. 6c and d."}]

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