Moisture-induced surface degradation mechanism of argyrodite Li6PS5Cl under dry-room conditions

Moisture-induced surface degradation mechanism of argyrodite Li6PS5Cl under dry-room conditions | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Research Article Moisture-induced surface degradation mechanism of argyrodite Li 6 PS 5 Cl under dry-room conditions Yoon-Seong Kim, Dong-Hwa Seo, Jeong-Doo Yi, Sihyeon Sung This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-7583174/v1 This work is licensed under a CC BY 4.0 License Status: Posted Version 1 posted You are reading this latest preprint version Abstract Argyrodite Li₆PS₅Cl (LPSC) possesses both high Li-ion conductivity (~ 10 mS cm⁻¹ at room temperature) and mechanical softness, positioning it as a flagship solid electrolyte for next-generation all-solid-state batteries (ASSBs). However, even trace amounts of moisture in industrial dry rooms (dew-point − 60 to − 70°C) rapidly degrade its surface, diminishing ionic transport and impeding scalable processes. Here, we elucidate the moisture-triggered surface degradation mechanism of LPSC under dry-room conditions by combining first-principles calculations with depth-profiling X-ray photoelectron spectroscopy analysis. The combined analysis reveals a five-step sequence: (i) H 2 O adsorptions on S-rich surface, (ii) P–S bond weakening followed by thermodynamically favoured S-O substitutions, (iii) rotation of the O-substituted PS₄ tetrahedra that drives O migration into subsurface layers, (iv) formation of an O-rich Li₆PO₅Cl-like surface, and (v) volume-shrinking phase separation into LiCl, Li₃PO₄, Li₂SO₄, LiOH, and Li 2 CO 3 . The resulting porous, O-enriched layer fails to passivate the electrolyte, causing a 36% drop in ionic conductivity within three days. These mechanistic insights highlight polyhedral-rigidity tuning and moisture-blocking surface chemistries as complementary strategies for stabilizing thiophosphate electrolytes during practical cell fabrication. Materials Chemistry Computational Chemistry Materials Theory and Modeling sulfide argyrodites all-solid-state batteries dry room compatibility degradation mechanism Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Figure 6 Introduction All-solid-state batteries have recently attracted significant attention due to their superior thermal stability and potential high energy density. In contrast, conventional Li-ion batteries (LIBs) employing liquid electrolytes face enduring challenges, such as safety risks associated with flammable organic solvents and limited energy density due to the use of carbon-based anodes like graphite 1,2 . To address these challenges, inorganic solid-state electrolytes (SSEs) have emerged as promising alternatives, offering enhanced thermal stability and enabling the use of lithium metal as an anode 3 . Consequently, extensive research efforts have been devoted to discovering suitable solid electrolyte materials, leading to the development of diverse material classes with distinct structural and compositional characteristics. Among various candidates, sulfide-based inorganic SSEs have emerged as particularly promising due to their high ionic conductivity and high ductility, enabled by the polarizable nature of their anionic frameworks. Several sulfide electrolytes, including Li 7 P 3 S 11 , Li 10 GeP 2 S 12 (LGPS), and the argyrodite compound Li 7 − a PS 6 − a X a (X = Cl, Br, I), have demonstrated high ionic conductivities exceeding 10 mS cm − 1 4–6 . Particularly among the argyrodite compounds, the thio-antimonate iodide argyrodites, Li 6 + x M x Sb 1−x S 5 I (M = Si, Ge, Sn), have exhibited exceptional ionic conductivity of 24 mS cm − 1 7 . Moreover, the inherent ductility of sulfide SSEs helps reduce interfacial and grain boundary resistance. For example, the experimental measurements indicate that Li 6 PS 5 Cl, a prototypical argyrodite, has a low elastic modulus (28.0 ± 1.8 GPa), corroborated by theoretical values of bulk modulus (28.7 GPa) and Young’s modulus (22.1 GPa), facilitating intimate electrode-electrolyte contact under the relatively low fabrication and stack pressures 8–10 . Despite their desirable properties, the poor moisture stability of sulfide SSEs, which leads to the release of toxic H 2 S gas and chemical degradation, has impeded their practical application 11 . To prevent the generation of H 2 S gas, an inert atmosphere is required to avoid reactions between sulfides and moisture, which in turn increases the cost of the necessary infrastructure. Thus, improving compatibility under tightly controlled humidity conditions, such as those in dry rooms, is crucial for the industrial implementation of sulfide-based SSEs. However, even within these controlled environments, they have exhibited limited stability, as indicated by a decrease in ionic conductivity after 24 hours of exposure in the dry room (dew point − 45 ℃) 12 . To address moisture-induced degradation, various approaches have been proposed, including cationic or anionic substitutions based on the hard and soft acids and bases (HSAB) theory, surface engineering to inhibit H 2 O adsorption, and the use of H 2 S-absorbing additives 13–26 . Although the previous approaches have made meaningful progress in addressing the issue, developing new strategies for designing water-resistive materials remains challenging due to a limited understanding of detailed decomposition processes. Therefore, it is essential to establish a comprehensive understanding of surface reactions and their underlying mechanisms under dry room conditions. While various studies have attempted to chemically analyze these phenomena and identify their origin, they have largely focused on the surface-related changes following the air exposure or have been limited to theoretical investigations of the initial stage of the decomposition 27–33 . As a result, they often provide only a general trend of the reactions in the subsequent steps. To overcome these limitations, further studies are needed to reveal the full degradation pathways, from reaction onset to the formation of decomposition products, at both atomic and nanoscale levels. Such insights will enable the development of more intuitive and diverse strategies for enhancing the stability of sulfide electrolytes against air and dry room environments. In this work, we systematically elucidate the surface degradation mechanism of the argyrodite Li 6 PS 5 Cl (LPSC), a representative sulfide-based SSE, under controlled atmospheric conditions containing trace amounts of H 2 O. A combined approach of first-principles calculations and X-ray photoelectron spectroscopy (XPS) depth profiling was employed to gain mechanistic insight into the decomposition process. Thermodynamic and kinetic analyses were performed using first-principles calculations to investigate the reactive behavior of the LPSC surface at the atomic scale and revealed that degradation is initiated by H 2 O adsorption followed by S-O substitution, which locally alters the surface bonding environments. Energetic comparisons of possible substitution mechanisms further clarified the thermodynamic feasibility of these reactions and identified the preferred pathways. In addition, ab initio molecular dynamics (AIMD) simulations tracked the kinetic evolution of O-substituted species, showing that polyhedral rotations of PS 4 units facilitate deeper oxygen penetration and drive phase separation. This process ultimately leads to the formation of a porous surface layer due to volume contraction upon decomposition. To validate our theoretical predictions, XPS depth profiling was performed on LPSC samples after controlled exposure. The resulting phase distribution as a function of depth confirmed the presence of surface-localized reactions and distinct decomposition products compared to the bulk. The progressive changes in composition with increasing depth strongly support the proposed mechanism of volume-shrinking phase separation. Overall, this dual approach combining atomic-scale calculations with experimental depth-resolved analysis provides a comprehensive understanding of surface degradation in LPSC under dry room conditions. These insights offer a foundation for developing more robust sulfide electrolytes with enhanced environmental stability. Results and discussion Surface degradation in the dry room After 3 days of exposure to a dry room environment (dew point − 60 to − 70 ℃), the electrochemical impedance spectroscopy (EIS) analysis results (Supplementary Fig. 1) show a significant decrease in ionic conductivity, from 3.13 mS cm − 1 to 2.00 mS cm − 1 (36.1% reduction). X-ray diffraction (XRD) patterns of the pristine and exposed samples of LPSC (Supplementary Fig. 2) exhibit notable peak broadening in the 2θ range of 24°-33° after 7 days of exposure. This broadening likely indicates the formation of nanoscale decomposition products consistent with previous ex-situ TEM studies on argyrodite degradation under humidity-controlled conditions 28 . XPS analysis (Supplementary Fig. 3) further supports this degradation behavior. Phase distribution at the surface of LPSC, LPSC 5d , and LPSC 25d (representing 5 and 25 days of dry room exposure, respectively) showed an increasing trend in the amount of decomposition products (P 2 S n , SO n 2− , PO 4 3− ), accompanied by a decline in the pristine PS 4 3− signal. These results suggest that the nanoparticles observed in XRD are composed of the decomposition products formed during prolonged exposure. Collectively, the evidence points to surface degradation and nanoparticle formation as key contributors to the observed decline in ionic conductivity. Additionally, theoretical reaction energies supported the experimental surface degradation. Three representative reactions, previously reported for LPSC degradation, were considered 28 : Li 6 PS 5 Cl + 6H 2 O → LiCl + Li 3 PO 4 + 2LiOH + 5H 2 S (1) 2LiOH + CO 2 → Li 2 CO 3 + H 2 O (2) Li 6 PS 5 Cl + 4H 2 O + 2O 2 → LiCl + Li 3 PO 4 + Li 2 SO 4 + 4H 2 S (3) All three reactions exhibit negative density functional theory (DFT)-computed reaction energies (E reaction ), confirming that they are thermodynamically favorable (Supplementary Table 1). Reaction (1), with an E reaction of − 6.222 eV, describes the initial hydrolysis of LPSC in the presence of water, forming LiOH and H 2 S among other products. The generated LiOH can further react with atmospheric CO 2 , as described in reaction (2), yielding Li 2 CO 3 with an E reaction of − 1.465 eV. Notably, reaction (3), which includes O 2 as a reactant, exhibits a significantly larger thermodynamic driving force (E reaction = − 18.991 eV) due to the formation of a stable Li 2 SO 4 . Theoretical trends suggest that the reactions promoted by H 2 O in conjunction with CO 2 and O 2 result in surface deterioration, which greatly affects the formation of decomposition phases and serves as a primary cause of the ionic conduction drop. H 2 O adsorption on the LPSC surface Building on the preceding findings, it is clear that dry room exposure triggers decomposition reactions. Therefore, a comprehensive investigation into the mechanisms underlying this phenomenon is essential. Given experimental limitations in capturing surface reactions that occur at extremely short timescales, DFT calculations were employed to achieve a detailed understanding. To elucidate the decomposition mechanism from the onset of reaction, it is important to observe structural variations and local coordination geometries at the uppermost surface. As shown in Supplementary Fig. 3, a key feature of the deterioration is the oxygen involvement at the surface. Notably, based on reactions (1) and (3), the consistent formation of Li 3 PO 4 and H 2 S gas indicates that the exchange of two atomic species, S and O, fundamentally occurs at PS 4 tetrahedra. To verify the thermodynamic preference of this surface behavior, surface calculations were performed using the slab model of the LPSC (001) surface (Fig. 1 a, b, Supplementary Table 2) 32,34,35 . These calculations focused on the first five atomic layers, with the first layer defined as the layer closest to the surface (Supplementary Fig. 4). As shown in Fig. 1 a, c, two sulfur atomic sites exist on the surface: a non-bonding site (Wyckoff 4d) surrounded by a lithium cage and a P–S bonding site (Wyckoff 16e) within the PS 4 tetrahedron. We evaluated the thermodynamic driving force of the S-O substitution reaction energy (E sub ) at these two potential reaction sites using Eq. (4): E sub = {(E substituted – E initial ) – n(µ O – µ S )} / n (eV O atom − 1 ) (4) In Eq. (4), E substituted and E initial refer to the energies of various O-substituted and pre-substituted slab structures, respectively. The variable n refers to the number of substituted O atoms, and µ O and µ S are the chemical potential of O and S, respectively. As a result, substitution at both sites is energetically favorable, attributed to the increased stability from the more stable Li–O and P–O bonds. Between them, the P–S bonding site (–2.915 eV O atom − 1 ) exhibits a lower energy value than the non-bonding site (–2.535 eV O atom- 1 ), suggesting that it is the more reactive site for oxygen substitution. Prior to the S-O substitution, the adsorption of H 2 O, a key initiator of the surface degradation, is regarded as the most likely event of the reaction, followed by structural changes. Although O 2 and CO 2 are involved in reaction processes (Supplementary Table 1), the principal origin of the decomposition is the H 2 O-induced initiation, as evidenced by discrepancies in reactivity depending on the humidity and atmospheric conditions 26,30 . Accordingly, the influence of initial H 2 O adsorption on the subsequent reactions was investigated to understand the substitution pathway by evaluating the adsorption energy (E ads ) and structural changes. The E ads was calculated based on the energy difference before and after adsorption, as shown in Eq. (5): E ads = E LPSC+H2O – E LPSC – E H2O (5) As shown in Fig. 2 a, spontaneous adsorption occurs regardless of the neighboring sulfur site, although the energetically more favorable reaction occurs near the P–S bonding site. After adsorption of the single H 2 O molecule, changes in surface bonding environments, including weakened P–S bond and newly formed H–S bond, should precede the subsequent S-O substitution and release of H 2 S gas. To verify this, we examined the bond length and integrated crystal orbital Hamiltonian population (ICOHP), which provides insight into bond strength, where a higher ICOHP value indicates weaker bond strength (Fig. 2 a). Compared to adsorption near the non-bonding (4d) site, adsorption near the P–S bonding (16e) site resulted in a significantly longer distance between P and S atoms (2.993 vs. 2.246 Å) of the neighboring PS 4 tetrahedron and a higher ICOHP value (–0.481 vs. − 3.847 eV). This bond weakening originates from the newly formed H–S bond (ICOHP: − 1.027 eV), which occurs between a neighboring PS 4 tetrahedron and the adsorbed H 2 O molecule. This weakened P–S bonding facilitates subsequent S-O substitutions through oxygen sources available in the surrounding environment. Consequently, the substitution reaction ultimately promotes the generation and release of H 2 S gas. These results confirm the hypothesis that the initial S-O substitution, driven by adsorption of H 2 O near the P–S bonding site, dominates the early stage of the degradation process. Similarly, Han et al. reported thermodynamic preference for anionic exchange on the P–S bonding site and the generation of the sulfur-based gases, which were accelerated by the presence of preadsorbed water molecule 32 . This finding implies a spontaneous tendency toward multiple oxygen substitutions triggered by the continuous H 2 O adsorption. Given the favorable formation of the P–O bond (Fig. 1 c, Supplementary Fig. 5), the adsorption effects on the second layer-substituted structure (O-substituted LPSC, LPSCO) were evaluated. In Fig. 2 b, spontaneous surface adsorption (E ads : − 1.766 eV) occurs with the H–S bond formation (bond length: 2.008 Å, ICOHP: − 1.140 eV). Compared to the adsorption on the pristine LPSC surface, the shorter distance between H and S atoms (2.008 vs. 2.357 Å) and a lower ICOHP value (–1.140 vs. − 0.342 eV) were observed during adsorption near the non-bonding site. To reveal the underlying reason for this result, a sum of ICOHP values for Li–S bonds (Total ICOHP Li–S ) was evaluated (Fig. 2 c). The total ICOHP Li–S value of LPSCO is higher than that of LPSC (–4.903 vs. − 5.434 eV), while H 2 O-adsorbed LPSCO (H 2 O-LPSCO) exhibits an even higher value (–4.440 eV). This weakening likely results from the structural3 instability caused by the distortion and Li–S bond elongation owing to the substitution with smaller O 2− (ionic radius: 140 pm) compared to S 2− (ionic radius: 184 pm), as demonstrated in Supplementary Fig. 6. Sano et al. showed that LPSC exposed to atmosphere containing trace amounts of H 2 O (dew point − 20 ℃) and oxygen decomposed into sulfonates, while maintaining the PS 4 3− amount equivalent to the case under the same condition but without O 2 30 . Interpreting this within our framework, it appears that the moisture primarily drives initial surface degradation, subsequently facilitating surface P–S bond weakening and further SO n 2− formation upon reaction with O 2 molecules. Consequently, as these reactions continuously progress, the surface will become entirely O-substituted regardless of the sulfur atomic sites. S-O substitution mechanism During the subsequent stages of decomposition at the surface, two possible scenarios of multiple S-O substitutions can occur: the penetration of substituted O species into neighboring atomic sites, and additional surface substitutions at the exposed sulfur sites. As illustrated in Fig. 3 a, we calculated E sub for single-atomic substitution, which showed that E sub became increasingly negative with increasing depth, indicating a thermodynamic driving force favoring penetration. Furthermore, we identified the most energetically favorable reaction pathway by comparing the E sub of two possible mechanisms at each reaction order grounded in an assumption of sequential substitutions (Fig. 3 b). They consisted of the additional substitution taking place at the 1st or 2nd atomic layer, and the migration of the substituted O to an internal region. Unlike the situation in which only additional substitution could occur in the first reaction order, two mechanisms proceeded competitively from the second order. Given that a more negative E sub indicates greater energetic preference, the additional substitution is preferred in the second reaction order (–3.023 vs. − 0.143 eV O atom − 1 ), which continues until the sixth order, referring to entire O substitutions at the exposed sulfur atoms. Based on the thermodynamic favorability, we can expect that the penetration is preceded by the complete series of O substitutions at the 1st and 2nd layers (LPSCO 12 ), despite the presence of the driving force for the penetration. Along with the spontaneous substitutions occurring at the topmost surface, we carried out AIMD simulations of the LPSCO 12 slab structure to confirm how penetration proceeds. Surprisingly, Fig. 3 c captures a positional reversal along the c-axis (perpendicular to the surface) between underlying sulfur and upper oxygen (O1) atoms in the PS 2 O 2 tetrahedron. A snapshot after 30 ps of simulation (Fig. 3 d) reveals polyhedral rotation. We attribute this result to the anionic size effect, wherein the sulfur anion migrates toward the vacuum to relieve structural stress induced by its larger ionic radius than that of O 2− . As a result, the sulfur initially positioned in the lower layer migrates to the top surface and undergoes continuous substitution as previously confirmed by the greater thermodynamic stability of P–O bond formation through successive anionic exchanges. Following the tetrahedral rotation and subsequent substitutions, further simulations were performed to examine structures involving oxygen substitution extending to the third atomic layer (Fig. 3 e, f). Here, the sulfur atoms diffuse along the positive c-axis (toward the topmost surface), although the rotation requires more simulation time to reach the crossing-point with respect to the vertical height. This delay can be explained by the oxygen chemical potential gradient, consistent with Le Chatelier’s principle. Consequently, the final exposed sulfur species are expected to undergo additional substitutions, which leads to the formation of an O-substituted surface extending down to the third atomic layer (LPSCO 123 ). Further decomposition during the continuous exposure If the fully O-substituted surface functioned as a protective layer against H 2 O molecules, the amount of decomposition products would not continuously increase with exposure time. However, XPS analysis (Supplementary Fig. 3) reveals the opposite trend, indicating that further investigation into the underlying reason for the continuous surface degradation is required. One plausible mechanism is the penetration of the substituted O atom in the third layer of LPSCO 123 . After the depletion of exposed sulfur species through O substitutions, polyhedral rotations of the O-substituted tetrahedron caused the subsequent substitutions at sulfur atoms even in the subsurface during the preceding sequential reactions. If there had been no reappearance of sulfur species on the surface, the oxygen substitution would have been terminated. Thus, given that the presence of sulfur is essential for the progression of the degradation, further downward diffusion of oxygen species from the third layer to the fourth or fifth layer, inducing an exchange with internal sulfur, can be one of the reaction pathways during the continuous surface decomposition. However, as shown in Fig. 4 a, more positive E sub values of the atomic substitution at the fourth (–1.767 eV) and fifth layers (–2.919 eV) compared to the one of the third layer (–3.058 eV) indicate the existence of only a weak thermodynamic driving force for this penetration from the pristine structure. Moreover, since penetration requires the cleavage of stable P–O bonds and the subsequent substitution with the internal sulfur source, we carefully evaluated the feasibility of this coupled reaction using several analyses. Among possible sulfur sites within the interior layers, three distinct atomic sites were considered: i) P–S bonding (16e) sites, ii) non-bonding (4d) sites, and iii) a broken P–S bonding site generated after the formation of LPSCO 123 (Supplementary Fig. 7). In Fig. 4 b, the ICOHP values of each bond relevant to oxygen and sulfur sites from the third to the fifth layer (P–O: substituted O, P–S: P–S bonding site, Li–S: non–bonding site, broken P–S: broken P–S bonding site) are evaluated. Among various bond types, P–O bonds exhibit the lowest values (–7.943 to − 8.160 eV), followed in increasing order by P–S bonds (–5.793 to − 6.188 eV), Li–S bonds (–0.759 to − 1.151 eV), and the broken P–S bond (–0.517 eV). Since both P–O and P–S bonds are intrinsically stable, bond breakages followed by the transformation between these bonding interactions are kinetically unfavorable without any reaction initiators from an external source, such as the adsorption of gas molecules on the surface. Additionally, the thermodynamic favorability of S-O substitutions at sulfurs of non-bonding and broken P–S bonding sites that exhibited weaker bond strength was evaluated. In Supplementary Fig. 8, the structural instability derived from the higher structural energies of S-O exchanged structures compared to the LPSCO 123 is observed, which indicates the absence of a thermodynamic driving force favoring penetration. Furthermore, to verify the non-existence of oxygen migration into deeper layers from a kinetic perspective, AIMD simulations were conducted at 500 K for 100 ps (Fig. 4 c). The result shows that O atoms do not exhibit a downward diffusion tendency: instead, the O-enriched surface is maintained without further reaction. Therefore, the continuous surface decomposition must proceed through an alternative pathway. A plausible mechanism is the formation of an inhomogeneous surface layer as a result of degradation. Following sufficient progression of the oxygen substitution, the Li 6 PS 5 − x ClO x ’s O-rich surface will undergo structural collapse due to inherent instability induced by lattice mismatch between surface Li 6 PO 5 Cl and internal LPSC phases. To quantify this instability, we calculated the surface energy (E surface ) according to Eq. (6): E surface = (E slab – N * E bulk ) / 2A (6) where E slab and E bulk are the energies of slab and bulk structures, respectively, N represents the number of bulk formula units contained within the slab structure, and A is the surface area. Using this equation, the reacted O-rich surface (LPSCO 123 ) exhibits significantly higher surface energy (0.213 eV Å −2 ) compared to the pristine LPSC surface (0.014 eV Å −2 ). This notable increase in surface energy indicates that the reacted O-rich surface is inherently unstable. Furthermore, Li 6 PO 5 Cl is thermodynamically unstable against hydrolysis (Eq. (7)), with a negative DFT-computed reaction energy of − 2.033 eV: Li 6 PO 5 Cl + H 2 O → LiCl + Li 3 PO 4 + 2LiOH (7) Consequently, additional decomposition reactions on the destabilized O-rich surface are expected. Multiple decomposition products are generated, corresponding to the reactions presented earlier (equations (1), (2), and (3)), all of which possess smaller atomic volumes than the pristine LPSC (Fig. 4 d). Considering the volumetric changes upon phase deterioration, the shrinkage of decomposition products leads to the formation of void 36 . Therefore, as schematically illustrated in Fig. 4 e, the volume reduction of the oxidized surface does not adequately protect underlying pristine atomic layers, resulting in continuous surface decomposition and failure to serve as a stable protective layer. Surface with concentrated decomposition products In the absence of oxygen penetration into deeper layers, a surface degradation is driven by the formation of a surface layer enriched with decomposition products resulting from volume-shrinking phase separation. To experimentally validate this scenario, a comparative analysis of the distribution of decomposition phases between the topmost surface and deeper layers was performed by using XPS depth profiling. Prior to this analysis, different intensities of etching beams were tested to identify optimal analytical conditions. If its intensity is excessively strong, unintended decomposition of the pristine PS 4 3− phase could occur, resulting in misleading results. As demonstrated in Supplementary Fig. 9, both the 0.5 kV and 1.0 kV etching beams yielded similar phase distributions to the non-etched LPSC sample. Thus, the 0.5 kV beam was applied for LPSC 5d analysis, which had a relatively mild degree of decomposition, while the 1.0 kV beam was selected for LPSC 25d analysis to ensure time- and cost-efficiency. During the XPS depth analysis, the presence of SO n 2− (including sulfites and sulfates) and PO 4 3− is a clear indicator for assessing changes in phase distribution, as these phases are absent in pristine LPSC (Fig. 5 a, b, Supplementary Fig. 9). Especially for SO n 2− , the peak is completely separated from other products, thereby eliminating potential deconvolution errors. As shown in Fig. 5 a, b, after 3 minutes of etching, these decomposition products in LPSC 25d , where the decomposition is more pronounced, are effectively removed, as confirmed by XPS analysis. Similarly, most reaction products in LPSC 5d were fully removed within the same short etching time (Supplementary Fig. 10). As a result, a significant decrease in relative intensity of decomposition products compared to the PS 4 3− was observed, clearly indicating the formation of a reaction product-enriched surface layer (Fig. 5 c, d). Subsequent etching for 40 minutes removed oxide-related phases from the etched LPSC 25d surface, confirming the termination of decomposition (Supplementary Fig. 11). Given that decomposition products were substantially or completely removed within a short etching period for both LPSC 5d and LPSC 25d , it is evident that the reacted phases are concentrated primarily at the topmost surface. Therefore, the degradation mechanism is predominantly driven by surface-dependent reactions characterized by the entire sequence of steps leading to phase separation rather than downward anionic migration into the deeper atomic layer. Overall surface degradation mechanism As previously discussed, we propose a detailed decomposition mechanism to explain the dry room incompatibility of the LPSC. Initial H 2 O adsorption onto the surface induced the formation of an H–S bond, simultaneously weakening the P–S bond in the neighboring PS 4 tetrahedron. This process facilitates subsequent S-O substitutions, driven by energetic favorability, as indicated by the incorporation of oxygen into the decomposition products at the LPSC surface, identified through XPS analysis. Therefore, based on this fundamental reason, preventing parasitic gas adsorption through surface engineering and introduction of H 2 S absorbents or stabilizing sulfur bonding through compositional modifications could effectively mitigate material deterioration, aligning well with previous research approaches 13–20,22−26 . Further understanding of the degradation mechanism beyond the initial stages is crucial, given that the reaction intensifies with increased exposure time. Atomic-scale simulations demonstrated sequential reactions, where oxygen substitutions at exposed sulfur sites and their diffusion into subsurface atomic layers through polyhedral rotations led to further degradation from the topmost to the deeper layers. Thus, polyhedral rotation plays a pivotal role in altering surface properties by increasing the degree of oxygen involvement. Moreover, the formation of an O-rich, Li 6 PO 5 Cl-like surface and subsequent phase separations into smaller decomposition products due to inherent phase instability and moisture sensitivity resulted in an inhomogeneous surface layer. This structural collapse prevents the formation of a uniform surface layer that could protect underlying pristine atomic layers, causing continued degradation with prolonged exposure. Conclusively, our findings highlight the necessity for caution when handling sulfide SSEs under industrial dry-room conditions. The overall degradation mechanism involves (Fig. 6 ): (i) H 2 O adsorptions, (ii) S-O substitutions, (iii) polyhedral rotations, (iv) formation of O-enriched surface, and (v) volume-shrinking phase separations that result in the continuous surface decomposition. Based on these mechanistic insights, we propose novel material design strategies to improve dry-room compatibility for sulfide-based electrolytes containing PS 4 tetrahedra. In addition to the previous approaches, compositional designs that inhibit the polyhedral rotations of the O-substituted tetrahedra can be effective. Additionally, selecting appropriate dopants that promote decomposition products with sufficient atomic volumes could lead to surfaces serving as stable protective layers. Even after the occurrence of phase separations, decomposition byproducts with large atomic volumes could physically inhibit further gas adsorption and subsequent degradation without volume-shrinking phase separations. Conclusion In summary, this study elucidated the detailed mechanism underlying moisture-induced surface degradation of the sulfide-based argyrodite Li 6 PS 5 Cl (LPSC) under dry-room conditions. By combining first-principles calculations and experimental XPS depth-profiling analyses, we successfully identified the comprehensive sequence of surface reactions: (i) H 2 O adsorption, (ii) thermodynamically favored S-O substitutions, (iii) polyhedral rotations driving subsurface oxygen migration, (iv) formation of an unstable O-rich Li 6 PO 5 Cl-like surface, and (v) volume-shrinking phase separations leading to continuous surface decomposition. These reactions explain the observed 36% drop in ionic conductivity after short-term dry-room exposure, highlighting the critical vulnerability of LPSC in practical operating environments. Our mechanistic insights reveal three pivotal factors behind the continuous degradation process: initial H 2 O-induced weakening of the P–S bond, subsequent polyhedral rotations facilitating deeper oxygen involvement, and the inhomogeneous surface layer driven by the generation of smaller decomposition products. Based on these findings, we propose targeted strategies for enhancing dry-room compatibility of sulfide solid-state electrolytes. Specifically, beyond established methods such as inhibiting parasitic gas adsorption through surface engineering and stabilizing sulfur bonding via compositional modifications, we suggest designing compositions that prevent polyhedral rotations of O-substituted tetrahedra. Additionally, selecting dopant elements capable of forming decomposition products with sufficiently large atomic volumes could create stable protective surface layers. Such decomposition byproducts with larger volumes would physically inhibit further gas adsorption and subsequent degradation, thus significantly improving the compatibility of sulfide electrolytes within industrial dry-room environments. Methods Computational details All computational calculations were performed through the Vienna Ab initio Simulation Package (VASP) 37 . The generalized gradient approximation (GGA) with the Perdew-Burke-Ernzerhof (PBE) functional was applied 38 . An energy cutoff of a plane-wave basis set was 520 eV. A 4×4×4 k-point grid for the LPSC structure and a 3×3×1 k-point grid for the slab were used for the sampling of the Brillouin zone. Ionic positions of all structures were fully relaxed until the interatomic force was less than 0.02 eV Å −1 . The cubic bulk structure of the LPSC was prepared from the Materials Project database 39 with a space group of F4̅3m, and the slab structure of the material with the (001) surface was applied in the simulations. Precomputed energies and volumes from the Materials Project database were employed in theoretical reaction energy calculations in Supplementary Table 1 and Eq. (7), and comparison of atomic volumes in Fig. 4 d. The materials and their corresponding material identifiers as follows: Li 6 PS 5 Cl: mp-985592, LiCl: mp-22905, Li 3 PO 4 : mp-13725, LiOH: mp-23856, Li 2 CO 3 : mp-3054, and Li 2 SO 4 : mp-4556. The energies of the H 2 O, H 2 S, CO 2 , and O 2 molecules for reaction energies were calculated with computational input parameters generated by MPRelaxSet in the Python Materials Genomics (Pymatgen) library 40 . The energy of O 2 at 0 K was adjusted by a 1.36 eV energy shift suggested by Ceder et al. 41 and an enthalpy difference compared to the value at 298.15 K in NIST-JANAF Thermochemical tables to correct the GGA binding energy 42 . During the adsorption calculations, the DFT-D3 method with Becke-Johnson damping function was used to apply van der Waals correction 43 . The energy of H 2 O was recalculated based on this method and the same computational parameters as those for slab structures. Additionally, for E sub calculations, the oxygen chemical potential (–4.309 eV) was derived from the energy of the O 2 molecule obtained by changing into the computational settings consistent with parameters for the slab structures. The sulfur chemical potential (–4.134 eV) was calculated using bulk sulfur in accordance with the sulfur-rich condition at the surface of the LPSC. Identical input parameters with the slab calculations were applied. ICOHP calculations were carried out by using the Local Orbital Basis Suite Towards Electronic-Structure Reconstruction (LOBSTER) software 44–47 . Kinetic approaches for movements of substituted oxygen and surface reactions in slab structures were observed through AIMD simulations. They were performed with a 1×1×1 k-point grid and the NVT ensemble using a Nose-Hoover thermostat 48 . Post analyses for the atomic displacement along the c-axis were performed by employing the Pymatgen library and smoothed using a Savitzky-Golay filter. Material characterization Argyrodite materials, including pristine and dry room-exposed samples, were provided by Samsung SDI. For a structural characterization, XRD was carried out with a Rigaku with a scan rate of 1° min − 1 in a 5–70° 2θ range. Ionic conductivities of the samples were measured by conducting EIS analysis with a Solartron Analytical after pelletizing the solid electrolytes. All of the analysis data was given by Samsung SDI. XPS depth profiling XPS depth profiling was carried out with a K-Alpha (Thermo Fisher Scientific) for depth analysis. A monochromatic Al Kα (1486.6 eV) was used as an X-ray beam source, and the spot size of the X-ray was 400 µm. An etching area was 4 mm², and the acceleration voltages of Ar + ion etching beams were tested between 0.5 kV and 1.0 kV. We set binding energy values of materials in the S 2p spectra as PS 4 3− (a pristine phase of the LPSC)-161.6 eV, P 2 S n -163.5 eV, Li 2 S-160.3 eV, SO n 2− (sulfites and sulfates)-over than 166.0 eV, and in the P 2p spectra as PS 4 3− -132.2 eV, P 2 S n -133.1 eV, PO 4 3− -134.0 eV 49–52 . Binding energy scales were calibrated with the C 1s peak at 285.0 eV for the initial surface detection. During the depth analysis, since no change was observed after the phase decomposition, the 198.7 eV peak in the Cl 2p spectra was employed for the calibration of the binding energy 53,54 . To prevent unnecessary air exposure, we used a vacuum transfer module, one of the options of the K-Alpha (Thermo Fisher Scientific). It enabled samples not to be exposed to the air from the Ar-filled glove box to the analysis chamber of the spectroscopy. Declarations Ethics declarations The authors declare no competing interests. Acknowledgements This research was funded by Samsung SDI. The computational work was supported by the Supercomputing Center/Korea Institute of Science and Technology Information with supercomputing resources, including technical support (KSC-2024-CRE-0463). The depth experiment was supported by UNIST UCRF. Data Availability The data that support the findings of this study are available from the corresponding author upon reasonable request. Contributions Y.-S.K., J.-D.L., and D.-H.S. planned the project. D.-H.S. supervised all aspects of the research. 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Additional Declarations The authors declare no competing interests. Supplementary Files SupplementaryInformation.docx Moisture-induced surface degradation mechanism of argyrodite Li 6 PS 5 Cl under dry-room conditions Cite Share Download PDF Status: Posted Version 1 posted You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. We do this by developing innovative software and high quality services for the global research community. Our growing team is made up of researchers and industry professionals working together to solve the most critical problems facing scientific publishing. Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-7583174","acceptedTermsAndConditions":true,"allowDirectSubmit":true,"archivedVersions":[],"articleType":"Research Article","associatedPublications":[],"authors":[{"id":513408672,"identity":"1a4a2da3-1dad-4039-99e3-8fcc6fddf17c","order_by":0,"name":"Yoon-Seong Kim","email":"","orcid":"","institution":"Korea Advanced Institute of Science and Technology (KAIST)","correspondingAuthor":false,"prefix":"","firstName":"Yoon-Seong","middleName":"","lastName":"Kim","suffix":""},{"id":513409020,"identity":"6bf61949-7bac-49a8-b62a-82c4c3006f5f","order_by":1,"name":"Dong-Hwa 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(KATECH)","correspondingAuthor":false,"prefix":"","firstName":"Jeong-Doo","middleName":"","lastName":"Yi","suffix":""},{"id":513409022,"identity":"eab5a664-24a6-46b0-b663-45e54dff1b96","order_by":3,"name":"Sihyeon Sung","email":"","orcid":"","institution":"Samsung SDI","correspondingAuthor":false,"prefix":"","firstName":"Sihyeon","middleName":"","lastName":"Sung","suffix":""}],"badges":[],"createdAt":"2025-09-10 13:08:00","currentVersionCode":1,"declarations":{"humanSubjects":false,"vertebrateSubjects":false,"conflictsOfInterestStatement":false,"humanSubjectEthicalGuidelines":false,"humanSubjectConsent":false,"humanSubjectClinicalTrial":false,"humanSubjectCaseReport":false,"vertebrateSubjectEthicalGuidelines":false},"doi":"10.21203/rs.3.rs-7583174/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-7583174/v1","draftVersion":[],"editorialEvents":[],"editorialNote":"","failedWorkflow":false,"files":[{"id":91206706,"identity":"80eb21b3-de70-4ba7-a625-a9d4f9140734","added_by":"auto","created_at":"2025-09-12 16:49:58","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":580763,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eLi\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e6\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003ePS\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e5\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003eCl (LPSC) structures and theoretical S-O substitution energies (E\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003esub\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003e) depending on the sulfur atomic site.\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003ea\u003c/strong\u003e\u0026nbsp;The initial bulk structure (green-Li, purple-P, yellow-S, cyanite-Cl). \u003cstrong\u003eb\u003c/strong\u003e The (001) surface slab structures viewed from different angles. \u003cstrong\u003ec\u003c/strong\u003e E\u003csub\u003esub\u003c/sub\u003e difference plot between the non-bonding (Wyckoff 4d) site and P–S bonding (Wyckoff 16e) site.\u003c/p\u003e","description":"","filename":"floatimage1.png","url":"https://assets-eu.researchsquare.com/files/rs-7583174/v1/cb8c024e496d4dd49afeef38.png"},{"id":91207409,"identity":"d02b5b13-ba4e-4baf-8397-07982e9a7dc2","added_by":"auto","created_at":"2025-09-12 16:57:59","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":567829,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eTheoretical results of the H\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e2\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003eO adsorption on the LPSC surface.\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003ea\u003c/strong\u003e\u0026nbsp;Illustration for H\u003csub\u003e2\u003c/sub\u003eO adsorptions on two sulfur atomic sites, including the adsorption energy (E\u003csub\u003eads\u003c/sub\u003e) and bond properties of P–S and H–S bonds (green-Li, purple-P, yellow-S, cyanite-Cl, red-O, pink-H). \u003cstrong\u003eb\u003c/strong\u003e Evolution of the H–S bonding on the Wyckoff 4d site after the H\u003csub\u003e2\u003c/sub\u003eO adsorption on the surface of the O-substituted LPSC (2\u003csup\u003end\u003c/sup\u003e layer-substituted structure), including the integrated crystal orbital Hamiltonian population (ICOHP) values. The isosurface level, expressing charge density for H–S bonding, is set to be 0.03 in (\u003cstrong\u003ea\u003c/strong\u003e, \u003cstrong\u003eb\u003c/strong\u003e). \u003cstrong\u003ec\u003c/strong\u003e Plot for comparing the sum of ICOHP values for Li–S bonds at Wyckoff 4d site among the pristine LPSC, O-substituted LPSC (LPSCO), and H\u003csub\u003e2\u003c/sub\u003eO adsorption on LPSCO (H\u003csub\u003e2\u003c/sub\u003eO-LPSCO).\u003c/p\u003e","description":"","filename":"floatimage2.png","url":"https://assets-eu.researchsquare.com/files/rs-7583174/v1/d60ff860ba678dd3a6e3021f.png"},{"id":91206707,"identity":"cfa0822b-d427-4e02-b02b-7634c5ae3914","added_by":"auto","created_at":"2025-09-12 16:49:59","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":604315,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eComputational simulations for the demonstration of oxygen-substitution mechanisms.\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eE\u003csub\u003esub\u003c/sub\u003e of \u003cstrong\u003ea\u003c/strong\u003e an atomic substitution from the 1\u003csup\u003est\u003c/sup\u003e to the 3\u003csup\u003erd\u003c/sup\u003e S atomic layer and \u003cstrong\u003eb\u003c/strong\u003e mechanisms including additional oxygen-substitution and penetration at each reaction order during sequential substitutions. Plots of the displacement along the c-axis and structural changes through ab initio molecular dynamics (AIMD) simulation for \u003cstrong\u003ec\u003c/strong\u003e, \u003cstrong\u003ed\u003c/strong\u003e the 1\u003csup\u003est\u003c/sup\u003e and 2\u003csup\u003end\u003c/sup\u003e layer substituted structure, and \u003cstrong\u003ee\u003c/strong\u003e, \u003cstrong\u003ef\u003c/strong\u003e the additionally substituted structure involving the atomic substitution at the 3\u003csup\u003erd\u003c/sup\u003e layer.\u003c/p\u003e","description":"","filename":"floatimage3.png","url":"https://assets-eu.researchsquare.com/files/rs-7583174/v1/3c08ea270a06112f8459d97b.png"},{"id":91206714,"identity":"f183a2bd-dee6-4733-a549-c8bbb17a4b55","added_by":"auto","created_at":"2025-09-12 16:49:59","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":342695,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eDecomposition reaction pathway following the complete surface substitution.\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003ea\u003c/strong\u003e E\u003csub\u003esub\u003c/sub\u003e of atomic substitution from the 3\u003csup\u003erd\u003c/sup\u003e to the 5\u003csup\u003eth\u003c/sup\u003e S atomic layer in pristine LPSC structure. \u003cstrong\u003eb\u003c/strong\u003e ICOHP plot for P–S, P–O, Li–S, and broken P–S bonds related to the 3\u003csup\u003erd\u003c/sup\u003e through 5\u003csup\u003eth\u003c/sup\u003e atomic layers of the structure, completely O-substituted at the first three atomic layers (LPSCO\u003csub\u003e123\u003c/sub\u003e). \u003cstrong\u003ec\u003c/strong\u003e LPSCO\u003csub\u003e123\u003c/sub\u003e structure after AIMD simulation at 500 K for 100 ps. \u003cstrong\u003ed\u003c/strong\u003e The graph for the atomic volume of anticipated decomposition products compared to the pristine LPSC. \u003cstrong\u003ee\u003c/strong\u003e Schematic representation for the volume shrinkage of the reacted surface.\u003c/p\u003e","description":"","filename":"floatimage4.png","url":"https://assets-eu.researchsquare.com/files/rs-7583174/v1/a9ca10391fbfdf6f060f827b.png"},{"id":91206713,"identity":"e9f2d9ce-1c7f-424a-8442-eea4dcc0d725","added_by":"auto","created_at":"2025-09-12 16:49:59","extension":"png","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":288485,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eX-ray photoelectron spectroscopy (XPS) depth analysis of LPSC\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e25d\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003e.\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe change in decomposition phase compositions of LPSC\u003csub\u003e25d\u003c/sub\u003e after 3 minutes etching with a l.0 kV beam in \u003cstrong\u003ea\u003c/strong\u003e S 2p spectra and \u003cstrong\u003eb\u003c/strong\u003e P 2p spectra. The difference of relative bond intensity between the pre-etched and the 3-minute etched LPSC\u003csub\u003e25d\u003c/sub\u003e in \u003cstrong\u003ec\u003c/strong\u003e S 2p spectra and \u003cstrong\u003ed\u003c/strong\u003e P 2p spectra.\u003c/p\u003e","description":"","filename":"floatimage5.png","url":"https://assets-eu.researchsquare.com/files/rs-7583174/v1/6db3181ed2b2caf4ef9c137b.png"},{"id":91206509,"identity":"56983900-dd80-4478-9dff-4c4e787a4182","added_by":"auto","created_at":"2025-09-12 16:41:59","extension":"png","order_by":6,"title":"Figure 6","display":"","copyAsset":false,"role":"figure","size":320395,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eMechanistic illustration for the H\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e2\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003eO-induced surface degradation processes and proposition of the material design principles for dry-room stable sulfides grounded in the mechanism.\u003c/strong\u003e\u003c/p\u003e","description":"","filename":"floatimage6.png","url":"https://assets-eu.researchsquare.com/files/rs-7583174/v1/1b4ed82c2d3ee86697a9fa59.png"},{"id":91332404,"identity":"a12dea52-143d-421d-a6de-bd30580a1ae6","added_by":"auto","created_at":"2025-09-15 11:11:30","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":3604928,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-7583174/v1/e0fe6fd7-e0d3-4373-bfb2-f512419984ed.pdf"},{"id":91206507,"identity":"d983e25a-3c61-4891-ac30-701914c2f403","added_by":"auto","created_at":"2025-09-12 16:41:58","extension":"docx","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":2212394,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eMoisture-induced surface degradation mechanism of argyrodite Li\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e6\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003ePS\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e5\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003eCl under dry-room conditions\u003c/strong\u003e\u003c/p\u003e","description":"","filename":"SupplementaryInformation.docx","url":"https://assets-eu.researchsquare.com/files/rs-7583174/v1/6a9b42e65fb7ba1a9f370ef3.docx"}],"financialInterests":"The authors declare no competing interests.","formattedTitle":"\u003cp\u003e\u003cstrong\u003eMoisture-induced surface degradation mechanism of argyrodite Li\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e6\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003ePS\u003c/strong\u003e\u003csub\u003e\u003cstrong\u003e5\u003c/strong\u003e\u003c/sub\u003e\u003cstrong\u003eCl under dry-room conditions\u003c/strong\u003e\u003c/p\u003e","fulltext":[{"header":"Introduction","content":"\u003cp\u003eAll-solid-state batteries have recently attracted significant attention due to their superior thermal stability and potential high energy density. In contrast, conventional Li-ion batteries (LIBs) employing liquid electrolytes face enduring challenges, such as safety risks associated with flammable organic solvents and limited energy density due to the use of carbon-based anodes like graphite\u003csup\u003e1,2\u003c/sup\u003e. To address these challenges, inorganic solid-state electrolytes (SSEs) have emerged as promising alternatives, offering enhanced thermal stability and enabling the use of lithium metal as an anode\u003csup\u003e3\u003c/sup\u003e. Consequently, extensive research efforts have been devoted to discovering suitable solid electrolyte materials, leading to the development of diverse material classes with distinct structural and compositional characteristics.\u003c/p\u003e\u003cp\u003eAmong various candidates, sulfide-based inorganic SSEs have emerged as particularly promising due to their high ionic conductivity and high ductility, enabled by the polarizable nature of their anionic frameworks. Several sulfide electrolytes, including Li\u003csub\u003e7\u003c/sub\u003eP\u003csub\u003e3\u003c/sub\u003eS\u003csub\u003e11\u003c/sub\u003e, Li\u003csub\u003e10\u003c/sub\u003eGeP\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003e12\u003c/sub\u003e (LGPS), and the argyrodite compound Li\u003csub\u003e7\u0026thinsp;\u0026minus;\u0026thinsp;a\u003c/sub\u003ePS\u003csub\u003e6\u0026thinsp;\u0026minus;\u0026thinsp;a\u003c/sub\u003eX\u003csub\u003ea\u003c/sub\u003e (X\u0026thinsp;=\u0026thinsp;Cl, Br, I), have demonstrated high ionic conductivities exceeding 10 mS cm\u003csup\u003e\u0026minus;\u0026thinsp;1 4\u0026ndash;6\u003c/sup\u003e. Particularly among the argyrodite compounds, the thio-antimonate iodide argyrodites, Li\u003csub\u003e6\u0026thinsp;+\u0026thinsp;x\u003c/sub\u003eM\u003csub\u003ex\u003c/sub\u003eSb\u003csub\u003e1\u0026minus;x\u003c/sub\u003eS\u003csub\u003e5\u003c/sub\u003eI (M\u0026thinsp;=\u0026thinsp;Si, Ge, Sn), have exhibited exceptional ionic conductivity of 24 mS cm\u003csup\u003e\u0026minus;\u0026thinsp;1 7\u003c/sup\u003e. Moreover, the inherent ductility of sulfide SSEs helps reduce interfacial and grain boundary resistance. For example, the experimental measurements indicate that Li\u003csub\u003e6\u003c/sub\u003ePS\u003csub\u003e5\u003c/sub\u003eCl, a prototypical argyrodite, has a low elastic modulus (28.0\u0026thinsp;\u0026plusmn;\u0026thinsp;1.8 GPa), corroborated by theoretical values of bulk modulus (28.7 GPa) and Young\u0026rsquo;s modulus (22.1 GPa), facilitating intimate electrode-electrolyte contact under the relatively low fabrication and stack pressures\u003csup\u003e8\u0026ndash;10\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eDespite their desirable properties, the poor moisture stability of sulfide SSEs, which leads to the release of toxic H\u003csub\u003e2\u003c/sub\u003eS gas and chemical degradation, has impeded their practical application\u003csup\u003e11\u003c/sup\u003e. To prevent the generation of H\u003csub\u003e2\u003c/sub\u003eS gas, an inert atmosphere is required to avoid reactions between sulfides and moisture, which in turn increases the cost of the necessary infrastructure. Thus, improving compatibility under tightly controlled humidity conditions, such as those in dry rooms, is crucial for the industrial implementation of sulfide-based SSEs. However, even within these controlled environments, they have exhibited limited stability, as indicated by a decrease in ionic conductivity after 24 hours of exposure in the dry room (dew point \u0026minus;\u0026thinsp;45 ℃)\u003csup\u003e12\u003c/sup\u003e. To address moisture-induced degradation, various approaches have been proposed, including cationic or anionic substitutions based on the hard and soft acids and bases (HSAB) theory, surface engineering to inhibit H\u003csub\u003e2\u003c/sub\u003eO adsorption, and the use of H\u003csub\u003e2\u003c/sub\u003eS-absorbing additives\u003csup\u003e13\u0026ndash;26\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eAlthough the previous approaches have made meaningful progress in addressing the issue, developing new strategies for designing water-resistive materials remains challenging due to a limited understanding of detailed decomposition processes. Therefore, it is essential to establish a comprehensive understanding of surface reactions and their underlying mechanisms under dry room conditions. While various studies have attempted to chemically analyze these phenomena and identify their origin, they have largely focused on the surface-related changes following the air exposure or have been limited to theoretical investigations of the initial stage of the decomposition\u003csup\u003e27\u0026ndash;33\u003c/sup\u003e. As a result, they often provide only a general trend of the reactions in the subsequent steps. To overcome these limitations, further studies are needed to reveal the full degradation pathways, from reaction onset to the formation of decomposition products, at both atomic and nanoscale levels. Such insights will enable the development of more intuitive and diverse strategies for enhancing the stability of sulfide electrolytes against air and dry room environments.\u003c/p\u003e\u003cp\u003eIn this work, we systematically elucidate the surface degradation mechanism of the argyrodite Li\u003csub\u003e6\u003c/sub\u003ePS\u003csub\u003e5\u003c/sub\u003eCl (LPSC), a representative sulfide-based SSE, under controlled atmospheric conditions containing trace amounts of H\u003csub\u003e2\u003c/sub\u003eO. A combined approach of first-principles calculations and X-ray photoelectron spectroscopy (XPS) depth profiling was employed to gain mechanistic insight into the decomposition process. Thermodynamic and kinetic analyses were performed using first-principles calculations to investigate the reactive behavior of the LPSC surface at the atomic scale and revealed that degradation is initiated by H\u003csub\u003e2\u003c/sub\u003eO adsorption followed by S-O substitution, which locally alters the surface bonding environments. Energetic comparisons of possible substitution mechanisms further clarified the thermodynamic feasibility of these reactions and identified the preferred pathways. In addition, ab initio molecular dynamics (AIMD) simulations tracked the kinetic evolution of O-substituted species, showing that polyhedral rotations of PS\u003csub\u003e4\u003c/sub\u003e units facilitate deeper oxygen penetration and drive phase separation. This process ultimately leads to the formation of a porous surface layer due to volume contraction upon decomposition.\u003c/p\u003e\u003cp\u003eTo validate our theoretical predictions, XPS depth profiling was performed on LPSC samples after controlled exposure. The resulting phase distribution as a function of depth confirmed the presence of surface-localized reactions and distinct decomposition products compared to the bulk. The progressive changes in composition with increasing depth strongly support the proposed mechanism of volume-shrinking phase separation. Overall, this dual approach combining atomic-scale calculations with experimental depth-resolved analysis provides a comprehensive understanding of surface degradation in LPSC under dry room conditions. These insights offer a foundation for developing more robust sulfide electrolytes with enhanced environmental stability.\u003c/p\u003e"},{"header":"Results and discussion","content":"\u003cp\u003eSurface degradation in the dry room\u003c/p\u003e\n\u003cp\u003eAfter 3 days of exposure to a dry room environment (dew point \u0026minus;\u0026thinsp;60 to \u0026minus;\u0026thinsp;70 ℃), the electrochemical impedance spectroscopy (EIS) analysis results (Supplementary Fig.\u0026nbsp;1) show a significant decrease in ionic conductivity, from 3.13 mS cm\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e to 2.00 mS cm\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e (36.1% reduction). X-ray diffraction (XRD) patterns of the pristine and exposed samples of LPSC (Supplementary Fig.\u0026nbsp;2) exhibit notable peak broadening in the 2\u0026theta; range of 24\u0026deg;-33\u0026deg; after 7 days of exposure. This broadening likely indicates the formation of nanoscale decomposition products consistent with previous ex-situ TEM studies on argyrodite degradation under humidity-controlled conditions\u003csup\u003e28\u003c/sup\u003e.\u003c/p\u003e\n\u003cp\u003eXPS analysis (Supplementary Fig.\u0026nbsp;3) further supports this degradation behavior. Phase distribution at the surface of LPSC, LPSC\u003csub\u003e5d\u003c/sub\u003e, and LPSC\u003csub\u003e25d\u003c/sub\u003e (representing 5 and 25 days of dry room exposure, respectively) showed an increasing trend in the amount of decomposition products (P\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003en\u003c/sub\u003e, SO\u003csub\u003en\u003c/sub\u003e\u003csup\u003e2\u0026minus;\u003c/sup\u003e, PO\u003csub\u003e4\u003c/sub\u003e\u003csup\u003e3\u0026minus;\u003c/sup\u003e), accompanied by a decline in the pristine PS\u003csub\u003e4\u003c/sub\u003e\u003csup\u003e3\u0026minus;\u003c/sup\u003e signal. These results suggest that the nanoparticles observed in XRD are composed of the decomposition products formed during prolonged exposure. Collectively, the evidence points to surface degradation and nanoparticle formation as key contributors to the observed decline in ionic conductivity.\u003c/p\u003e\n\u003cp\u003eAdditionally, theoretical reaction energies supported the experimental surface degradation. Three representative reactions, previously reported for LPSC degradation, were considered\u003csup\u003e28\u003c/sup\u003e:\u003c/p\u003e\n\u003cdiv class=\"BlockQuote\"\u003e\n\u003cp\u003eLi\u003csub\u003e6\u003c/sub\u003ePS\u003csub\u003e5\u003c/sub\u003eCl\u0026thinsp;+\u0026thinsp;6H\u003csub\u003e2\u003c/sub\u003eO \u0026rarr; LiCl\u0026thinsp;+\u0026thinsp;Li\u003csub\u003e3\u003c/sub\u003ePO\u003csub\u003e4\u003c/sub\u003e\u0026thinsp;+\u0026thinsp;2LiOH\u0026thinsp;+\u0026thinsp;5H\u003csub\u003e2\u003c/sub\u003eS (1)\u003c/p\u003e\n\u003cp\u003e2LiOH\u0026thinsp;+\u0026thinsp;CO\u003csub\u003e2\u003c/sub\u003e \u0026rarr; Li\u003csub\u003e2\u003c/sub\u003eCO\u003csub\u003e3\u003c/sub\u003e\u0026thinsp;+\u0026thinsp;H\u003csub\u003e2\u003c/sub\u003eO (2)\u003c/p\u003e\n\u003cp\u003eLi\u003csub\u003e6\u003c/sub\u003ePS\u003csub\u003e5\u003c/sub\u003eCl\u0026thinsp;+\u0026thinsp;4H\u003csub\u003e2\u003c/sub\u003eO\u0026thinsp;+\u0026thinsp;2O\u003csub\u003e2\u003c/sub\u003e \u0026rarr; LiCl\u0026thinsp;+\u0026thinsp;Li\u003csub\u003e3\u003c/sub\u003ePO\u003csub\u003e4\u003c/sub\u003e\u0026thinsp;+\u0026thinsp;Li\u003csub\u003e2\u003c/sub\u003eSO\u003csub\u003e4\u003c/sub\u003e\u0026thinsp;+\u0026thinsp;4H\u003csub\u003e2\u003c/sub\u003eS (3)\u003c/p\u003e\n\u003c/div\u003e\n\u003cp\u003eAll three reactions exhibit negative density functional theory (DFT)-computed reaction energies (E\u003csub\u003ereaction\u003c/sub\u003e), confirming that they are thermodynamically favorable (Supplementary Table\u0026nbsp;1). Reaction (1), with an E\u003csub\u003ereaction\u003c/sub\u003e of \u0026minus;\u0026thinsp;6.222 eV, describes the initial hydrolysis of LPSC in the presence of water, forming LiOH and H\u003csub\u003e2\u003c/sub\u003eS among other products. The generated LiOH can further react with atmospheric CO\u003csub\u003e2\u003c/sub\u003e, as described in reaction (2), yielding Li\u003csub\u003e2\u003c/sub\u003eCO\u003csub\u003e3\u003c/sub\u003e with an E\u003csub\u003ereaction\u003c/sub\u003e of \u0026minus;\u0026thinsp;1.465 eV. Notably, reaction (3), which includes O\u003csub\u003e2\u003c/sub\u003e as a reactant, exhibits a significantly larger thermodynamic driving force (E\u003csub\u003ereaction\u003c/sub\u003e = \u0026minus;\u0026thinsp;18.991 eV) due to the formation of a stable Li\u003csub\u003e2\u003c/sub\u003eSO\u003csub\u003e4\u003c/sub\u003e. Theoretical trends suggest that the reactions promoted by H\u003csub\u003e2\u003c/sub\u003eO in conjunction with CO\u003csub\u003e2\u003c/sub\u003e and O\u003csub\u003e2\u003c/sub\u003e result in surface deterioration, which greatly affects the formation of decomposition phases and serves as a primary cause of the ionic conduction drop.\u003c/p\u003e\n\u003cp\u003eH\u003csub\u003e2\u003c/sub\u003eO adsorption on the LPSC surface\u003c/p\u003e\n\u003cp\u003eBuilding on the preceding findings, it is clear that dry room exposure triggers decomposition reactions. Therefore, a comprehensive investigation into the mechanisms underlying this phenomenon is essential. Given experimental limitations in capturing surface reactions that occur at extremely short timescales, DFT calculations were employed to achieve a detailed understanding.\u003c/p\u003e\n\u003cp\u003e\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eTo elucidate the decomposition mechanism from the onset of reaction, it is important to observe structural variations and local coordination geometries at the uppermost surface. As shown in Supplementary Fig.\u0026nbsp;3, a key feature of the deterioration is the oxygen involvement at the surface. Notably, based on reactions (1) and (3), the consistent formation of Li\u003csub\u003e3\u003c/sub\u003ePO\u003csub\u003e4\u003c/sub\u003e and H\u003csub\u003e2\u003c/sub\u003eS gas indicates that the exchange of two atomic species, S and O, fundamentally occurs at PS\u003csub\u003e4\u003c/sub\u003e tetrahedra. To verify the thermodynamic preference of this surface behavior, surface calculations were performed using the slab model of the LPSC (001) surface (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e1\u003c/span\u003ea, b, Supplementary Table\u0026nbsp;2)\u003csup\u003e32,34,35\u003c/sup\u003e. These calculations focused on the first five atomic layers, with the first layer defined as the layer closest to the surface (Supplementary Fig.\u0026nbsp;4).\u003c/p\u003e\n\u003cp\u003eAs shown in Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e1\u003c/span\u003ea, c, two sulfur atomic sites exist on the surface: a non-bonding site (Wyckoff 4d) surrounded by a lithium cage and a P\u0026ndash;S bonding site (Wyckoff 16e) within the PS\u003csub\u003e4\u003c/sub\u003e tetrahedron. We evaluated the thermodynamic driving force of the S-O substitution reaction energy (E\u003csub\u003esub\u003c/sub\u003e) at these two potential reaction sites using Eq.\u0026nbsp;(4):\u003c/p\u003e\n\u003cp\u003eE\u003csub\u003esub\u003c/sub\u003e = {(E\u003csub\u003esubstituted\u003c/sub\u003e \u0026ndash; E\u003csub\u003einitial\u003c/sub\u003e) \u0026ndash; n(\u0026micro;\u003csub\u003eO\u003c/sub\u003e \u0026ndash; \u0026micro;\u003csub\u003eS\u003c/sub\u003e)} / n (eV O atom\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e) (4)\u003c/p\u003e\n\u003cp\u003eIn Eq.\u0026nbsp;(4), E\u003csub\u003esubstituted\u003c/sub\u003e and E\u003csub\u003einitial\u003c/sub\u003e refer to the energies of various O-substituted and pre-substituted slab structures, respectively. The variable n refers to the number of substituted O atoms, and \u0026micro;\u003csub\u003eO\u003c/sub\u003e and \u0026micro;\u003csub\u003eS\u003c/sub\u003e are the chemical potential of O and S, respectively. As a result, substitution at both sites is energetically favorable, attributed to the increased stability from the more stable Li\u0026ndash;O and P\u0026ndash;O bonds. Between them, the P\u0026ndash;S bonding site (\u0026ndash;2.915 eV O atom\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e) exhibits a lower energy value than the non-bonding site (\u0026ndash;2.535 eV O atom-\u003csup\u003e1\u003c/sup\u003e), suggesting that it is the more reactive site for oxygen substitution.\u003c/p\u003e\n\u003cp\u003ePrior to the S-O substitution, the adsorption of H\u003csub\u003e2\u003c/sub\u003eO, a key initiator of the surface degradation, is regarded as the most likely event of the reaction, followed by structural changes. Although O\u003csub\u003e2\u003c/sub\u003e and CO\u003csub\u003e2\u003c/sub\u003e are involved in reaction processes (Supplementary Table\u0026nbsp;1), the principal origin of the decomposition is the H\u003csub\u003e2\u003c/sub\u003eO-induced initiation, as evidenced by discrepancies in reactivity depending on the humidity and atmospheric conditions\u003csup\u003e26,30\u003c/sup\u003e. Accordingly, the influence of initial H\u003csub\u003e2\u003c/sub\u003eO adsorption on the subsequent reactions was investigated to understand the substitution pathway by evaluating the adsorption energy (E\u003csub\u003eads\u003c/sub\u003e) and structural changes. The E\u003csub\u003eads\u003c/sub\u003e was calculated based on the energy difference before and after adsorption, as shown in Eq.\u0026nbsp;(5):\u003c/p\u003e\n\u003cp\u003eE\u003csub\u003eads\u003c/sub\u003e = E\u003csub\u003eLPSC+H2O\u003c/sub\u003e \u0026ndash; E\u003csub\u003eLPSC\u003c/sub\u003e \u0026ndash; E\u003csub\u003eH2O\u003c/sub\u003e (5)\u003c/p\u003e\n\u003cp\u003eAs shown in Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e2\u003c/span\u003ea, spontaneous adsorption occurs regardless of the neighboring sulfur site, although the energetically more favorable reaction occurs near the P\u0026ndash;S bonding site.\u003c/p\u003e\n\u003cp\u003eAfter adsorption of the single H\u003csub\u003e2\u003c/sub\u003eO molecule, changes in surface bonding environments, including weakened P\u0026ndash;S bond and newly formed H\u0026ndash;S bond, should precede the subsequent S-O substitution and release of H\u003csub\u003e2\u003c/sub\u003eS gas. To verify this, we examined the bond length and integrated crystal orbital Hamiltonian population (ICOHP), which provides insight into bond strength, where a higher ICOHP value indicates weaker bond strength (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e2\u003c/span\u003ea). Compared to adsorption near the non-bonding (4d) site, adsorption near the P\u0026ndash;S bonding (16e) site resulted in a significantly longer distance between P and S atoms (2.993 vs. 2.246 \u0026Aring;) of the neighboring PS\u003csub\u003e4\u003c/sub\u003e tetrahedron and a higher ICOHP value (\u0026ndash;0.481 vs. \u0026minus;\u0026thinsp;3.847 eV). This bond weakening originates from the newly formed H\u0026ndash;S bond (ICOHP: \u0026minus;\u0026thinsp;1.027 eV), which occurs between a neighboring PS\u003csub\u003e4\u003c/sub\u003e tetrahedron and the adsorbed H\u003csub\u003e2\u003c/sub\u003eO molecule. This weakened P\u0026ndash;S bonding facilitates subsequent S-O substitutions through oxygen sources available in the surrounding environment. Consequently, the substitution reaction ultimately promotes the generation and release of H\u003csub\u003e2\u003c/sub\u003eS gas. These results confirm the hypothesis that the initial S-O substitution, driven by adsorption of H\u003csub\u003e2\u003c/sub\u003eO near the P\u0026ndash;S bonding site, dominates the early stage of the degradation process.\u003c/p\u003e\n\u003cp\u003eSimilarly, Han et al. reported thermodynamic preference for anionic exchange on the P\u0026ndash;S bonding site and the generation of the sulfur-based gases, which were accelerated by the presence of preadsorbed water molecule\u003csup\u003e32\u003c/sup\u003e. This finding implies a spontaneous tendency toward multiple oxygen substitutions triggered by the continuous H\u003csub\u003e2\u003c/sub\u003eO adsorption. Given the favorable formation of the P\u0026ndash;O bond (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e1\u003c/span\u003ec, Supplementary Fig.\u0026nbsp;5), the adsorption effects on the second layer-substituted structure (O-substituted LPSC, LPSCO) were evaluated.\u003c/p\u003e\n\u003cp\u003eIn Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e2\u003c/span\u003eb, spontaneous surface adsorption (E\u003csub\u003eads\u003c/sub\u003e: \u0026minus;\u0026thinsp;1.766 eV) occurs with the H\u0026ndash;S bond formation (bond length: 2.008 \u0026Aring;, ICOHP: \u0026minus;\u0026thinsp;1.140 eV). Compared to the adsorption on the pristine LPSC surface, the shorter distance between H and S atoms (2.008 vs. 2.357 \u0026Aring;) and a lower ICOHP value (\u0026ndash;1.140 vs. \u0026minus;\u0026thinsp;0.342 eV) were observed during adsorption near the non-bonding site. To reveal the underlying reason for this result, a sum of ICOHP values for Li\u0026ndash;S bonds (Total ICOHP\u003csub\u003eLi\u0026ndash;S\u003c/sub\u003e) was evaluated (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e2\u003c/span\u003ec). The total ICOHP\u003csub\u003eLi\u0026ndash;S\u003c/sub\u003e value of LPSCO is higher than that of LPSC (\u0026ndash;4.903 vs. \u0026minus;\u0026thinsp;5.434 eV), while H\u003csub\u003e2\u003c/sub\u003eO-adsorbed LPSCO (H\u003csub\u003e2\u003c/sub\u003eO-LPSCO) exhibits an even higher value (\u0026ndash;4.440 eV). This weakening likely results from the structural3 instability caused by the distortion and Li\u0026ndash;S bond elongation owing to the substitution with smaller O\u003csup\u003e2\u0026minus;\u003c/sup\u003e (ionic radius: 140 pm) compared to S\u003csup\u003e2\u0026minus;\u003c/sup\u003e (ionic radius: 184 pm), as demonstrated in Supplementary Fig.\u0026nbsp;6. Sano et al. showed that LPSC exposed to atmosphere containing trace amounts of H\u003csub\u003e2\u003c/sub\u003eO (dew point \u0026minus;\u0026thinsp;20 ℃) and oxygen decomposed into sulfonates, while maintaining the PS\u003csub\u003e4\u003c/sub\u003e\u003csup\u003e3\u0026minus;\u003c/sup\u003e amount equivalent to the case under the same condition but without O\u003csub\u003e2\u003c/sub\u003e\u003csup\u003e30\u003c/sup\u003e. Interpreting this within our framework, it appears that the moisture primarily drives initial surface degradation, subsequently facilitating surface P\u0026ndash;S bond weakening and further SO\u003csub\u003en\u003c/sub\u003e\u003csup\u003e2\u0026minus;\u003c/sup\u003e formation upon reaction with O\u003csub\u003e2\u003c/sub\u003e molecules. Consequently, as these reactions continuously progress, the surface will become entirely O-substituted regardless of the sulfur atomic sites.\u003c/p\u003e\n\u003cp\u003eS-O substitution mechanism\u0026nbsp;\u003c/p\u003e\n\u003cp\u003eDuring the subsequent stages of decomposition at the surface, two possible scenarios of multiple S-O substitutions can occur: the penetration of substituted O species into neighboring atomic sites, and additional surface substitutions at the exposed sulfur sites. As illustrated in Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e3\u003c/span\u003ea, we calculated E\u003csub\u003esub\u003c/sub\u003e for single-atomic substitution, which showed that E\u003csub\u003esub\u003c/sub\u003e became increasingly negative with increasing depth, indicating a thermodynamic driving force favoring penetration. Furthermore, we identified the most energetically favorable reaction pathway by comparing the E\u003csub\u003esub\u003c/sub\u003e of two possible mechanisms at each reaction order grounded in an assumption of sequential substitutions (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e3\u003c/span\u003eb). They consisted of the additional substitution taking place at the 1st or 2nd atomic layer, and the migration of the substituted O to an internal region. Unlike the situation in which only additional substitution could occur in the first reaction order, two mechanisms proceeded competitively from the second order. Given that a more negative E\u003csub\u003esub\u003c/sub\u003e indicates greater energetic preference, the additional substitution is preferred in the second reaction order (\u0026ndash;3.023 vs. \u0026minus;\u0026thinsp;0.143 eV O atom\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e), which continues until the sixth order, referring to entire O substitutions at the exposed sulfur atoms. Based on the thermodynamic favorability, we can expect that the penetration is preceded by the complete series of O substitutions at the 1st and 2nd layers (LPSCO\u003csub\u003e12\u003c/sub\u003e), despite the presence of the driving force for the penetration.\u003c/p\u003e\n\u003cp\u003eAlong with the spontaneous substitutions occurring at the topmost surface, we carried out AIMD simulations of the LPSCO\u003csub\u003e12\u003c/sub\u003e slab structure to confirm how penetration proceeds. Surprisingly, Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e3\u003c/span\u003ec captures a positional reversal along the c-axis (perpendicular to the surface) between underlying sulfur and upper oxygen (O1) atoms in the PS\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e2\u003c/sub\u003e tetrahedron. A snapshot after 30 ps of simulation (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e3\u003c/span\u003ed) reveals polyhedral rotation. We attribute this result to the anionic size effect, wherein the sulfur anion migrates toward the vacuum to relieve structural stress induced by its larger ionic radius than that of O\u003csup\u003e2\u0026minus;\u003c/sup\u003e. As a result, the sulfur initially positioned in the lower layer migrates to the top surface and undergoes continuous substitution as previously confirmed by the greater thermodynamic stability of P\u0026ndash;O bond formation through successive anionic exchanges.\u003c/p\u003e\n\u003cp\u003eFollowing the tetrahedral rotation and subsequent substitutions, further simulations were performed to examine structures involving oxygen substitution extending to the third atomic layer (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e3\u003c/span\u003ee, f). Here, the sulfur atoms diffuse along the positive c-axis (toward the topmost surface), although the rotation requires more simulation time to reach the crossing-point with respect to the vertical height. This delay can be explained by the oxygen chemical potential gradient, consistent with Le Chatelier\u0026rsquo;s principle. Consequently, the final exposed sulfur species are expected to undergo additional substitutions, which leads to the formation of an O-substituted surface extending down to the third atomic layer (LPSCO\u003csub\u003e123\u003c/sub\u003e).\u003c/p\u003e\n\u003cp\u003eFurther decomposition during the continuous exposure\u003c/p\u003e\n\u003cp\u003eIf the fully O-substituted surface functioned as a protective layer against H\u003csub\u003e2\u003c/sub\u003eO molecules, the amount of decomposition products would not continuously increase with exposure time. However, XPS analysis (Supplementary Fig.\u0026nbsp;3) reveals the opposite trend, indicating that further investigation into the underlying reason for the continuous surface degradation is required.\u003c/p\u003e\n\u003cp\u003eOne plausible mechanism is the penetration of the substituted O atom in the third layer of LPSCO\u003csub\u003e123\u003c/sub\u003e. After the depletion of exposed sulfur species through O substitutions, polyhedral rotations of the O-substituted tetrahedron caused the subsequent substitutions at sulfur atoms even in the subsurface during the preceding sequential reactions. If there had been no reappearance of sulfur species on the surface, the oxygen substitution would have been terminated. Thus, given that the presence of sulfur is essential for the progression of the degradation, further downward diffusion of oxygen species from the third layer to the fourth or fifth layer, inducing an exchange with internal sulfur, can be one of the reaction pathways during the continuous surface decomposition. However, as shown in Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e4\u003c/span\u003ea, more positive E\u003csub\u003esub\u003c/sub\u003e values of the atomic substitution at the fourth (\u0026ndash;1.767 eV) and fifth layers (\u0026ndash;2.919 eV) compared to the one of the third layer (\u0026ndash;3.058 eV) indicate the existence of only a weak thermodynamic driving force for this penetration from the pristine structure.\u003c/p\u003e\n\u003cp\u003eMoreover, since penetration requires the cleavage of stable P\u0026ndash;O bonds and the subsequent substitution with the internal sulfur source, we carefully evaluated the feasibility of this coupled reaction using several analyses. Among possible sulfur sites within the interior layers, three distinct atomic sites were considered: i) P\u0026ndash;S bonding (16e) sites, ii) non-bonding (4d) sites, and iii) a broken P\u0026ndash;S bonding site generated after the formation of LPSCO\u003csub\u003e123\u003c/sub\u003e (Supplementary Fig.\u0026nbsp;7). In Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e4\u003c/span\u003eb, the ICOHP values of each bond relevant to oxygen and sulfur sites from the third to the fifth layer (P\u0026ndash;O: substituted O, P\u0026ndash;S: P\u0026ndash;S bonding site, Li\u0026ndash;S: non\u0026ndash;bonding site, broken P\u0026ndash;S: broken P\u0026ndash;S bonding site) are evaluated. Among various bond types, P\u0026ndash;O bonds exhibit the lowest values (\u0026ndash;7.943 to \u0026minus;\u0026thinsp;8.160 eV), followed in increasing order by P\u0026ndash;S bonds (\u0026ndash;5.793 to \u0026minus;\u0026thinsp;6.188 eV), Li\u0026ndash;S bonds (\u0026ndash;0.759 to \u0026minus;\u0026thinsp;1.151 eV), and the broken P\u0026ndash;S bond (\u0026ndash;0.517 eV). Since both P\u0026ndash;O and P\u0026ndash;S bonds are intrinsically stable, bond breakages followed by the transformation between these bonding interactions are kinetically unfavorable without any reaction initiators from an external source, such as the adsorption of gas molecules on the surface.\u003c/p\u003e\n\u003cp\u003eAdditionally, the thermodynamic favorability of S-O substitutions at sulfurs of non-bonding and broken P\u0026ndash;S bonding sites that exhibited weaker bond strength was evaluated. In Supplementary Fig.\u0026nbsp;8, the structural instability derived from the higher structural energies of S-O exchanged structures compared to the LPSCO\u003csub\u003e123\u003c/sub\u003e is observed, which indicates the absence of a thermodynamic driving force favoring penetration. Furthermore, to verify the non-existence of oxygen migration into deeper layers from a kinetic perspective, AIMD simulations were conducted at 500 K for 100 ps (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e4\u003c/span\u003ec). The result shows that O atoms do not exhibit a downward diffusion tendency: instead, the O-enriched surface is maintained without further reaction.\u003c/p\u003e\n\u003cp\u003eTherefore, the continuous surface decomposition must proceed through an alternative pathway. A plausible mechanism is the formation of an inhomogeneous surface layer as a result of degradation. Following sufficient progression of the oxygen substitution, the Li\u003csub\u003e6\u003c/sub\u003ePS\u003csub\u003e5\u0026thinsp;\u0026minus;\u0026thinsp;x\u003c/sub\u003eClO\u003csub\u003ex\u003c/sub\u003e\u0026rsquo;s O-rich surface will undergo structural collapse due to inherent instability induced by lattice mismatch between surface Li\u003csub\u003e6\u003c/sub\u003ePO\u003csub\u003e5\u003c/sub\u003eCl and internal LPSC phases. To quantify this instability, we calculated the surface energy (E\u003csub\u003esurface\u003c/sub\u003e) according to Eq.\u0026nbsp;(6):\u003c/p\u003e\n\u003cp\u003eE\u003csub\u003esurface\u003c/sub\u003e = (E\u003csub\u003eslab\u003c/sub\u003e \u0026ndash; N * E\u003csub\u003ebulk\u003c/sub\u003e) / 2A (6)\u003c/p\u003e\n\u003cp\u003ewhere E\u003csub\u003eslab\u003c/sub\u003e and E\u003csub\u003ebulk\u003c/sub\u003e are the energies of slab and bulk structures, respectively, N represents the number of bulk formula units contained within the slab structure, and A is the surface area. Using this equation, the reacted O-rich surface (LPSCO\u003csub\u003e123\u003c/sub\u003e) exhibits significantly higher surface energy (0.213 eV \u0026Aring;\u003csup\u003e\u0026minus;2\u003c/sup\u003e) compared to the pristine LPSC surface (0.014 eV \u0026Aring;\u003csup\u003e\u0026minus;2\u003c/sup\u003e). This notable increase in surface energy indicates that the reacted O-rich surface is inherently unstable.\u003c/p\u003e\n\u003cp\u003eFurthermore, Li\u003csub\u003e6\u003c/sub\u003ePO\u003csub\u003e5\u003c/sub\u003eCl is thermodynamically unstable against hydrolysis (Eq.\u0026nbsp;(7)), with a negative DFT-computed reaction energy of \u0026minus;\u0026thinsp;2.033 eV:\u003c/p\u003e\n\u003cp\u003eLi\u003csub\u003e6\u003c/sub\u003ePO\u003csub\u003e5\u003c/sub\u003eCl\u0026thinsp;+\u0026thinsp;H\u003csub\u003e2\u003c/sub\u003eO \u0026rarr; LiCl\u0026thinsp;+\u0026thinsp;Li\u003csub\u003e3\u003c/sub\u003ePO\u003csub\u003e4\u003c/sub\u003e\u0026thinsp;+\u0026thinsp;2LiOH (7)\u003c/p\u003e\n\u003cp\u003eConsequently, additional decomposition reactions on the destabilized O-rich surface are expected. Multiple decomposition products are generated, corresponding to the reactions presented earlier (equations (1), (2), and (3)), all of which possess smaller atomic volumes than the pristine LPSC (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e4\u003c/span\u003ed). Considering the volumetric changes upon phase deterioration, the shrinkage of decomposition products leads to the formation of void\u003csup\u003e36\u003c/sup\u003e. Therefore, as schematically illustrated in Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e4\u003c/span\u003ee, the volume reduction of the oxidized surface does not adequately protect underlying pristine atomic layers, resulting in continuous surface decomposition and failure to serve as a stable protective layer.\u003c/p\u003e\n\u003cp\u003eSurface with concentrated decomposition products\u003c/p\u003e\n\u003cp\u003eIn the absence of oxygen penetration into deeper layers, a surface degradation is driven by the formation of a surface layer enriched with decomposition products resulting from volume-shrinking phase separation. To experimentally validate this scenario, a comparative analysis of the distribution of decomposition phases between the topmost surface and deeper layers was performed by using XPS depth profiling. Prior to this analysis, different intensities of etching beams were tested to identify optimal analytical conditions. If its intensity is excessively strong, unintended decomposition of the pristine PS\u003csub\u003e4\u003c/sub\u003e\u003csup\u003e3\u0026minus;\u003c/sup\u003e phase could occur, resulting in misleading results. As demonstrated in Supplementary Fig.\u0026nbsp;9, both the 0.5 kV and 1.0 kV etching beams yielded similar phase distributions to the non-etched LPSC sample. Thus, the 0.5 kV beam was applied for LPSC\u003csub\u003e5d\u003c/sub\u003e analysis, which had a relatively mild degree of decomposition, while the 1.0 kV beam was selected for LPSC\u003csub\u003e25d\u003c/sub\u003e analysis to ensure time- and cost-efficiency.\u003c/p\u003e\n\n\u003cp\u003eDuring the XPS depth analysis, the presence of SO\u003csub\u003en\u003c/sub\u003e\u003csup\u003e2\u0026minus;\u003c/sup\u003e (including sulfites and sulfates) and PO\u003csub\u003e4\u003c/sub\u003e\u003csup\u003e3\u0026minus;\u003c/sup\u003e is a clear indicator for assessing changes in phase distribution, as these phases are absent in pristine LPSC (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e5\u003c/span\u003ea, b, Supplementary Fig.\u0026nbsp;9). Especially for SO\u003csub\u003en\u003c/sub\u003e\u003csup\u003e2\u0026minus;\u003c/sup\u003e, the peak is completely separated from other products, thereby eliminating potential deconvolution errors. As shown in Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e5\u003c/span\u003ea, b, after 3 minutes of etching, these decomposition products in LPSC\u003csub\u003e25d\u003c/sub\u003e, where the decomposition is more pronounced, are effectively removed, as confirmed by XPS analysis. Similarly, most reaction products in LPSC\u003csub\u003e5d\u003c/sub\u003e were fully removed within the same short etching time (Supplementary Fig.\u0026nbsp;10). As a result, a significant decrease in relative intensity of decomposition products compared to the PS\u003csub\u003e4\u003c/sub\u003e\u003csup\u003e3\u0026minus;\u003c/sup\u003e was observed, clearly indicating the formation of a reaction product-enriched surface layer (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e5\u003c/span\u003ec, d).\u003c/p\u003e\n\u003cp\u003eSubsequent etching for 40 minutes removed oxide-related phases from the etched LPSC\u003csub\u003e25d\u003c/sub\u003e surface, confirming the termination of decomposition (Supplementary Fig.\u0026nbsp;11). Given that decomposition products were substantially or completely removed within a short etching period for both LPSC\u003csub\u003e5d\u003c/sub\u003e and LPSC\u003csub\u003e25d\u003c/sub\u003e, it is evident that the reacted phases are concentrated primarily at the topmost surface. Therefore, the degradation mechanism is predominantly driven by surface-dependent reactions characterized by the entire sequence of steps leading to phase separation rather than downward anionic migration into the deeper atomic layer.\u003c/p\u003e\n\u003cp\u003eOverall surface degradation mechanism\u003c/p\u003e\n\u003cp\u003eAs previously discussed, we propose a detailed decomposition mechanism to explain the dry room incompatibility of the LPSC. Initial H\u003csub\u003e2\u003c/sub\u003eO adsorption onto the surface induced the formation of an H\u0026ndash;S bond, simultaneously weakening the P\u0026ndash;S bond in the neighboring PS\u003csub\u003e4\u003c/sub\u003e tetrahedron. This process facilitates subsequent S-O substitutions, driven by energetic favorability, as indicated by the incorporation of oxygen into the decomposition products at the LPSC surface, identified through XPS analysis. Therefore, based on this fundamental reason, preventing parasitic gas adsorption through surface engineering and introduction of H\u003csub\u003e2\u003c/sub\u003eS absorbents or stabilizing sulfur bonding through compositional modifications could effectively mitigate material deterioration, aligning well with previous research approaches\u003csup\u003e13\u0026ndash;20,22\u0026minus;26\u003c/sup\u003e.\u003c/p\u003e\n\u003cp\u003eFurther understanding of the degradation mechanism beyond the initial stages is crucial, given that the reaction intensifies with increased exposure time. Atomic-scale simulations demonstrated sequential reactions, where oxygen substitutions at exposed sulfur sites and their diffusion into subsurface atomic layers through polyhedral rotations led to further degradation from the topmost to the deeper layers. Thus, polyhedral rotation plays a pivotal role in altering surface properties by increasing the degree of oxygen involvement. Moreover, the formation of an O-rich, Li\u003csub\u003e6\u003c/sub\u003ePO\u003csub\u003e5\u003c/sub\u003eCl-like surface and subsequent phase separations into smaller decomposition products due to inherent phase instability and moisture sensitivity resulted in an inhomogeneous surface layer. This structural collapse prevents the formation of a uniform surface layer that could protect underlying pristine atomic layers, causing continued degradation with prolonged exposure.\u003c/p\u003e\n\u003cp\u003eConclusively, our findings highlight the necessity for caution when handling sulfide SSEs under industrial dry-room conditions. The overall degradation mechanism involves (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e6\u003c/span\u003e): (i) H\u003csub\u003e2\u003c/sub\u003eO adsorptions, (ii) S-O substitutions, (iii) polyhedral rotations, (iv) formation of O-enriched surface, and (v) volume-shrinking phase separations that result in the continuous surface decomposition. Based on these mechanistic insights, we propose novel material design strategies to improve dry-room compatibility for sulfide-based electrolytes containing PS\u003csub\u003e4\u003c/sub\u003e tetrahedra. In addition to the previous approaches, compositional designs that inhibit the polyhedral rotations of the O-substituted tetrahedra can be effective. Additionally, selecting appropriate dopants that promote decomposition products with sufficient atomic volumes could lead to surfaces serving as stable protective layers. Even after the occurrence of phase separations, decomposition byproducts with large atomic volumes could physically inhibit further gas adsorption and subsequent degradation without volume-shrinking phase separations.\u003c/p\u003e"},{"header":"Conclusion","content":"\u003cp\u003eIn summary, this study elucidated the detailed mechanism underlying moisture-induced surface degradation of the sulfide-based argyrodite Li\u003csub\u003e6\u003c/sub\u003ePS\u003csub\u003e5\u003c/sub\u003eCl (LPSC) under dry-room conditions. By combining first-principles calculations and experimental XPS depth-profiling analyses, we successfully identified the comprehensive sequence of surface reactions: (i) H\u003csub\u003e2\u003c/sub\u003eO adsorption, (ii) thermodynamically favored S-O substitutions, (iii) polyhedral rotations driving subsurface oxygen migration, (iv) formation of an unstable O-rich Li\u003csub\u003e6\u003c/sub\u003ePO\u003csub\u003e5\u003c/sub\u003eCl-like surface, and (v) volume-shrinking phase separations leading to continuous surface decomposition. These reactions explain the observed 36% drop in ionic conductivity after short-term dry-room exposure, highlighting the critical vulnerability of LPSC in practical operating environments.\u003c/p\u003e\u003cp\u003eOur mechanistic insights reveal three pivotal factors behind the continuous degradation process: initial H\u003csub\u003e2\u003c/sub\u003eO-induced weakening of the P\u0026ndash;S bond, subsequent polyhedral rotations facilitating deeper oxygen involvement, and the inhomogeneous surface layer driven by the generation of smaller decomposition products. Based on these findings, we propose targeted strategies for enhancing dry-room compatibility of sulfide solid-state electrolytes. Specifically, beyond established methods such as inhibiting parasitic gas adsorption through surface engineering and stabilizing sulfur bonding via compositional modifications, we suggest designing compositions that prevent polyhedral rotations of O-substituted tetrahedra. Additionally, selecting dopant elements capable of forming decomposition products with sufficiently large atomic volumes could create stable protective surface layers. Such decomposition byproducts with larger volumes would physically inhibit further gas adsorption and subsequent degradation, thus significantly improving the compatibility of sulfide electrolytes within industrial dry-room environments.\u003c/p\u003e"},{"header":"Methods","content":"\u003cp\u003eComputational details\u003c/p\u003e\u003cp\u003eAll computational calculations were performed through the Vienna Ab initio Simulation Package (VASP)\u003csup\u003e37\u003c/sup\u003e. The generalized gradient approximation (GGA) with the Perdew-Burke-Ernzerhof (PBE) functional was applied\u003csup\u003e38\u003c/sup\u003e. An energy cutoff of a plane-wave basis set was 520 eV. A 4\u0026times;4\u0026times;4 k-point grid for the LPSC structure and a 3\u0026times;3\u0026times;1 k-point grid for the slab were used for the sampling of the Brillouin zone. Ionic positions of all structures were fully relaxed until the interatomic force was less than 0.02 eV \u0026Aring;\u003csup\u003e\u0026minus;1\u003c/sup\u003e. The cubic bulk structure of the LPSC was prepared from the Materials Project database\u003csup\u003e39\u003c/sup\u003e with a space group of F4̅3m, and the slab structure of the material with the (001) surface was applied in the simulations. Precomputed energies and volumes from the Materials Project database were employed in theoretical reaction energy calculations in Supplementary Table\u0026nbsp;1 and Eq.\u0026nbsp;(7), and comparison of atomic volumes in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ed. The materials and their corresponding material identifiers as follows: Li\u003csub\u003e6\u003c/sub\u003ePS\u003csub\u003e5\u003c/sub\u003eCl: mp-985592, LiCl: mp-22905, Li\u003csub\u003e3\u003c/sub\u003ePO\u003csub\u003e4\u003c/sub\u003e: mp-13725, LiOH: mp-23856, Li\u003csub\u003e2\u003c/sub\u003eCO\u003csub\u003e3\u003c/sub\u003e: mp-3054, and Li\u003csub\u003e2\u003c/sub\u003eSO\u003csub\u003e4\u003c/sub\u003e: mp-4556. The energies of the H\u003csub\u003e2\u003c/sub\u003eO, H\u003csub\u003e2\u003c/sub\u003eS, CO\u003csub\u003e2\u003c/sub\u003e, and O\u003csub\u003e2\u003c/sub\u003e molecules for reaction energies were calculated with computational input parameters generated by MPRelaxSet in the Python Materials Genomics (Pymatgen) library\u003csup\u003e40\u003c/sup\u003e. The energy of O\u003csub\u003e2\u003c/sub\u003e at 0 K was adjusted by a 1.36 eV energy shift suggested by Ceder et al.\u003csup\u003e41\u003c/sup\u003e and an enthalpy difference compared to the value at 298.15 K in NIST-JANAF Thermochemical tables to correct the GGA binding energy\u003csup\u003e42\u003c/sup\u003e.\u003c/p\u003e\u003cp\u003eDuring the adsorption calculations, the DFT-D3 method with Becke-Johnson damping function was used to apply van der Waals correction\u003csup\u003e43\u003c/sup\u003e. The energy of H\u003csub\u003e2\u003c/sub\u003eO was recalculated based on this method and the same computational parameters as those for slab structures. Additionally, for E\u003csub\u003esub\u003c/sub\u003e calculations, the oxygen chemical potential (\u0026ndash;4.309 eV) was derived from the energy of the O\u003csub\u003e2\u003c/sub\u003e molecule obtained by changing into the computational settings consistent with parameters for the slab structures. The sulfur chemical potential (\u0026ndash;4.134 eV) was calculated using bulk sulfur in accordance with the sulfur-rich condition at the surface of the LPSC. Identical input parameters with the slab calculations were applied.\u003c/p\u003e\u003cp\u003eICOHP calculations were carried out by using the Local Orbital Basis Suite Towards Electronic-Structure Reconstruction (LOBSTER) software\u003csup\u003e44\u0026ndash;47\u003c/sup\u003e. Kinetic approaches for movements of substituted oxygen and surface reactions in slab structures were observed through AIMD simulations. They were performed with a 1\u0026times;1\u0026times;1 k-point grid and the NVT ensemble using a Nose-Hoover thermostat\u003csup\u003e48\u003c/sup\u003e. Post analyses for the atomic displacement along the c-axis were performed by employing the Pymatgen library and smoothed using a Savitzky-Golay filter.\u003c/p\u003e\u003cp\u003eMaterial characterization\u003c/p\u003e\u003cp\u003eArgyrodite materials, including pristine and dry room-exposed samples, were provided by Samsung SDI. For a structural characterization, XRD was carried out with a Rigaku with a scan rate of 1\u0026deg; min\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e in a 5\u0026ndash;70\u0026deg; 2θ range. Ionic conductivities of the samples were measured by conducting EIS analysis with a Solartron Analytical after pelletizing the solid electrolytes. All of the analysis data was given by Samsung SDI.\u003c/p\u003e\u003cp\u003eXPS depth profiling\u003c/p\u003e\u003cp\u003eXPS depth profiling was carried out with a K-Alpha (Thermo Fisher Scientific) for depth analysis. A monochromatic Al Kα (1486.6 eV) was used as an X-ray beam source, and the spot size of the X-ray was 400 \u0026micro;m. An etching area was 4 mm\u0026sup2;, and the acceleration voltages of Ar\u003csup\u003e+\u003c/sup\u003e ion etching beams were tested between 0.5 kV and 1.0 kV. We set binding energy values of materials in the S 2p spectra as PS\u003csub\u003e4\u003c/sub\u003e\u003csup\u003e3\u0026minus;\u003c/sup\u003e (a pristine phase of the LPSC)-161.6 eV, P\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003en\u003c/sub\u003e-163.5 eV, Li\u003csub\u003e2\u003c/sub\u003eS-160.3 eV, SO\u003csub\u003en\u003c/sub\u003e\u003csup\u003e2\u0026minus;\u003c/sup\u003e (sulfites and sulfates)-over than 166.0 eV, and in the P 2p spectra as PS\u003csub\u003e4\u003c/sub\u003e\u003csup\u003e3\u0026minus;\u003c/sup\u003e-132.2 eV, P\u003csub\u003e2\u003c/sub\u003eS\u003csub\u003en\u003c/sub\u003e-133.1 eV, PO\u003csub\u003e4\u003c/sub\u003e\u003csup\u003e3\u0026minus;\u003c/sup\u003e-134.0 eV\u003csup\u003e49\u0026ndash;52\u003c/sup\u003e. Binding energy scales were calibrated with the C 1s peak at 285.0 eV for the initial surface detection. During the depth analysis, since no change was observed after the phase decomposition, the 198.7 eV peak in the Cl 2p spectra was employed for the calibration of the binding energy\u003csup\u003e53,54\u003c/sup\u003e. To prevent unnecessary air exposure, we used a vacuum transfer module, one of the options of the K-Alpha (Thermo Fisher Scientific). It enabled samples not to be exposed to the air from the Ar-filled glove box to the analysis chamber of the spectroscopy.\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003ch2\u003eEthics declarations\u003c/h2\u003e\u003cp\u003eThe authors declare no competing interests.\u003c/p\u003e\u003c/p\u003e\u003ch2\u003eAcknowledgements\u003c/h2\u003e\u003cp\u003eThis research was funded by Samsung SDI. The computational work was supported by the Supercomputing Center/Korea Institute of Science and Technology Information with supercomputing resources, including technical support (KSC-2024-CRE-0463). The depth experiment was supported by UNIST UCRF.\u003c/p\u003e\n\u003ch3\u003eData Availability\u003c/h3\u003e\n\u003cp\u003eThe data that support the findings of this study are available from the corresponding author upon reasonable request.\u003c/p\u003e\n\u003cp\u003eContributions\u003c/p\u003e\n\u003cp\u003eY.-S.K., J.-D.L., and D.-H.S. planned the project. D.-H.S. supervised all aspects of the research. Y.-S.K. performed the overall experiments and first-principles calculations, with the interpretation of the results. J.-D.L. supported the preparation of samples. S.S. assisted in experimental works. Y.-S.K. and D.-H.S. wrote the manuscript.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\u003cli\u003e\u003cspan\u003e1. Goodenough, J. B. \u0026amp; Kim, Y. 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Narayanan, S. \u003cem\u003eet al.\u003c/em\u003e Effect of current density on the solid electrolyte interphase formation at the lithium∣ Li6PS5Cl interface. \u003cem\u003eNature Communications\u003c/em\u003e \u003cb\u003e13\u003c/b\u003e, 7237 (2022).\u003c/span\u003e\u003c/li\u003e\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":true,"hideJournal":true,"highlight":"","institution":"Korea Advanced Institute of Science and Technology","isAcceptedByJournal":false,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"researchsquare","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":true,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"/submission","title":"Research Square","twitterHandle":"researchsquare","acdcEnabled":true,"dfaEnabled":false,"editorialSystem":"","reportingPortfolio":"","inReviewEnabled":false,"inReviewRevisionsEnabled":true},"keywords":"sulfide argyrodites, all-solid-state batteries, dry room compatibility, degradation mechanism","lastPublishedDoi":"10.21203/rs.3.rs-7583174/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-7583174/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eArgyrodite Li₆PS₅Cl (LPSC) possesses both high Li-ion conductivity (~\u0026thinsp;10 mS cm⁻\u0026sup1; at room temperature) and mechanical softness, positioning it as a flagship solid electrolyte for next-generation all-solid-state batteries (ASSBs). However, even trace amounts of moisture in industrial dry rooms (dew-point \u0026minus;\u0026thinsp;60 to \u0026minus;\u0026thinsp;70\u0026deg;C) rapidly degrade its surface, diminishing ionic transport and impeding scalable processes. Here, we elucidate the moisture-triggered surface degradation mechanism of LPSC under dry-room conditions by combining first-principles calculations with depth-profiling X-ray photoelectron spectroscopy analysis. The combined analysis reveals a five-step sequence: (i) H\u003csub\u003e2\u003c/sub\u003eO adsorptions on S-rich surface, (ii) P\u0026ndash;S bond weakening followed by thermodynamically favoured S-O substitutions, (iii) rotation of the O-substituted PS₄ tetrahedra that drives O migration into subsurface layers, (iv) formation of an O-rich Li₆PO₅Cl-like surface, and (v) volume-shrinking phase separation into LiCl, Li₃PO₄, Li₂SO₄, LiOH, and Li\u003csub\u003e2\u003c/sub\u003eCO\u003csub\u003e3\u003c/sub\u003e. The resulting porous, O-enriched layer fails to passivate the electrolyte, causing a 36% drop in ionic conductivity within three days. These mechanistic insights highlight polyhedral-rigidity tuning and moisture-blocking surface chemistries as complementary strategies for stabilizing thiophosphate electrolytes during practical cell fabrication.\u003c/p\u003e","manuscriptTitle":"Moisture-induced surface degradation mechanism of argyrodite Li6PS5Cl under dry-room conditions","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2025-09-12 16:41:54","doi":"10.21203/rs.3.rs-7583174/v1","editorialEvents":[{"type":"communityComments","content":0}],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"researchsquare","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":true,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"/submission","title":"Research Square","twitterHandle":"researchsquare","acdcEnabled":true,"dfaEnabled":false,"editorialSystem":"","reportingPortfolio":"","inReviewEnabled":false,"inReviewRevisionsEnabled":true}}],"origin":"","ownerIdentity":"1b07365a-f0a8-46be-8d0d-29fc1df2eeca","owner":[],"postedDate":"September 12th, 2025","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"posted","subjectAreas":[{"id":54540236,"name":"Materials Chemistry"},{"id":54540237,"name":"Computational Chemistry"},{"id":54540238,"name":"Materials Theory and Modeling"}],"tags":[],"updatedAt":"2025-09-12T16:41:54+00:00","versionOfRecord":[],"versionCreatedAt":"2025-09-12 16:41:54","video":"","vorDoi":"","vorDoiUrl":"","workflowStages":[]},"version":"v1","identity":"rs-7583174","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-7583174","identity":"rs-7583174","version":["v1"]},"buildId":"8U1c8b4HqxoKbykW_rLl7","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}