Zinc-based metal halide electrolytes for all-solid-state zinc-metal batteries

Domain: chemistry
Published: 2026-01-19
Editorial outcome: major_revision

Abstract
The advancement of electrolyte materials is fundamental to the advancement of batteries. While zinc-metal batteries offer compelling advantages in safety, cost, and environmental sustainability, their performance is constrained by conventional aqueous electrolytes. All-solid-state zinc-ion electrolytes, which can inherently address the issues associated with aqueous electrolytes, are promising but yet currently in their nascent stages. Here, we show a composition engineering strategy to design metal halide-based zinc-ion electrolytes, and two types of zinc-based hybrid metal halides, C4H12N2ZnBr4 and C4H12N2ZnCl4, are obtained. Both solid-state electrolytes exhibit good electronically insulating properties and low zinc-migration energy barriers. C4H12N2ZnBr4 demonstrates a high ionic conductivity of 2.9×10−4 S cm−1 at 25 °C and a wide electrochemical window of ~3.74 V, making it suitable for all-solid-state zinc-metal batteries. The C4H12N2ZnBr4 electrolyte also establishes a tight contact with the zinc electrode, thereby minimizing interface resistance for enhanced electrochemical performance. As a result, an all-solid-state Zn||I2 cell using C4H12N2ZnBr4 electrolyte achieves a high capacity retention of 234.5 mAh g−1 after 200 cycles at 0.2 mA cm−2, with a low capacity degradation rate of 0.056% per cycle. This work provides a design approach for metal halide-based solid electrolytes in zinc-metal batteries.

Introduction
In recent years, increasing renewable energy production has highlighted the imperative to advance battery technology, with a particular emphasis on alternatives beyond lithium-ion batteries1. Electrolyte materials, which serve as a key medium for ion transport between negative electrode and positive electrode, play a critical role in battery devices, and battery performance is heavily influenced by the properties of electrolytes. Among various new battery systems, aqueous zinc-metal batteries (ZMBs) become a hot topic due to their high volumetric capacity (5855 mAh·cm−3), inherent safety, cost-effectiveness, and environmental sustainability2,3. Nevertheless, the wide application of aqueous ZMBs has not been achieved due to the inherent high activity of aqueous electrolytes4–6, the narrow electrochemical stability window, and the dissolution of positive electrode materials7,8.

Solid-state electrolytes are promising alternatives to aqueous electrolytes due to their potential advantages of safety, energy density, and mechanical properties9. It is very exciting that the introduction of solid-state electrolytes can effectively address the issues of dendrite formation and gas generation associated with zinc negative electrodes in aqueous ZMBs. Inspired by the research of lithium-ion batteries, most studies on solid-state zinc-ion electrolytes are based on polymer electrolyte materials10,11. Unfortunately, the ionic conductivity of polymer electrolytes, which relies on the random thermal motion of zinc ions, is relatively low. Zinc ions presents significant challenges for solid-state migration due to two key properties: their high charge valence (+2) and relatively low electronegativity (1.65). The high charge density associated with the +2 state induces strong electrostatic interactions with the anion lattice, resulting in high migration barriers. Simultaneously, the low electronegativity promotes stronger bonding with anions (e.g., O2−), which further restricts mobility. This dual impediment makes zinc ion conduction particularly difficult to achieve in solid electrolytes. Recently, inorganic-based solid-state zinc-ion electrolytes have been developed, such as fluorine-doped mesoporous zinc sulfide ZnSyF1-x12 and a phosphorus sulfide ZnPS313. These inorganic-based solid-state zinc-ion electrolytes possess high ionic conductivity and can effectively conduct zinc ions, endowing their batteries with good cycling performance. However, the research of solid-state zinc-ion electrolytes is still in its early stages, thereby the quest for applicable solid-state electrolytes remains of significant importance.

In this work, we proposed a composition engineering strategy for designing metal halide-based all-solid-state zinc-ion electrolytes. Based on the theoretical analysis between lithium- and zinc-based ternary metal halides, two zinc-based hybrid metal halides, C4H12N2ZnBr4 and C4H12N2ZnCl4 (denoted as PipZnBr4 and PipZnCl4, respectively, where Pip is the divalent ion of piperazine), were theoretically studied. These two materials show good electronically insulating properties and low zinc-migration barriers and may have the potential to be all-solid-state zinc-ion electrolytes. PipZnBr4, in particular, was experimentally verified and found to have high ionic conductivity, low activation energy, and a wide electrochemical window. It also forms a close contact with the zinc electrode to reduce interface resistance for promising electrochemical performance.

Results
Metal halide materials, a large family of materials composed of metal and halide ions (like chloride or bromide), have been widely used in various fields14–17. There are successful examples of metal halide materials for lithium solid electrolytes18,19. These studies verify that metal halide materials show good structural stability, high ionic conductivity, wide voltage windows, and long cycle stability at high voltages. The attempt to introduce metal halides into ZMBs is of great significance, as it can not only enrich the types of all-solid-state zinc-ion electrolyte materials but also deepen our understanding of their intrinsic mechanisms. However, it was found that metal halide materials for solid-state lithium batteries cannot be directly applicable to ZMBs. Specifically, lithium (Li) is the 3rd element and belongs to alkali metals, while zinc (Zn) is the 30th element and belongs to the transition metals. As shown in Fig. 1a, b and Supplementary Fig. 1, in the arrangement of electrons outside nucleus, Li ions have significantly fewer electrons than Zn ions. When Li or Zn elements are combined with halogen elements, both Li and Zn can usually form tetrahedral structures and octahedral ligand structures (Fig. 1c, d). When the coordination environments of Li and Zn ions with halogens are either 4-coordinated or 6-coordinated (Li/Zn surrounded by 4 or 6 halide atoms), their ionic radii are notably similar. In a four-coordination condition, the radius of Li+ is 59 pm, and that of Zn2+ is 60 pm; in a six-coordination condition, the radius of Li+ is 76 pm and that of Zn2+ is 74 pm20. This similarity implies that they occupy similar space in the crystal. To understand why these structurally similar materials behave differently as electrolytes, a comparative analysis of the tetrahedral and octahedral structures of Li and Zn at the orbital level was performed to elucidate their similarities and differences. As shown in Supplementary Figs. 2−5 (Supplementary Data 1), when Li and Zn atoms bond with halogen atoms to form tetrahedral or octahedral structures, the nature of their interaction differs fundamentally. The electron orbitals of Li and halogen atoms exhibit clear separation, with minimal overlap between their electron clouds (Supplementary Figs. 2 and 3). This result means Li and halogen atoms retain a non-bonding state in their elemental orbitals, which suggests that the predominant binding mode between Li ions and halogens is ionic bonds. Conversely, when Zn ions coordinate with halogens to form tetrahedral or octahedral structures, their interaction exhibits stronger orbital overlap (Supplementary Figs. 4 and 5). This overlap electron cloud sharing generates distinct bonding and antibonding states, which means a covalent character in the Zn-halogen bond. To quantify the interaction, an in-depth analysis of the crystal orbital bond index (COBI)21 was conducted for Li−Cl and Zn−Cl as representative cases (Supplementary Data 2). As illustrated in Supplementary Fig. 6, the COBI values for 4-coordinated and 6-coordinated LiCl are 0.15 and 0.09, respectively, suggesting a strong ionic character. In contrast, the COBI values for 4-coordinated and 6-coordinated ZnCl2 are 0.22 and 0.18 (Supplementary Fig. 7), indicating a stronger covalent interaction between Zn and the halogens. These results suggest that the Li-based halides characterized by localized ionic bonding, which could facilitate Li ion liberation, render them potential candidates for solid-state Li electrolytes. In contrast, the relatively strong covalent bonding interactions in Zn-based halides may hinder Zn ion migration. Consequently, traditional Zn-based halides with a Zn-halogen structure may not be able to effectively conduct Zn ions, thereby rendering them unsuitable as solid-state Zn-ion electrolytes. This further suggests that the design structure of all-solid electrolytes in Zn batteries cannot be modeled after Li-ion halide solid electrolytes.

To validate the above point, the migration pathways of reported Li-ion solid-state halide electrolyte materials (Li-ion ternary halides)22–26 and Zn-ion ternary halides were calculated by using a bond valence site energy (BVSE) method27,28 (Supplementary Data 3). In the case of Li-ion solid electrolytes, β-Li2ZrCl6 serves as an example, as shown in Fig. 1e. It exhibits a significantly low migration energy barrier for a 2D ionic pathway, at only 0.47 eV, which is consistent with the previously reported study26. Supplementary Fig. 8 and Supplementary Table 1 have summarized the typically reported Li-ion solid-state electrolyte halides, which consistently demonstrate a pathway characterized by a relatively low migration energy barrier (0.47~0.57 eV). In contrast, the calculated migration energy barriers for Zn-ion ternary halide materials Cs2ZnCl4 exceed 1.23 eV (Fig. 1f). As further shown in Supplementary Fig. 9 and Supplementary Table 2, additional calculations suggest that the migration energy barriers of Zn-ion ternary halides (0.96~1.48 eV) are relatively higher than those of previously reported Li-ion solid-state electrolyte halides. These results indicate that the conventional Zn-ion ternary halides may not be suitable for solid-state Zn-ion electrolytes, and it is crucial to mitigate the strong interaction between Zn ions and halide anions. Subsequently, constructing a channel for Zn-ion migration is significant to obtain superior solid-state Zn-ion electrolytes.

We propose a composition engineering strategy that involves replacing the skeleton metal ions in Zn ion ternary halides with organics, aiming to construct potential Zn ion solid electrolytes. Figure 1g illustrates the specific methodology of this composition engineering, wherein the inorganic A-site cations in the original A2ZnX4 (A=Cs or Rb; X=halogen) are replaced with organic groups R (R=organic clusters) to form the RZnX4. Given that the original inorganic cations are spherical, their spatial positions tend to be relatively compact, making it difficult for these materials to form Zn ion channels. After component engineering design, the target material contains organic cationic groups, which exhibit a comparatively loose spatial configuration. This may facilitate transmission possibilities for Zn ions. In addition, the integration of organic cations may provide more abundant defect properties for materials, thereby enhancing the kinetics of Zn-ion transport. Based on the principle of charge balance, a series of possible cation organic clusters are chosen to combine with the [ZnX4]2− units to form soft Zn-based hybrid metal halides (Fig. 1h). Supplementary Fig. 10 shows the structures of the possible organic molecules and corresponding cation clusters.

After a series of experimental explorations based on the above strategy, two representative materials, i.e., PipZnBr4 and PipZnCl4, were synthesized through a facile solvothermal reaction. The parent compound halides were prepared by hydrothermal heating of commercially available zinc bromide (ZnBr2/ZnCl2) and piperazine in hydrobromic acid (HBr/HCl). The mixture solution was heated to 90 °C within 120 min, and kept for 2440 min, then cooling of the solution yielded soft co-crystalline halides. As shown in Fig. 2a and c, the crystals of both PipZnBr4 and PipZnCl4 belong to the same space group, P2221 (NO. 19), in the orthorhombic system, which is consistent with the previously reported results29 and the simulated X-ray diffraction (XRD) patterns from Cambridge Crystallographic Data Centre (CCDC) database (Supplementary Fig. 11). The computation of the chemical phase diagrams can provide additional evidence supporting the stability of these two materials. Given that the organic clusters effectively function as a cation in both materials, we consider the whole system as a three-component compound for defect calculations. The chemical phase diagrams of PipZnBr4 and PipZnCl4 were determined by the known competing phases in Supplementary Table 3 (Supplementary Data 4). Both PipZnBr4 and PipZnCl4 show trapezoidal thermochemical stability windows, indicating their good thermodynamic stability and would be synthesizable. (Fig. 2b and d). The chemical boundary points of PipZnBr4 are marked as A1, B1, C1, and D1, while those of PipZnCl4 are represented as A2, B2, C2, and D2. The calculated chemical potentials, that are, ΔμZn, ΔμBr, ΔμCl/ΔμPip, in the chemical phase diagrams are shown in the Supplementary Tables 4 and 5.

To evaluate the electronic conductivity of PipZnBr4 and PipZnCl4 materials, the electronic structures of both materials were investigated systematically. For PipZnBr4, the Heyd–Scuseria–Ernzerhof (HSE) calculated band structure shows a large direct band gap of (5.37 eV) at the Γ point as shown in Fig. 2e, implying a poor electrical conductivity. The density of states (DOS) analysis (Fig. 2f) reveals that the valence band maximum (VBM) is dominated by nonbonding states of Br 4p orbitals, whereas the adjacent valence band region consists of antibonding states from Zn 3d and Br 4p hybridization. The conduction band minimum (CBM) primarily consists of the antibonding states of Zn 4s and Br 4p orbitals, as well as partial p orbitals of organic cations (Fig. 2f). Additionally, the local charge density can provide a more intuitive representation of the orbital contributions to support this interpretation. As shown in Fig. 2g, the local charge density at CBM shows a near tetrahedron cloud around a Zn atom and semi-spindle electric clouds around a Br atom, which intuitively illustrates the antibonding states of the Zn 3s and Br 4p orbitals. The local charge density at VBM is composed solely of spindle-like p orbital electric clouds on the Br atoms (Fig. 2h), corresponding to their non-bonding orbitals. For PipZnCl4, the higher electronegativity of Cl relative to Br lowers the energy of Cl 3p orbitals, which results in stronger Zn 4s–Cl 3p orbital hybridization. This effect enhances the splitting between bonding and anti-bonding states, significantly elevating the energy of the anti-bonding orbital at the CBM. Consequently, PipZnCl4 exhibits a larger band gap (6.31 eV) compared to PipZnBr4 (5.37 eV), as depicted in Fig. 2i. The DOS and local charge density of PipZnBr4 are very similar to those of PipZnCl4, except that the orbital contribution changes from Br 4p to Cl 3p (Fig. 2j−l). These results imply that both materials exhibit good stability and a high degree of electronic insulation, thereby suggesting their potential as solid-state Zn-ion electrolytes.

The defect properties and the Zn-ion migrations of PipZnBr4 and PipZnCl4 were calculated to evaluate their potential carrier mobility and ionic conductivity. Because the piperazine cluster is much larger than Zn and Br/Cl atoms, nine intrinsic point defects are considered, i.e., three vacancies (VPip, VZn, and VBr/VCl), four antisites (BrPip/ClPip, BrZn/ClZn, ZnBr/ZnCl, and ZnPip/ZnPip), and two interstitials (Zn, and Br/Cl) (Supplementary Data 5). The defect models used a 2×2×2 supercell structure (Supplementary Fig. 12). The formation enthalpies (ΔH) of the above-mentioned intrinsic defects as a function of Fermi level (EF) at each boundary point based on Fig. 2b and d are plotted in Fig. 3a–h, and the corresponding transition levels ε(q/q’) of acceptor-like and donor-like defects are summarized in Supplementary Figs. 13 and 14. Among the calculated nine intrinsic point defects, VPip, VBr/Cl, VZn, and Br/Cl i are the dominant defects, which determine the electrical properties of PipZnBr4 and PipZnCl4. Under the Zn-rich and Br/Cl-poor condition (Point A1 and A2), the primary donor defect is halogen vacancy (VBr/VCl), while the domain acceptor defects are VZn, VPip, and Br/Cl i, indicating complex defect chemistry in this condition. The equilibrium Fermi levels (EF,e) for PipZnBr4 and PipZnCl4 at Points A1/A2 are 3.61 eV and 4.27 eV above the VBM, respectively (Fig. 3a and e). As the chemical potential gradually shifts from Br-poor to Br-rich (from point A1 to D1), the EF,e of PipZnBr4 moves correspondingly from near the VBM towards the CBM (Fig. 3a−d). Similarly, for PipZnCl4, as the chemical potential changes from Cl-poor to Cl-rich (from point A2 to D2), its EF,e also has a similar shift trend (Fig. 3e−h). Crucially, no matter how the EF,e of PipZnBr4 and PipZnCl4 changes, it is far away from the band edge, and the corresponding hole density (Nh) or electron density (Ne) is very low (Supplementary Tables 6 and 7). These results further show the electronically insulating nature of these two materials. It should be noted that the dominant donor defects are always the halogen vacancy defects (VBr for PipZnBr4 and VCl for PipZnCl4), while the dominant acceptor defects change with different chemical potential. The dominant halogen vacancy defects suggest that both materials are susceptible to breaking the original Zn-halogen tetrahedral structure due to these defects, which may provide more space for Zn ion conduction.

To evaluate the Zn-ion mobility in PipZnBr4 and PipZnCl4, the Zn ion pathways for a Zn atom confined within two materials were investigated by the climbing image nudged elastic band (Cl-NEB) method30 (Supplementary Data 6). Exploration of migration pathways along all three axis directions revealed that the lowest energy barriers occur along the a-direction (Fig. 3i−l and Supplementary Figs. 15 and 16). As shown in Fig. 3i and k, the migration of the Zn-ion along the a-axis occurs almost in the direction of the calculated supercells, specifically through the gap between the [ZnX4]2− units and the organic cations. The corresponding energy barriers for Zn ion diffusion along the a-direction in PipZnBr4 and PipZnCl4 are 0.37 eV (Fig. 3j) and 0.39 eV (Fig. 3l), respectively, which means Zn ions can easily migrate in the two materials. The migration results of Zn-ion in PipZnBr4 and PipZnCl4 indicate that both materials exhibit good Zn-ion conductivity. The above defect properties and the Zn-ion migration calculation results suggest that both materials have good electronically insulating properties and Zn-ion conductivity, which shows great potential as superior solid-state Zn-ion electrolytes.

Based on the above theoretical calculation guidance, we use PipZnBr4 as a case study to further explore the practical application of Zn-ion electrolytes. Defect regulation was achieved by adding excess ZnBr2 compared to piperazine to increase the ionic conduction. Specifically, the feed molar ratio of piperazine to ZnBr2 was regulated from 1:1 to 1:5, corresponding denoted as PipZnBr4 (1:1) and PipZnBr4. As characterized by scanning electron microscopy (SEM) in Fig. 4a and Supplementary Fig. 17, PipZnBr4 shows irregular polygonal particles with a diameter of 10–80 μm. The XRD analysis of the prepared hybrid Zn halides was implemented to verify the crystal structure. As shown in Fig. 4b and Supplementary Fig. 18a, the experimentally obtained powder XRD patterns of PipZnBr4 correspond well with the simulated pattern (# CCDC PipZnBr4: 284070), especially PipZnBr4 (1:1), indicating excess ZnBr2 does not affect the crystal structure of compound PipZnBr429. The chemical structures of the as-synthesized PipZnBr4 sample were studied by Raman spectroscopy and Fourier transform infrared (FTIR) tests. As shown in Fig. 4c, compared with pure piperazine, the PipZnBr4 exhibits an N−H single peak (~816 cm−1) and a new peak of −NH2+ (~1540 cm−1), manifesting the formation of piperazine ions31. The inversion symmetry of C−H bond peaks (2700~3200 cm−1) also illustrates the formation of the compound PipZnBr432. The Raman peak of the prepared PipZnBr4 (1:1) has a similar situation (Supplementary Fig. 18b). Additionally, the presence of −NH2+ peak in FTIR spectra of PipZnBr4 and PipZnBr4 (1:1) verifies the formation of the soft PipZnBr4 (Supplementary Fig. 19). Significantly, the FTIR spectra also show that the addition of excess ZnBr2 does not impact the structure of the compound PipZnBr4. The Zn content (12.73 wt%) of the synthesized PipZnBr4 sample with excess ZnBr2 is only slightly increased compared to the PipZnBr4 (1:1) (11.99 wt%) by inductively coupled plasma (ICP) measurements (Supplementary Fig. 20), implying that it only exists a small number of vacancies to accelerate ion migration12. Additionally, the weight and XRD of PipZnBr4 remain almost unchanged when the temperature rises to 150 °C in the thermal gravimetric test (Supplementary Fig. 21). Even when the temperature increases to 300 °C, the weight of the sample can remain at 96.5 wt%, indicating the good thermal stability of PipZnBr4.

The synthesized PipZnBr4 sample with excess ZnBr2 was pressed into a 16 mm diameter pellet with a thickness of ~430 μm as a solid-state Zn-ion electrolyte for ZMBs (Fig. 4d). The ionic conductivity (σ) of the pressed pellet of PipZnBr4 was investigated at different temperatures with 5 °C intervals by electrochemical impedance spectroscopy (EIS) test. The σ value can be calculated according to the total resistance determined by the intercept on the real axis based on fitting Nyquist plots (Supplementary Fig. 22a−m). The temperature-dependent ionic conductivities of PipZnBr4, within the range of 0~60 °C, are depicted in Supplementary Fig. 22n. The elevated ion conductivity indicates that PipZnBr4 serves as a good Zn-ion conductor, aligning with prior computational predictions. Notably, the high σ values of 2.9 and 4.7×10−4 S cm−1 at 25 °C and 40 °C were observed for the PipZnBr4 pellet, indicating its potential suitability as an all-solid-state Zn-ion electrolyte. Its Arrhenius curve is plotted in Fig. 4e to calculate the activation energy (Ea) of PipZnBr4 electrolyte based on the Arrhenius equation. The result shows that PipZnBr4 has a small activation energy (0.13 eV), indicating its promising application in all-solid-state ZMBs. To verify the high ionic conductivity of PipZnBr4, the ionic conductivity of ZnBr2 has been assessed at 25 °C (Supplementary Fig. 23). Meanwhile, PipZnBr4 also exhibits a potential range from −1.46 to 2.28 V (Fig. 4f), which is much wider than the electrochemical window of 1.2 M zinc trifluoromethanesulfonate (Zn(OTf)2) aqueous electrolyte (2.02 V)33 and hydrogel electrolyte (2.13 V)34. More significantly, PipZnBr4 as an all-solid-state Zn-ion electrolyte can display a Zn-ion transfer number (tZn2+) of 0.61 (Fig. 4g), which effectively regulates Zn-ion flux. While this value suggests contributions from additional ionic migration pathways (e.g., Br− ions) in PipZnBr4 electrolytes, Zn-ion transport remains the dominant mechanism and the Zn-ion transfer number is comparable to that of other solid Zn-ion conductors35–37. To investigate the electrochemical process of the Zn electrode, chronoamperometry (CA) tests and voltage-time curves were evaluated. The current density of the PipZnBr4 electrolyte rises instantaneously and then quickly remains stable (Fig. 4h), implying a three-dimensional (3D) dominating Zn-ion diffusion. Additionally, the PipZnBr4-based Zn||Zn solid-state cell shows low initial nucleation overpotentials at 0.1, 0.5, and 1.0 mA cm−2 (Supplementary Fig. 24). This phenomenon suggests that Zn ions are prone to be reduced and nucleated at the active sites, resulting in a uniform and dense Zn deposition. To study the anti-corrosion property of PipZnBr4, the Tafel polarization curve was tested (Fig. 4i). The Zn electrode in the PipZnBr4 electrolyte delivers a lower corrosion potential (Ecorr = −851.4 mV vs. Zn2+/Zn) and a much lower corrosion current density (Icorr = 0.037 μA cm−2) than a 1.2 M Zn(OTf)2 aqueous solution33, implying the weak adverse reaction.

The impact of PipZnBr4 electrolyte on the Zn deposition behavior was first observed by SEM. As shown in Fig. 5a and Supplementary Fig. 25, the uniform and compact spherical deposition of Zn is exhibited in PipZnBr4 Zn-ion electrolyte after 20 cycles at 0.1 mA cm−2 and 0.1 mAh cm−2, suggesting good Zn deposition behavior. The spherical morphology and Young’s modulus of the Zn electrode after different cycles were further characterized by atomic force microscopy (AFM). The mean modulus value and average height of the Zn negative electrode following 20 cycles are 2.1 GPa and 0.9 μm (Fig. 5b), while the Zn foil after 50 cycles at 0.1 mA cm−2 and 0.1 mAh cm−2 shows a slightly larger modulus value (3.6 GPa) and average height (2.3 μm) (Supplementary Fig. 26). These results illustrate that the PipZnBr4 electrolyte induces a deposition behavior of densely compacted Zn grains. To characterize the Zn crystal orientation and crystal structure regularity, electron backscatter diffraction (EBSD) was implemented38. As shown in Fig. 5c, the Zn grains from Zn foil in the PipZnBr4 electrolyte after 20 cycles at 0.1 mA cm−2 and 0.1 mAh cm−2 exhibit irregularity, and the grain boundaries are interlaced, implying polycrystalline and random orientation characteristics39 (Supplementary Fig. 27). From the ex situ 3D confocal laser scanning microscopy (CLSM), the Zn negative electrodes in PipZnBr4 electrolyte after different cycles at 0.1 mA cm−2 and 0.1 mAh cm−2 also present uniform and compact Zn grain deposition (Fig. 5di−iii), which is consistent with the AFM results. The Zn negative electrode surface after 10 cycles at 0.1 mA cm−2 shows a larger surface roughness (2.8 μm), higher surface volume (1.9×107 μm3), and larger surface area (4.6×106 μm2) than those of the Zn negative electrode after 30 and 50 cycles (Supplementary Fig. 28 and Supplementary Table 8), confirming the notably even Zn deposition with the cycling. By contrast, the surface roughness (4.9 μm) is the largest for the Zn electrode after 20 cycles at 0.1 mA cm−2 and 0.1 mAh cm−2, which may be because the electrode is slightly less dense in the early deposition. Intuitively, the self-correlation function of the Zn electrode surface roughness after 50 cycles shows the region with an unclear boundary (Fig. 5div) due to the weak self-correlation, indirectly indicating uniform and dense Zn deposition34.

To further investigate the internal structure and non-destructive Zn deposition morphology, Computed tomography (CT) technology was implemented for the PipZnBr4-based solid-state Zn||Zn cell after 20 cycles at 0.1 mA cm−2 and 0.1 mAh cm−2. The PipZnBr4-based battery is divided into five parts to understand detailly the contact between the Zn electrode and electrolyte (Fig. 5e). Obviously, the CT of the 1-Zn electrode presents a smooth texture (Fig. 5e(1)), which suggests that the Zn foil here is not involved in the electrochemical reaction. Figure 5e(2) presents a uniform and flat interface between the Zn electrode and electrolyte, implying good interface contact. Further closer to the electrolyte, it is found that a small number of uneven pits appear at the interface (Fig. 5e(3)), which may be caused by the solid and particulate PipZnBr4 electrolyte. As shown in Fig. 5e(4), the CT of the solid electrolyte further verifies the uniform interface. To observe the interfacial contact of the entire battery system, the interface CT was also studied by dynamically recording the CT images of the cross-section of the solid-state Zn symmetrical cell (Supplementary Fig. 29). It is observed that the Zn electrodes and electrolyte maintain close contact after 20 cycles, which contributes to retaining the integrity of the electrolyte and restraining the occurrence of Zn dendrites and other side reactions. Due to the intimate contact at the Zn/PipZnBr4 electrolyte interface and the good reaction kinetics, the Zn surface exhibits uniform and compact Zn spherical deposition (Fig. 5e(5)), which is consistent with SEM and AFM results. The PipZnBr4-based Zn symmetrical cell was disassembled after the CT test, and then the properties of the electrolyte and the Zn electrodes were investigated. Compared with pure PipZnBr4 powder (Supplementary Fig. 17), the PipZnBr4 electrolyte after Zn plating/stripping at 0.1 mA cm−2 and 0.1 mAh cm−2 exhibits a smaller particle size (Supplementary Fig. 30), which may be because the high pressure used in the preparation of the all-solid-state Zn-ion electrolyte pellet causes the PipZnBr4 powder to crush. The XRD spectra of the electrolyte after Zn plating/stripping do not change significantly compared to pure PipZnBr4 powder (Supplementary Fig. 31). After 20 cycles at 0.1 mA cm−2 and 0.1 mAh cm−2, the peak positions of C1s, N1s, Br3d, ZnLMM, and Zn2p X-ray photoelectron spectroscopy (XPS) shift to a high energy level as a whole (Supplementary Fig. 32). This may be caused by the application of voltage to the cell during Zn plating/stripping, inducing environment change of the elements in the PipZnBr4 electrolyte, implying that the PipZnBr4 has good structural stability. Additionally, the surface composition of the Zn electrode after 20 cycles was characterized. Compared with the bare Zn foil, the electrode surface after Zn plating/stripping at 0.1 mA cm−2 and 0.1 mAh cm−2 shows some new peaks (Supplementary Fig. 33). The XPS test with Ar+ sputtering was further used to measure the interphase chemistry on the Zn electrode surface in all-solid-state Zn-ion electrolyte after 20 cycles at 0.1 mA cm−2 and 0.1 mAh cm−2 (Supplementary Fig. 34). It is found that the in situ formed solid electrolyte interphase (SEI) layer on Zn surface features an organic-rich outer layer and an inorganic-rich inner layer, which contributes to suppressing dendrite formation and promoting Zn-ion diffusion40.

To explore the change of element valence state in the PipZnBr4 electrolyte and its stability before and after Zn plating/stripping at 0.1 mA cm−2 and 0.1 mAh cm−2, X-ray absorption fine structure (XAFS) tests were executed. Figure 6a shows the normalized X-ray absorption near-edge structure (XANES) spectra at Zn K-edge for pure ZnBr2, Zn-foil, pure PipZnBr4 and PipZnBr4 after 20 cycles. Obviously, the Zn-foil demonstrates the lowest edge absorption energy in XANES, a characteristic primarily attributable to the inherent zero-valence state of the Zn-foil41. By contrast, the absorption edges of the remaining materials are essentially identical, suggesting the divalent of Zn in them. The Fourier transform spectra of the extended XAFS (EXAFS) curves in R-space show that the bond length of Zn−Br in PipZnBr4 is longer than that in pure ZnBr2 (Fig. 6b), which suggests a relatively weaker Zn−Br bond in PipZnBr4. Additionally, the results of the XANES (Fig. 6a), Fourier-transformed EXAFS (Fig. 6b), and the wavelet-transformed k3-weighted Zn K-edge EXAFS spectra (Fig. 6c and Supplementary Fig. 35) of PipZnBr4 suggest that the Zn-Br bond length remains unchanged before and after Zn plating/stripping at 0.1 mA cm−2 and 0.1 mAh cm−2. This implies that PipZnBr4 as a solid-state electrolyte exhibits robust electrochemical stability.

To further assess the interfacial and electrochemical stability of the Zn/PipZnBr4, Zn solid-state cells with a PipZnBr4 electrolyte were investigated systematically. Figure 6d shows the cycling performance of the Zn symmetric cell for over 1100 h at 0.1 mA cm−2. The corresponding polarization potential increases gradually with the increase of the Zn plating/stripping time, attributing to the high specific kinetic obstacle and the slow Zn plating/stripping42. When the current density is increased to 0.5 mA cm−2, the PipZnBr4-based Zn symmetric cell still works over 400 h (Supplementary Fig. 36a), implying the stable ionic migration. At an increased current density of 1.0 mA cm−2, long Zn plating/stripping process with voltage hysteresis (~860 mV) can still be attained in the symmetric Zn cell (Supplementary Fig. 36b−d). Meanwhile, the cycling stability of Zn||Zn cell using com ZnBr2 as a solid-state electrolyte was also tested at 0.1 mA cm−2. As shown in Supplementary Fig. 37, the short cycle life and large voltage polarization were delivered, suggesting poor ionic conduction. To understand the reversibility of the Zn2+ plating/stripping process, Zn|PipZnBr4|Ti asymmetric cells were tested. Supplementary Figs. 38 and 39 show the rate cyclic voltammetry (CV) curves of the Zn||Ti cell at 25 °C. The extracted curves suggest that reduction potential decreases (ΔV=70 mV) as the scanning rate increases (Supplementary Fig. 39), implying the weak parasitic reaction at the Zn/PipZnBr4 interface, which is different from the continuous reduction potential increase of 2 M ZnSO4 aqueous electrolyte43–45. Figure 6e exhibits the long-term cycle performance of a Zn||Ti cell with PipZnBr4 solid-state electrolyte. It is found that the cell yields a stable and high average CE of 98.9% at 2 mA cm−2 and 0.5 mA cm−2 after 350 cycles associated with the close contact at the Zn/PipZnBr4 interface. The corresponding voltage profiles at different cycles are presented in Supplementary Fig. 40, and deliver a lower voltage hysteresis (~170 mV) than that of the Zn||Ti cell (~184 mV) with 1 M ZnSO4 electrolyte due to inhibition of side reaction33. As the cutoff capacity is further increased to 1.0 mAh cm−2, the Zn||Ti cell still realizes a long lifespan over 115 cycles with an average CE of ~99.6% (Supplementary Fig. 41a), and the corresponding voltage profiles are shown in Supplementary Fig. 41b. These prospective results provide support for using PipZnBr4 as a solid-state electrolyte in highly reversible ZMBs.

To evaluate the feasibility of PipZnBr4 electrolyte in practical applications, a Zn|PipZnBr4|iodine (I2) full all-solid-state Zn-metal battery was assembled, in which I2 is adopted as a positive electrode and the mass of I2 loading is ~3 mg cm−2. Figure 6f shows the long-term cycle performance of the Zn||I2 cell. A residual capacity of 234.5 mAh g−1 and a low capacity degradation rate of 0.056% per cycle is achieved after 200 cycles at 0.2 mA cm−2. The corresponding galvanostatic charge/discharge (GCD) curves were observed and found that these curves exhibit good overlap, except for the second cycle (Fig. 6g), implying the high reversibility of Zn2+ in PipZnBr4 solid-state electrolyte. Supplementary Fig. 42 illustrates the CV curves at 0.1 and 0.2 mV s−1 of full cells with distinct iodine-related redox peaks, which exhibit a diminished voltage gap between the cathodic and anodic peaks, indicating an improvement in reaction kinetics. This phenomenon is also corroborated by EIS analysis, which demonstrates reduced charge transfer resistance (Supplementary Fig. 43). When the current density is increased to 0.5 mA cm−2, the Zn|PipZnBr4|I2 cell can still operate 100 cycles after the pre-activation process at the current density of 0.5 mA cm−2 for 3 cycles (Supplementary Fig. 44a) and corresponding GCD curves are exhibited in Supplementary Fig. 44b. These good performances are attributed to the uniform Zn plating/stripping at the Zn/PipZnBr4 interface due to the close contact of the interface, inducing homogeneous Zn2+ flux. Additionally, the FTIR and XPS spectra of the I2 positive electrode were also tested to exclude the possibility of halogen ion conductivity. Compared with the pristine I2 electrode, the cycled I2 positive electrode exhibits no additional peaks (Supplementary Fig. 45), particularly in the wavenumber ranges of 520~600 cm−1 and 1300~1500 cm−1, where peaks associated with the Br element would typically appear. XPS spectra test with Ar+ sputtering was also used to assess the chemical component on the I2 positive electrode surface, when the Zn|PipZnBr4|I2 full cell has undergone 10 cycles at 0.2 mA cm−2 and then is discharged to 1.0 V (Supplementary Fig. 46). More significantly, as shown in Supplementary Fig. 47, PipZnCl4 as electrolyte was also evaluated in a Zn symmetric cell and a Zn||I2 full cell. Finally, to exclude the possibility of halide ion conduction, CsSnBr3 and CsSnCl3 samples as the blocking layer were applied in the solid-state PipZnBr4 and PipZnCl4 electrolytes (Supplementary Fig. 48). It is found that the CV curces of Zn|PipZnCl4|CsSnCl3|I2 and Zn|PipZnCl4|CsSnCl3|I2 full cells at 0.1 mV s−1 have no oxidation-reduction peak and Zn||Zn symmetrical cells at 0.1 mA cm−2 and 0.1 mAh cm−2 are almost unable to operate (Supplementary Fig. 49). These results demonstrate that although halide ions could migrate in PipZnBr4 and PipZnCl4 electrolytes, Zn-ion conduction in the halide-based ZMBs play a dominant contribution.

Methods
Orbital analysis and interaction diagram
The orbital analysis of tetrahedral and octahedral structures was performed with the Gaussian16 package46. The Becke 3-paeameter-Lee–Yang–Parr (B3LYP)47 correlation functional at def2-TZVP48 level was used for the optimization and self-consistent field calculations, with density functional theory-D3 (Becke-Johnson)49,50 method for the van der Waals interactions. The visualization of the orbital results was conducted using Multiwfn51 and Visual Molecular Dynamics52 software. The irreducible representations, such as a1, a1g, t1g, eg, were also obtained from the Gaussian16 calculation results. The bonding interactions illustrated in the interaction diagram were deduced based on the orbital symmetry information and the spatial distribution/overlap characteristics of the visualized orbitals.

COBI analysis
The structural models of 4-coordinated/6-coordinated LiCl sourced from the Materials Project database53. The optimization calculations of the structures were performed using the projector-augmented wave54 method, as incorporated in the Vienna Abinitio Simulation Package (VASP)55 code. The sutoff energy was set to 500 eV. Γ-centered k-meshes, with a k-spacing of 0.2 Å−1, were employed for the sampling of the Brillouin zones. The generalized gradient approximation (GGA) Perdew-Burke-Ernzerhof (PBE)56 was used as the exchange-correlation functional. The energy convergence criterion for COBI analysis is 10−7 eV. The COBI values were calculated by using the Lobster57 code.

BVSE calculations
BVSE calculations were conducted utilizing the softBV27,28 software, employing structural models sourced from the Inorganic Crystal Structure Database. The energies of Li or Zn sites within structures were calculated on a three-dimensional grid of points, each with a resolution of 0.1 Å. The pathways of Li+ or Zn2+ migration were automatically obtained by the software.

Density functional theory (DFT) calculations for electronic structures
DFT calculations were performed using the VASP55 code. The cut-off energy, k-points, and exchange-correlation functional settings for all the structures (PipZnBr4, PipZnCl4, and their possible secondary phases) optimizations were the same as those used in the structure optimization for COBI analysis. The crystal structures of PipZnBr4 and PipZnCl4 were fully relaxed until the total force on each atom was <0.03 eV Å−1. To rectify the underestimation of the band gaps in GGA-PBE methods, the band structure, DOS, and charge density (CBM and VBM) were calculated by using the HSE0658 hybrid functional. The visualization of the charge density is carried out using the VESTA59 software. The chemical phase diagrams for PipZnBr4 and PipZnCl4 were drawn by the Chesta60 code.

Defect calculations
For modeling the intrinsic defects in PipZnBr4 and PipZnCl4, a 2×2×2 supercell containing 736 atoms and Γ-only k-mesh was used. The perfect and defect structures were fully relaxed by the GGA-PBE method until the total force on each atom was <0.03 eV Å−1. The formation enthalpy (ΔHD,q) for a defect (D) in a charge state q was calculated by refs. 61,62
ΔHD,q = ED,q − Eh + q(EV + EF) + n iμi + Ecorr (1)
where ED,q is the total energy of the supercell containing the defect (D) in the charge q. Eh is the total energy of the perfect host supercell. EV is the VBM level, while EF is the Fermi level referred to EV. The parameter ni indicates the number of i atoms added (ni < 0) or removed (ni > 0) to form a defect (D). μi is the chemical potential of the atom i. Ecorr represents the total correction for ΔHD,q, including the band-filling correction, the potential alignment correction, and the image charge correction63,64. To address the band edge problem using the GGA-PBE method61,65, corrections were applied for both the valence band maximum (VBM) overestimation and the CBM underestimation. For PipZnBr4, the corrections are ΔVBM=0.89 eV and ΔCBM=0.41 eV, while for PipZnCl4, they are ΔVBM=1.02 eV and ΔCBM=0.44 eV. These corrections are incorporated by adding ΔVBM (or ΔCBM) to the formation energy of deep acceptor (or donor) defects calculated using the HSE hybrid functional method. For shallow defects, the formation energy is adjusted to maintain its relative position to the edge. The charge transition level and the equilibrium Fermi level (EF,e) were determined using methods identical to those employed in our previous research66.

Zn-ion migration energy barrier calculations
The initial sites of the Zn-ion migration calculations were determined through numerous DFT energy calculations. Subsequently, the final position was identified as the symmetrical position in the corresponding direction of the supercell. For the migration energy barrier calculations along the a, b, and c directions, supercells of dimensions 2×1×1, 1×2×1, and 1×1×2 were employed, respectively. The diffusion pathways and barrier energies of Zn-ion diffusion throughout the PipZnBr4 and PipZnCl4 were calculated by using the Cl-NEB30 method implemented in VASP. The Cl-NEB calculations, with a k-spacing of 0.2 Å−1, were determined to converge when the maximum force of each image was <0.03 eV Å−1.

Preparation of PipZnBr4 solid electrolyte and electrode
Preparation of PipZnBr4-based solid electrolyte. The solid electrolyte was synthesized by hydrothermal method. Stoichiometric amounts of ZnBr2 (over 98% purity, Shanghai Adamasi Reagents Co., Ltd.) and piperazine (over 99% purity, Shanghai Adamasi Reagents Co., Ltd.) were dissolved in hydrobromic acid (HBr, 48 wt%, Shanghai Aladdin Biochemical Technology Co., Ltd.). The mixture solution was added into the Teflon autoclave and was heated to 90 °C within 120 min, and kept for 2440 min. Next, the reactor was cooled to 25 ± 0.5 °C within 1200 min. Finally, the solution was filtered and collected to yield the soft co-crystalline halides. Specifically, to synthesize PipZnBr4 (1:1) and PipZnBr4 electrolyte, including ZnBr2 (0.6 mmol) and piperazine (0.6 mmol) or ZnBr2 (3.0 mmol) and piperazine (0.6 mmol) were added in HBr (3 ml), respectively. Then the solution was transferred to a 20 ml Teflon autoclave. The as-synthesized PipZnBr4 (1:1) and PipZnBr4 powders were collected and stored in a 2 ml centrifuge tube for further characterization and electrochemical testing. To assemble the cell, the PipZnBr4 powders were pressed at around 20 MPa for 2 min in a stainless steel mold (ϕ=16 mm) by Infrared Pellet Press (PC-30, Tianjin Jingtuo Instrument Technology Co., Ltd) under an air-conditioned room at 25 ± 0.5 °C. To evaluate the performance of the PipZnBr4-based solid electrolyte, the thickness of the electrolyte was controlled to approximately 430 μm. The reference 1.2 M Zn(OTf)2 electrolyte was prepared by dissolving Zn(OTf)2 (98%, Shanghai Aladdin Biochemical Technology Co., Ltd.) in deionized water33. The hydrogel electrolyte was prepared according to reference 34. First, gelatin (2 g, 98%, Shanghai Aladdin Biochemical Technology Co., Ltd.) was dissolved in 40 ml deionized water under stirring at 60 °C for 30 min. Subsequently, a specified quantity of oxidized dextran (OD) solution was introduced into the gelatin solution. The reaction system was then cooled below 25 ± 0.5 °C and held for 4 h to promote the coordination-driven modification of gelatin chain by OD. Following this, phosphate-buffered saline (20×, Shanghai Aladdin Biochemical Technology Co., Ltd.) was incorporated, and the mixture was stirred at 60 °C for 1.5 h. Methacrylic anhydride (15%, w/v, Shanghai Aladdin Biochemical Technology Co., Ltd.) and the photo-initiator lithium phenyl-2, 4, 6-trimethylbenzoylphosphinate (LAP, 0.05% (w/v), 98%, Macklin, Shanghai Biochemical Technology Company) were added dropwise in sequence. The resulting prepolymer solution was dialyzed in water at 40 °C, transferred into a polytetrafluoroethylene (PTFE) mold, and frozen at 4 °C. Finally, the ODGelMA hydrogel film was obtained after drying and photocrosslinking under 365 nm UV light.

Positive electrode preparation
The positive electrode was fabricated according to the method we reported earlier29,34. Carbon black (YP-50, Kuraray Co., Ltd), acetylene (Suzhou Sinero Technology Co., Ltd), and PTFE powder (Guangdong Canrd New Energy Technology Co., Ltd.) with a mass ratio of 8:1:1 were mixed and hot-pressed onto a thin sheet (denoted as C). The as-synthesized C and Ti mesh (100 mesh, Hebei Mengshang Metal Mesh Industry Co., Ltd) were pressed to build a positive electrode matrix (Ti@C). A mixed solution with iodine (I2) as active material was prepared by dissolving 1 M potassium iodide (KI, 99%, Shanghai Aladdin Biochemical Technology Co., Ltd.) and 0.5 M zinc sulfate (ZnSO4, 99%, Shanghai Aladdin Biochemical Technology Co., Ltd.). Finally, the I2 electrode was obtained by depositing I2 onto the Ti@C at 5 mA for 6 min and dried at 40 ± 0.5 °C for 24 h. The I2 loading mass is about 3 mg cm−2.

Characterizations
The structures of the PipZnBr4 (1:1) and PipZnBr4 solid electrolyte, as well as the electrode after cycling Zn plating/stripping, were characterized by FTIR spectra (AVATAR 370 mode) and Raman spectroscopy (523 nm, Horiba LabRAM HR800). The XRD (Rigaku, D/Max-2500) pattern was scanned from 5 to 80° with a Cu Kα X-rays (λ=1.5406 Å) radiation at a scan rate of 10°/min and a step of 0.02°. The PipZnBr4 solid-state electrolyte and the surface morphology of the Zn electrode were observed by SEM (Hitachi SU-4800 FE-SEM) with an EDS detector (UItim Max 20) and EBSD (Nordlys Max 3 detection system) under the electron current of 13 pA and accelerating voltage of 2 kV. The surface feature of the cycled Zn electrode was observed by CLSM (Olympus LEXT OLS5000-SAF) and AFM (Bruker, Dimension icon). Chemical bonding information was obtained via XPS measurements performed on a PHI 5000 VersaProbe II spectrometer, utilizing monochromatized Al Kα radiation for excitation. An integrated argon ion gun was utilized for sample sputtering, with a sputter rate calibrated against SiO2 (25 nm min‒1) over a 1×1 mm region. The C–C peak at 284.8 eV was used as a reference to calibrate the binding energies presented in this work. Non-destructive visualization of the solid-state battery’s interior and the Zn electrode’s deposition morphology was characterized using CT (Phoenix V MDTC-EQ-M60-01). The system operated at 140 kV and 180 μA, capturing data at 3 μm spatial resolution with 334 ms exposures. A 3D reconstruction model and corresponding cross-sectional views were obtained by integrating the multi-angle scan datasets. The Zn element content in PipZnBr4 (1:1) and PipZnBr4 solid electrolyte was assessed in parallel three times by ICP (Thermo Fisher iCAP Q). ATHENA software was used to process and normalize the XAFS spectra67. Wavelet transform analysis was then carried out on the EXAFS data employing the HAMA program.

Electrochemical measurements
Zn foil (100 μm, 99%, Wuhan Luojia Yanxin Technology Co., Ltd.) was polished with abrasive paper (20000 mesh) and cleaned with ethanol (AR, Shanghai Lingfeng Chemical Reagent Co., Ltd.) to remove the inactive layer. Zn foil was cut into a disc (ϕ=12 mm) as an electrode. Ti foil (10 μm, 99.99%, Wuhan Luojia Yanxin Technology Co., Ltd.) was cut into a disc (ϕ=16 mm) to be applied as a positive electrode for Zn||Ti cells. All the testing CR2032 coin cells (stainless steel, Guangdong Canrd New Energy Technology Co., Ltd.) consisted of a stainless steel gasket (0.05 mm) and a shrapnel (0.1 mm), and they were assembled under a pressure of 50 ± 1 kgf cm−2 in an open-air environment by using PipZnBr4 solid electrolyte and tested on a battery test system (Neware) at 40 ± 1 °C. All the batteries were subjected to electrochemical characterization without any additional pressure being applied. The GCD cycles of Zn||I2 using PipZnBr4 solid electrolyte were conducted in a voltage range of 0.3~1.3 V (Zn2+/Zn). EIS tests were performed on a VSP-3e Potentiostat electrochemical workstation (BioLogic) operating in a potentiostatic mode. A frequency range of 100 kHz–50 MHz was explored, collecting 10 points per decade. To measure the EIS of the solid-state electrolyte at different temperatures, the electrolyte was measured in parallel three times. CV measurements were performed on Zn||Ti coin cells, employing 12 mm Zn foil and 16 mm Ti foil. The testes covered a potential range of −0.2 to 1.0 V at different scan rates (0.1, 0.2, 0.5, and 1.0 mV s−1) on an electrochemical workstation. The nucleation overpotential of the Zn electrode was studied on the electrochemical workstation based on the Zn|PipZnBr4|Zn cell. LSV test of Zn|PipZnBr4|Ti cell was conducted on workstations at 1 mV s−1 at −1.8~2.8 V. The Zn transference number is calculated by a current-time curve under constant polarization at −50 mV for 700 s. The Tafel profiles were operated on Zn|PipZnBr4|Ti cell. The ionic conductivity (σ) of the as-synthesized PipZnBr4 solid electrolyte was measured by EIS measurements. The σ was calculated based on the equation of σ=l/RS, where l represents the electrolyte thickness (430 μm), R is the total impedance according to EIS measurement, and S is the area (2.0 cm2). To gain the Arrhenius plots, carriable-temperature EIS was conducted from 0 to 60 °C. The activation energy (Ea) was calculated on the Arrhenius equation. The cells were characterized in terms of their galvanostatic charge/discharge tests, CV curves, and rate performance using a 12-mm diameter Zn foil negative electrode and a 1×1 cm I2 positive electrode with a controlled ~3 mg cm−2. The titanium mesh served as the substrate, with the I2 positive electrode material loaded on one side facing the solid-state electrolyte. Area current densities were calculated by the area of the negative electrode, while specific currents were determined based on the mass of iodine present in the positive electrode. All the assembled solid-state Zn cells with the sealing pressure of 50 ± 1 kgf cm−2 were left for 6 h before the electrochemical test. Then, the cells underwent charging and discharging tests without any additional pressure being applied. At least ten cells were tested for a single electrochemical experiment to guarantee the reproducibility. The CE of a cell is the ratio of the discharge capacity to the charging capacity of the preceding charge cycle.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This manuscript presents an interesting story that the zinc-based hybrid metal halides (ZHMHs), C4N2H12ZnBr4 and C4N2H12ZnCl4, could be used as zinc ion solid electrolytes, because of their analogies to the halide lithium ion electrolytes. The electronic structure analysis of these hybrids is ok to confirm the electronic insulation for the solid electrolyte. However, I have some doubts about the migration mechanism for the ionic conduction and the electrochemical process of the fabricated cells. To allow the publication of this manuscript without severe problems, I think the authors should clarify the zinc ion conduction mechanism and real electrochemical reactions by providing sufficient evidence.\nHere are my specific comments:\n1. For the proposed concept of an atomic substitution strategy, I do not think it accurately reflects the central idea that the authors intend to present. The modification of the halide solid electrolytes involves using the organic cations to replace some inorganic species, which can not be conceptualized as “atomic substitution”. This inconsistency would make the readers very confused about how advanced the proposed atomic substitution strategy.\n2. A heavy, uneasy concern comes from the ambiguous ionic migration analysis presented in this manuscript. Firstly, the proposed Zn²⁺ migration occurs primarily through interstitial sites rather than the experimentally identified stable Zn-ion site (4c site), which raises doubts about the validity and connectivity of these migration paths. Furthermore, the authors' proposed mechanism of \"defect-facilitated ion conduction\" should be supported by more comprehensive data, either computationally or experimentally. Lastly, given the presence of abundant halogen defects, the authors should consider/examine the possibility of halogen ion conduction.\n3. The Arrhenius plot of PipZnBr4 in Fig.4e exhibits significant deviations from linearity, indicating a poor correlation between the conductivity and temperature. Consequently, the derived activation energy of 0.11 eV (an unusually low value for solid-state electrolytes) appears questionable. This anomalous result may suggest poor thermal stability of the material and therefore additional temperature-dependent characterizations (e.g., thermal gravimetric analysis) are recommended. I also suggest that the authors repeat this experiment with tighter temperature control (e.g., 5℃ intervals). If the nonlinearity persists, a reasonable explanation for both the low activation energy and non-Arrhenius behavior should be provided.\n4. Besides the concern about the conduction of ionic species, the electrochemical process in the fabricated cells needs careful analysis and more convincing evidence. It is necessary to identify the real conducting ions in the reported electrolytes. A Zn2+ block layer for the impedance test should be employed. In addition, chloride based electrolytes should also be tested in the symmetric cells and full cells. In particular, the charge/discharge curves of chloride-type electrolyte based full cell can reveal the details of electrochemical reactions in the cell.\n5. There are many minor errors in the manuscript. For example, Fig.2j lacks the DOS data for Cl 3p; Line 231: an incorrect full name of CI-NEB; Line 453: duplicate \"PipZnBr4\"; Line 463: incorrect activation energy value. These errors could be avoided if the authors prepare the manuscript in a well-trained and multiple-checked way.", "reviewer_2": "Hu et al. proposed an interesting design rule for Zn-based metal halide electrolytes by replacing inorganic atoms with organic clusters. The hypothesis is that organic clusters may provide more abundant defect properties and thereby enhance the kinetics of zinc-ion transport. They investigated the hypothesis both computationally and experimentally. Guided by computational results, they synthesized PipZnBr4, which shows high conductivity. Overall, I found the work interesting and relevant for the battery community but there are a few technical concerns that need to be addressed before being considered for publication in Nature Communication.\nTechnical concerns:\n1. Applicability of BVSE can be limited and material-dependent, the authors need to provide a discussion on the limitation of the approach and how it can still be used.\n2. The author assumed that Pip is one “component” like Zn and Br/Cl when predicting chemical phase diagrams and defect formation energy. But this assumption becomes questionable when the results of defect formation enthalpy for point D1/D2 were shown, especially the formation enthalpy is very negative (at the order of 1 eV)\n3. Related to point 2 above, the formation energy of C,H,N vacancies can be used to test if the assumption is valid. In addition, what about C/H/N interstitials? Those can be charged defects and move the equilibrium Fermi energy.\n4. Since the band edge/gap correction is not included in the defect formation enthalpy, the uncertainty of the predicted values is unknown and thus not strong enough to fully support the discussion. The authors need to either add the correction to the results or discuss the limitation.\n5. One of the main conclusions from computational results is that PipZnBr4/PipZnCl4 can be good insulator based on the large band gap and defect calculations. While the results of equilibrium Fermi energy being at least about 1 eV away from band edges can support the argument, a more quantitative analysis on the number of electron/hole carriers as a function of synthesis temperature can be a stronger support.\n6. For NEB calculations, how are the initial path determined? In other words, how can we know this is the minimum energy path among all the possible pathways.\n7. “The broken tetrahedra leaves a low migration energy barrier for Zn ions in the material,…” It is unclear whether broken tetrahedra is included in the NEB calculations. If not, the statement is not fully supported\n8. For XRD results, do the authors identify competing phases? This can provide a hint whether the material is near the phase boundary, i.e. close to energy of the A,B,C points shown in Fig 2b.\n9. For the measured stability window, how does this compare to prediction (Fig 2b)? Related to point 2 above, this seems the indicate the breakdown of the assumption that treats pip as one component.\n10. Why is the measured activation barrier (~0.11 eV) much lower than the calculated migration barrier?\n11. Here is a general question since I see some connection between this work and hybrid perovskite. Can the authors comment on the stability of PipZnBr4 against water/moisture?\nReadability/Presentation:\n1. In abstract, the importance of Zn-based battery is unclear\n2. In abstract, the statement, “low energy barriers between organic groups and halogen tetrahedra”, should be more precise. I understand this statement only after seeing Fig 3 i/k. I originally thought this means Zn diffuse between organic groups and halogen tetrahedra.\n3. In abstract, minor grammar issue for the last sentence “… a new approach to designing ….”\n4. One general comment, the readership of nature communication is not limited to chemist. Some of the statements might be straightforward to chemist, but further elaboration can be helpful for a broader audience.\n5. In introduction, “Zinc ion possesses a high valence charge and mediocre electronegativity, making the study of solid electrolytes that utilize zinc ion transport as the primary conduction mechanism particularly challenging” the connection is not fully clear.\n6. In introduction, “These two materials show good electronic properties…” The statement is vague\n7. For FigS2 and S3, the meaning of red and blue dots is not defined.\n8. “The localized charge of ionic bonds suggests that Li-based halides can conduct lithium ions” elaboration is needed.\n9. On the related topic to the point 8 above, Crystal Orbital Bond Index (COBI) analysis could be a good way to quantify the covalency. (optional)\n10. How the authors created the interaction diagram is not fully clear in current format\n11. For many figures, color scheme for each atom is currently missing.\n12. On line 171, “The valence band maximum (VBM) mainly consists of the nonbonding states of Br 4p orbitals” This statement can only be drawn with further details discussed later in the paragraph.\n13. “For PipZnCl4, due to the electronegativity of the Cl atom is more negative than that of the Br atom, leading to a significantly higher anti-bonding orbital between Zn 4s and Cl 3p. The case results in a larger band gap (6.31 eV) for PipZnCl4 than that of PipZnBr4 (5.37 eV)” This is a correct statement but might not be intuitive.\n14. Likely a typo in figure 2j, it should be Cl instead B\n15. “The Fourier transform spectra of the extended XAFS (FT-EXAFS) curves in R-space show that the bond length of Zn−Br in PipZnBr4 is shorter than that in pure ZnBr2 (Fig. 6b), which suggests a relatively weaker Zn−Br bond in PipZnBr4.” Like a typo since Zn-Br is longer in Fig 6b.\n16. In the Discussion, there is a typo “low activation energy (0.01 eV), ...” since the value is different from ones in the Results section."}, "author_response": "Response to reviewers’ comments\nReviewer #1\nThis manuscript presents an interesting story that the zinc-based hybrid metal halides (ZHMHs),\nC N H ZnBr and C N H ZnCl , could be used as zinc ion solid electrolytes, because of their\n4 2 12 4 4 2 12 4\nanalogies to the halide lithium ion electrolytes. The electronic structure analysis of these hybrids\nis ok to confirm the electronic insulation for the solid electrolyte. However, I have some doubts\nabout the migration mechanism for the ionic conduction and the electrochemical process of the\nfabricated cells. To allow the publication of this manuscript without severe problems, I think\nthe authors should clarify the zinc ion conduction mechanism and real electrochemical reactions\nby providing sufficient evidence.\nHere are my specific comments:\nResponse: We sincerely thank Reviewer #1 for the comprehensive evaluation and constructive\nfeedback. We are appreciative of their acknowledgment of the electrical insulating property and\ninteresting potential of our ZHMHs as solid-state electrolytes. We also appreciate the concerns\nraised by Reviewer #1 on the migration mechanism of ion conduction and the electrochemical\nprocess of the fabricated cells. We are trying to solve these concerns point-by-point as follows.\n1. For the proposed concept of an atomic substitution strategy, I do not think it accurately\nreflects the central idea that the authors intend to present. The modification of the halide solid\nelectrolytes involves using the organic cations to replace some inorganic species, which can not\nbe conceptualized as “atomic substitution”. This inconsistency would make the readers very\nconfused about how advanced the proposed atomic substitution strategy.\nResponse: We sincerely appreciate the reviewer's insightful comment regarding the\nterminology “atomic substitution” strategy. We agree that the “atomic substitution” strategy did\nnot accurately reflect the central idea of our modification strategy, where organic cations replace\ninorganic species within the halide solid electrolytes. In our original manuscript, the proposed\n“atomic substitution” is a broad generalization originating from research on metal halide\nmaterials, which may create confusion for readers in other fields. As the reviewer rightly points\nout, our modification of the halide solid electrolytes indeed involves molecular/ionic groups\nrather than individual atoms, and using “atomic substitution” may confuse the strategy's essence.\nConsequently, following the reviewer’s valuable suggestion, we have replaced “atomic\nsubstitution” with the more precise term “composition engineering” throughout the revised\nmanuscript. This change better conveys our core concept: the design concept of the material's\ncomposition to transform inorganic zinc-based metal halides (originally lacking Zn²⁺ migration\nchannels) into organic-inorganic hybrid zinc-based metal halides. After the “composition\nengineering”, the organic-inorganic hybrid zinc-based metal halides create the necessary\ninternal pathways for Zn²⁺ migration. We hope this revised terminology significantly enhances\nthe clarity, accuracy, and rigor of our presentation, preventing potential reader confusion and\nmore accurately representing the advancement of the proposed strategy.\nWe thank the reviewer again for this valuable suggestion that has improved the readability\nand scientific precision of our work. We have modified the relevant statements in the abstract,\nintroduction, and discussion sections of the revised manuscript.\nRevisions in the manuscript:\n“Herein, we propose a universal composition engineering strategy to design …”\n“In this work, we proposed a composition engineering strategy for designing …”\n“… to form zinc ion channels. After component engineering design, the target material contains\norganic cationic groups, which exhibit a comparatively loose spatial configuration. …”\n“Fig. 1| Comparison of properties of Li/Zn-based halide materials and composition\nengineering design of zinc-ion solid electrolytes. … g Schematic of composition engineering\ndesign of zinc-ion solid electrolytes.”\n“In summary, a simple and effective composition engineering strategy for constructing halide\nzinc-ion electrolytes was proposed based on the theoretical prediction, …”\n2. A heavy, uneasy concern comes from the ambiguous ionic migration analysis presented in\nthis manuscript. Firstly, the proposed Zn²⁺ migration occurs primarily through interstitial sites\nrather than the experimentally identified stable Zn-ion site (4c site), which raises doubts about\nthe validity and connectivity of these migration paths. Furthermore, the authors' proposed\nmechanism of \"defect-facilitated ion conduction\" should be supported by more comprehensive\ndata, either computationally or experimentally. Lastly, given the presence of abundant halogen\ndefects, the authors should consider/examine the possibility of halogen ion conduction.\nResponse: We thank the reviewer for his/her valuable suggestion. We agree with the reviewer\nthat it is important to further prove the ionic migration in this manuscript. Following the\nsuggestion, we have done a detailed analysis through both theoretical calculations and\nexperiments, as detailed below.\n(1) Firstly, we apologize for the confusion caused by our previous description of the Zn-\nion migration sites. In our design, our original intention was to use the spatial structure of\norganic clusters to provide Zn-ion migration channels, while the zinc halide tetrahedral\nstructure was retained. Consequently, the feature that zinc ions are difficult to conduct in the\ntetrahedron is unchanged. As shown in Fig. R1, we have calculated the Zn²⁺ migration at the\n4c site of PipZnBr and PipZnCl . The results show that the migration energy barriers of\n4 4\nPipZnBr and PipZnCl are 1.40 eV and 1.26 eV, respectively. This result is in line with\n4 4\nexpectations. We would like to emphasize that there are no channels for zinc ions inside or\noutside of the tetrahedra in the inorganic crystal. To avoid confusion, we modified the relevant\nsentences in the revised version.\nFig. R1 | Zn2+-migrations schematic diagrams at 4c site of PipZnBr and PipZnCl and their\n4 4\ncalculated Zn-ion energy migration pathways. a PipZnBr . b PipZnCl .\n4 4\n(2) Secondly, we sincerely appreciate the reviewer pointing out the comment on “defect-\nfacilitated ion conduction”. In the defect analysis, we have considered all of the possible\nintrinsic point defects of PipZnBr₄ and PipZnCl₄. The calculation results reveal that both\nmaterials exhibit a complex defect chemistry, with halogen vacancies (V /V ) consistently\nBr Cl\nacting as the dominant donor defects. The equilibrium Fermi level under all chemical conditions\nremains far from both band edges, resulting in extremely low charge carrier concentrations and\nconfirming the intrinsic insulating nature of the compounds. The dominant halogen vacancy\ndefects suggest that both materials are susceptible to breaking the original zinc-halogen\ntetrahedral structure due to these defects, which could provide more space for zinc ion\nconduction. Here, the space formed by such vacancies is also adjacent to the organic groups, so\nthat the material may have broader ion migration channels. We have also conducted migration\ncalculation studies on materials with structural defects and found that the results were difficult\nto converge. Consequently, to ensure full rigor, we refine our initial affirmative assertion to\nmerely suggest that vacancies might offer increased space for the migration of Zn ions. In our\nwork, the migration channel for zinc ions is primarily provided by the composition engineering\nstrategy, thus it does not affect the conclusion of the entire paper.\n(3) Thirdly, we have tested the ionic conductivity of ZnBr and the synthesized PipZnBr\n2 4\nsamples at 25 ℃. As shown in Supplementary Figs. 22 and 23, the ionic conductivities of\nZnBr and PipZnBr are 0.09×10−3 S cm−1 and 0.29 ×10−3 S cm−1, respectively. Obviously, the\n2 4\nhigher ionic conductivity of the synthesized PipZnBr sample than that ZnBr verifies that the\n4 2\nsynthesized material contributes to facilitating ionic conduction. Meanwhile, the\nelectrochemical performance of ZnBr as a solid-state electrolyte has also been assessed by the\nZnBr sample can only work 400 h at current density of 0.1 mA cm−2 and plating capacity of\n2\n0.1 mAh cm−2 at 40 ℃, which is far lower than the cycle life of the synthesized PipZnBr\n4\nsamples (over 1100 h in Fig. 6d). Furthermore, as the cycle life increases, the voltage\npolarization of the Zn cell with ZnBr becomes increasingly larger, indicating decresed ionic\n2\nconductivity during cycling. Based on the above theoretical calculations and experimental\nresults, it is believed that the PipZnBr sample possesses a high ionic conduction.\n4\n0.1 mAh cm−2.\n(4) Lastly, to exclude the possibility of halogen ion conduction, FTIR and XPS spectra of\nthe I cathode were tested. As shown in Supplementary Fig. 45, the I cathode after cycling\n2 2\nexhibits no additional peaks compared to the pristine electrode, particularly in the wavenumber\nranges of 520 ~ 600 cm−1 and 1300 ~ 1500 cm−1, where peaks associated with the Br element\nwould typically appear. Additionally, XPS spectra test with Ar+ sputtering was used to assess\nthe chemical component on the I cathode surface, when the Zn||PipZnBr ||I cell has undergone\n2 4 2\n10 cycles and then is discharged to 1.0 V (Supplementary Fig. 46). On the top surface (before\nsputtering), Zn and Br signals involving Zn-Br and Br-Br bonds are detected, which is most\npossibly caused by the PipZnBr solid electrolyte adhering to the surface of the I cathode when\n4 2\nassembling the Zn||PipZnBr ||I cell. With the Ar+ sputtering (690 s), the signals of Br-related\n4 2\npeaks become very weak, whereas the Zn signals still remain very strong, indicating that the\nZn2+ ion conduction plays the dominant role in the halide-based zinc-ion batteries. In addition,\nwe have supplemented the ionic conductivity of ZnBr and the cycling performance of\n2 2\nand figures have been revised manuscript and supplementary information.\nRevisions in the manuscript:\n“The dominant halogen vacancy defects suggest that both materials are susceptible to breaking\nthe original zinc-halogen tetrahedral structure due to these defects, which may provide more\nspace for zinc ion conduction.”\n“… indicating its promising application in ASSZBs. To verify the excellent ionic conductivity\nof PipZnBr , the ionic conductivity of ZnBr has been assessed at 25 ℃ (Supplementary Fig.\n4 2\n23). …”\n“… the PipZnBr -based Zn symmetric cell still works over 400 h (Supplementary Fig. 36a),\n4\nimplying the stable ionic migration. At an increased current density of 1.0 mA cm−2, long\ncycling with voltage hysteresis (~860 mV) can still be attained in the symmetric Zn cell\n(Supplementary Figs. 36b−d). Meanwhile, the cycling stability of Zn||Zn cell using ZnBr as\n2\na solid-state electrolyte was also tested at 0.1 mA cm−2. As shown in Supplementary Fig. 37,\nthe short cycle life and large voltage polarization were delivered, suggesting poor ionic\nconduction. …”\n“… These excellent performances are attributed to the uniform Zn plating/stripping at the\nZn/PipZnBr interface due to the close contact of the interface, inducing homogeneous Zn2+\n4\nflux. Additionally, the FTIR and XPS spectra of the I cathode were also tested to exclude the\n2\npossibility of halogen ion conduction. Compared with the pristine I cathode, the cycled I\n2 2\ncathode exhibits no additional peaks (Supplementary Fig. 45), particularly in the wavenumber\nranges of 520 ~ 600 cm−1 and 1300 ~ 1500 cm−1, where peaks associated with the Br element\nwould typically appear. XPS spectra test with Ar+ sputtering was also used to assess the\nchemical component on the I cathode surface, when the Zn||PipZnBr ||I full cell has undergone\n2 4 2\n10 cycles and then is discharged to 1.0 V (Supplementary Fig. 46)…”\nRevisions in the supplementary information:\nSupplementary Fig. 23 | Nyquist plot of ZnBr at 25 ℃.\n2\nelectrolyte at current density of 0.1 mA cm−2 and plating capacity of 0.1 mAh cm−2 at 40 ℃.\nSupplementary Fig. 45 | Characterization of I cathode from Zn||PipZnBr ||I full cell. a\n2 4 2\nFTIR spectra of the I cathode before and after 10 cycles at 0.2 mA cm−2. b, c FTIR spectra of\n2\nthe I cathode where peaks associated with the Br element would typically appear.\n2 4 2\na-c Full spectra, Br 3d, and Zn 2p high-resolution XPS spectra of I cathode after 10 cycles. d-\n2\ne Full spectra, Br 3d, and Zn 2p high-resolution XPS spectra of I cathode after 10 cycles and\n2\nthen discharging to 1.0 V with Ar+ sputtering for 690 s.\n3. The Arrhenius plot of PipZnBr in Fig. 4e exhibits significant deviations from linearity,\n4\nindicating a poor correlation between the conductivity and temperature. Consequently, the\nderived activation energy of 0.11 eV (an unusually low value for solid-state electrolytes) appears\nquestionable. This anomalous result may suggest poor thermal stability of the material and\ntherefore additional temperature-dependent characterizations (e.g., thermal gravimetric analysis)\nare recommended. I also suggest that the authors repeat this experiment with tighter temperature\ncontrol (e.g., 5 ℃ intervals). If the nonlinearity persists, a reasonable explanation for both the\nlow activation energy and non-Arrhenius behavior should be provided.\nResponse: We sincerely appreciate your insightful suggestion to add temperature-dependent\ncharacterizations and repeat the experiment on ion conductivity.\nAccording to the reviewer’s suggestion, we have provided the thermal gravimetric analysis\nof PipZnBr . As shown in Supplementary Fig. 21a, the weight of PipZnBr remains almost\n4 4\nunchanged when the temperature rises to 150 ℃. After heating at 150 ℃ for 30 min, the XRD\npatterns of PipZnBr₄ also remain unchanged (Supplementary Fig. 21b), indicating the good\nthermal stability of the PipZnBr sample. Even when the temperature increases to 300 ℃, the\n4\nweight of the sample can remain at 96.5wt%. These results demonstrate that the PipZnBr\n4\nmaterial possesses excellent stability within the range of 300 ℃.\nSubsequently, the synthesized PipZnBr was pressed into a 16 mm diameter pellet with a\n4\nthickness of 430 μm to assess its ionic conductivity (Fig. 4d). The experiment of ion\nconductivity has been tested within the temperature range of 0 ℃ to 60 ℃ with 5 ℃ intervals.\nAs shown in Supplementary Figs. 22a−m, the Nyquist plot of PipZnBr pellet as a function\n4\nof temperature was exhibited, and the electrolyte underwent three EIS tests at each temperature.\nSupplementary Fig. 22n shows the corresponding ionic conductivity at different temperatures.\nNotably, the high ionic conductivity values of 0.29×10−3 S cm−1 and 0.47 ×10−3 S cm−1 at 25 ℃\nand 40 ℃ were observed for the PipZnBr pellet. Its Arrhenius curve was plotted in Fig. 4e to\n4\ncalculate the activation energy of the PipZnBr electrolyte. Arrhenius plot of PipZnBr exhibits\n4 4\ngood linearity, and the derived activation energy is 0.13 eV.\nWe have supplemented the thermal gravimetric analysis of PipZnBr and updated the\naccordingly.\nRevisions in the manuscript:\n“… a high ionic conductivity of 2.9×10−4 S cm−1 at 25 ℃…”\n“Additionally, the weight and XRD of PipZnBr remain almost unchanged when the\n4\ntemperature rises to 150 ℃ in the thermal gravimetric test (Supplementary Fig. 21). Even\nwhen the temperature increases to 300 ℃, the weight of the sample can remain at 96.5wt%,\nindicating the excellent thermal stability of PipZnBr .”\n4\n“… pellet with a thickness of ~430 μm as a solid-state zinc-ion electrolyte for zinc-ion\nbatteries (Fig. 4d). … at different temperatures with 5 ℃ intervals by … based on fitting\nNyquist plots (Supplementary Figs. 22a−m). … 0 ~ 60 ℃, are depicted in Supplementary\nFig. 22n. … Notably, the high σ values of 2.9 and 4.7 ×10−4 S cm−1 at 25 ℃ and\n40 ℃ …activation energy (0.13 eV). …”\n“ To verify excellent ionic conductivity of PipZnBr , the ionic conductivity of ZnBr has\n4 2\nbeen assessed at 25 ℃ (Supplementary Fig. 23). Meanwhile, PipZnBr also exhibits a\n4\npotential range …”\nFig. 4 | d Photograph of PipZnBr . e Arrhenius plots of PipZnBr as a function of temperature\n4 4\nfrom 0 ℃ to 60 ℃.\nRevisions in the supplementary information:\nSupplementary Fig. 21│ Thermal stability of PipZnBr crystal. a TG curve of PipZnBr\n4 4\nsample from room temperature to 800 ℃. b XRD patterns of PipZnBr₄ after heating at 150 ℃\nfor 30 min (red) and simulated XRD patterns based on the cifs file of CCDC 284070 (black).\nSupplementary Fig. 22 | Electrochemical characterization of the synthesized PipZnBr\n4\nsolid-state electrolyte. a−m Nyquist plot of PipZnBr electrolyte as a function of temperature\n4\nfrom 0 to 60 ℃. n Ionic conductivity of PipZnBr electrolyte at different temperatures.\n4\n4. Besides the concern about the conduction of ionic species, the electrochemical process in the\nfabricated cells needs careful analysis and more convincing evidence. It is necessary to identify\nthe real conducting ions in the reported electrolytes. A Zn2+ block layer for the impedance test\nshould be employed. In addition, chloride based electrolytes should also be tested in the\nsymmetric cells and full cells. In particular, the charge/discharge curves of chloride-type\nelectrolyte based full cell can reveal the details of electrochemical reactions in the cell.\nResponse: We sincerely appreciate the reviewer’s insightful suggestion. Following the\nsuggestions, we have tested the resistance of two SS||PipZnBr ||SS symmetrical cells by adding\n4\na single-layer PP separator on the electrolyte side and both sides. The PP membrane with a\nsmall pore diameter is often used as a separator in lithium-ion batteries. The resistance curves\ndiffer from SS||PipZnBr ||SS symmetrical cell without adding PP separator (Fig. R2).\nPipZnBr electrolyte, indicating that the PP membrane as a Zn2+ block layer can prevent Zn2+\n4\nmigration in the PipZnBr electrolyte. This test result suggests that the Zn2+ ion conduction\n4\nplays the dominant role in the halide-based zinc-ion batteries.\nFig. R2 | EIS curves of SS||PipZnBr ||SS symmetrical cell with adding PP separator on the\n4\nelectrolyte side and both sides.\nFurthermore, the cycling performance of Zn plating/stripping in Zn||Zn and Zn||I cells\n2\nusing PipZnCl as a solid-state electrolyte was tested. As shown in Supplementary Fig. 47a,\n4\nthe Zn||PipZnCl ||Zn cell can run over 600 h at 0.1 mA cm−2 and 0.1 mAh cm−2 at 40 ℃. The\n4\ncorresponding polarization potential remains almost unchanged with the increase of the cycling\ntime (Supplementary Fig. 47b), attributed to the stable Zn plating/stripping. Additionally, the\nZn||PipZnCl ||I full cell was assembled. As shown in Supplementary Fig. 47c, the full cell\n4 2\ncan operate about 70 cycles at 0.2 mA cm−2 at 40 ℃. In particular, the charge/discharge curves\nof Zn||PipZnCl ||I full cell exhibit the redox potential at ~1.2 V (Supplementary Fig. 47d),\n4 2\ncorresponding to the conversion of I to I−, which is consistent with a bromine-type electrolyte-\n2\nbased full cell. These test results suggest that the composition engineering strategy is versatile\nfor constructing halide zinc-ion electrolytes. Based on the above analysis, it is found that the\nreal conducting ion is Zn2+ in the halide-based solid-state zinc-ion batteries.\nIn the revised version, we have supplemented the electrochemical performance of a\nchloride-type electrolyte-based full cell. Finally, we have revised the manuscript and the\nsupplementary information accordingly.\nRevisions in the manuscript:\n“… More significantly, as shown in Supplementary Fig. 47, PipZnCl as electrolyte was\n4\nalso evaluated in a Zn symmetric cell and a Zn||I full cell. These results demonstrate that the\n2\nreal conducting ion is Zn2+ in the halide-based solid-state zinc-ion batteries.”\nRevisions in the supplementary information:\nSupplementary Fig. 47│Electrochemical characterization of the PipZnCl -type\n4\nelectrolyte. a Cycling performance of Zn||PipZnCl ||Zn cell at current density of 0.1 mA cm−2\n4\nand plating capacity of 0.1 mAh cm−2 at 40 ℃. b Voltage curves of Zn symmetrical cells. c\nCycling performance of Zn||PipZnCl ||I cell at 0.2 mA cm−2. d Corresponding to the\n4 2\ncharge/discharge curves of the full cell.\n“As shown in Supplementary Fig. 47a, the Zn||PipZnCl ||Zn cell can run over 600 h at 0.1 mA\n4\ncm−2 and 0.1 mAh cm−2 at 40 ℃. The corresponding polarization potential remains almost\nunchanged with the increase of the cycling time (Supplementary Fig. 47b), attributed to the\nstable Zn plating/stripping. Additionally, the Zn||PipZnCl ||I full cell was assembled. As shown\n4 2\nin Supplementary Fig. 47c, the full cell can operate about 70 cycles at 0.2 mA cm−2 at 40 ℃.\nIn particular, the charge/discharge curves of Zn||PipZnCl ||I full cell exhibit the redox potential\n4 2\nat ~1.2 V (Supplementary Fig. 47d), corresponding to the conversion of I to I−, which is\n2\nconsistent with a bromine-type electrolyte-based full cell. These test results suggest that the\ncomposition engineering strategy is versatile for constructing halide zinc-ion electrolytes.”\n5. There are many minor errors in the manuscript. For example, Fig. 2j lacks the DOS data for\nCl 3p; Line 231: an incorrect full name of CI-NEB; Line 453: duplicate \"PipZnBr \"; Line 463:\n4\nincorrect activation energy value. These errors could be avoided if the authors prepare the\nmanuscript in a well-trained and multiple-checked way.\nResponse: We deeply appreciate the reviewer for their meticulous reading of our manuscript\nand for identifying these minor errors. We sincerely apologize for these oversights. Following\nthe errors pointed out by the reviewer, we have carefully revised them one by one and also\nchecked for some other minor errors in the entire manuscript, listed as follows.\n(1) In Figs. 2f and 2j, the identification of the DOS for Br p should be Cl p, the\nidentification was modified in the modified version. To maintain consistency with the main text,\nwe further specified their orbital periods.\nRevisions in the manuscript:\nFig. 2 | e−l, Calculated (e, i) band structure, (f, j) total and projected DOSs, local charge density\nat (g, k) CBM and (h, l) VBM of PipZnBr and PipZnCl , respectively. The zero in band\n4 4\nstructures and DOSs is denoted as the VBM. The insets in f and j represent projected DOSs\nwhich have been enlarged by 10 times for clearer indication of their orbital contributions. The\nisosurfaces of local charge density are all set to 0.001 e Bohr−3\n.\n(2) In Line 231, the full name of Cl-NEB has been corrected.\nRevisions in the manuscript: “… climbing image nudged elastic band (Cl-NEB) …”\n(3) In Line 453, the subsequent duplicate PipZnBr has been modified to PipZnCl .\n4 4\nRevisions in the manuscript: “… i,e, PipZnBr and PipZnCl .”\n4 4\n(4) In Line 463, the activation energy value has been corrected.\nRevisions in the manuscript: “…low activation energy (0.13 eV) …”\n(5) In Line 188, the typos and the missing in the text of Figs. 2a and 2d have been corrected.\nRevisions in the manuscript:\nFig. 2 | b, d Chemical phase diagrams of (b) PipZnBr and (d) PipZnCl .\n4 4\n(6) In the supplementary information, some typos have been corrected.\nRevisions in the supplementary information:\n“ … Mn-doped C H N ZnX (X = Cl, Br) for X-ray …”\n4 12 2 4\n“… Li MCl (M = Tb−Lu, Y, Sc): Synthese …”\n3 6\n“… Li MBr (M=Sm−Lu, Y): Synthese …”\n3 6\n“… incommensurate (INC) Rb ZnBr phases.”\n2 4\nFinally, thank you again to the reviewer for their valuable comments on our manuscript,\nwhich are of great help to improve the quality of our manuscript.\nReviewer #2\nHu et al. proposed an interesting design rule for Zn-based metal halide electrolytes by replacing\ninorganic atoms with organic clusters. The hypothesis is that organic clusters may provide more\nabundant defect properties and thereby enhance the kinetics of zinc-ion transport. They\ninvestigated the hypothesis both computationally and experimentally. Guided by computational\nresults, they synthesized PipZnBr , which shows high conductivity. Overall, I found the work\n4\ninteresting and relevant for the battery community but there are a few technical concerns that\nneed to be addressed before being considered for publication in Nature Communication.\nResponse: We sincerely appreciate Reviewer #2 for his/her positive assessment of the novelty\nand relevance of our work to the battery community. We also thank the reviewer for his/her\ninsightful technical concerns and suggestions on readability and presentation, which can\nsignificantly improve the quality of our manuscript. We are trying to address these concerns\npoint-by-point as follows.\nTechnical concerns:\n1. Applicability of BVSE can be limited and material-dependent, the authors need to provide a\ndiscussion on the limitation of the approach and how it can still be used.\nResponse: We greatly appreciate the reviewer’s concern. In response to the concern about the\napplicability of the bond valence site energy (BVSE) method, we will proceed with a detailed\ndiscussion in two aspects:\n(1) The BVSE method is an effective simulation method that can predict the ion migration\nbehavior based on the crystal structure. When compared to the traditional DFT method, BVSE\nprovides a faster determination of ion migration barriers in materials. As a result, it is widely\nutilized in the field of solid-state batteries. The BVSE method boasts broad applicability and is\nsuited for materials containing cations such as Li+, Na+, K+, Ag+, Cu2+, Mg2+, Zn2+, Al3+, etc.\nFor example, Wang et al. employed the BVSE method to investigate the migration of Li ions\non chloride solid electrolytes Li ZrCl (Nat. Comm., 2021, 12(1): 4410). Nishitani et al.\n2 6\nconducted a study on the migration of Mg ions across six types of inorganic oxides (Solid State\nIonics, 2018, 315: 111-115), while Lv et al. examined Zn ion migration in ZnPS by using the\nmethod, Zhang et. Al established a database of ionic transport characteristics for over 29000\ncompounds, including Li+, Na+, K+, Ag+, Cu2+, Mg2+, Zn2+, Ca2+, Al3+, F−, and O2− (Adv. Funct.\nMater., 2020, 30(35): 2003087).\n(2) To further address the reviewers' concerns regarding the applicability range of the\nBVSE method, we performed complementary DFT calculations on the materials under\ninvestigation to enable mutual verification between the two methods. For Li-based compounds,\nwe selected Li₂S as a representative example, which has been reported in the reference (Nature\nMater., 2015, 14(10): 1026-1031). Similarly, Cs₂ZnBr₄ was chosen as the representative Zn-\nbased compound. As shown in Fig. R3, the BVSE calculation results show the Li⁺ migration\nbarriers of 0.30 eV for Li₂S and Zn²⁺ migration barriers of 1.09 ~1.32 eV for Cs₂ZnBr₄. As a\ncontrast, DFT results (Fig. R4) show corresponding barriers of 0.32 eV and 1.36 eV,\nrespectively. The close agreement between the BVSE and DFT migration barriers for both Li⁺\nand Zn²⁺ ions demonstrate the applicability of the BVSE method.\nFig. R3 | BVSE calculated Li/Zn-ion energy migration pathways. a Li S. b Cs ZnBr .\n2 2 4\nFig. R4 | Migrations schematic diagrams of Li/Zn-based examples and their DFT-calculated\nLi/Zn-ion energy migration pathways. a Li S. b Cs ZnBr .\n2 2 4\n2. The author assumed that Pip is one “component” like Zn and Br/Cl when predicting chemical\nphase diagrams and defect formation energy. But this assumption becomes questionable when\nthe results of defect formation enthalpy for point D1/D2 were shown, especially the formation\nenthalpy is very negative (at the order of 1 eV)\nResponse: We sincerely thank the reviewer for pointing out this important issue. Following the\nreviewer’s concern, we conducted a thorough examination of the entire phase diagram and the\ndefect calculation process to identify the reason for the negative formation enthalpy at point\nD1/D2.\nFirstly, we rechecked the structures of all the possible secondary phases from the databases.\nDuring our research of the Inorganic Crystal Structure Database (ICSD) and the Cambridge\nCrystallographic Data Centre (CCDC) database for the possible secondary phases, we found\nthat the possible phase structures of PipBr and PipCl had been missed in our previous\n2 2\ncalculations. We sincerely apologize for this oversight. Therefore, we rechecked all of the\ncalculated total energies of PipZnBr , PipZnCl , and their possible secondary phases, and then\n4 4\ndrew new chemical phase diagrams of PipZnBr and PipZnCl (Figs. 2b and 2d).\n4 4\nSecondly, based on the new phase diagrams of PipZnBr and PipZnCl , we recalculated\n4 4\nthe calculated formation enthalpies (ΔH) of intrinsic defects in PipZnBr and PipZnCl as a\n4 4\nfunction of the Fermi level (E ) at different chemical potential points (Figs. R5a−h). The\nF\nrevised results demonstrate that the exceptionally negative formation energy in D1/D2 has\nceased to exist (Figs. R5d and R5h). We gratefully thank the reviewer again for his/her astute\nobservation of this issue, which has allowed us to recognize and rectify the oversight.\nFig. R5 | a−h Calculated formation enthalpies (ΔH) of intrinsic defects in (a−d) PipZnBr and\n4\n(e−h) PipZnCl as a function of Fermi level (E ) at different chemical potential points: (a) point\nA , (b) point B , (c) point C , (d) point D , (e) point A , (f) point B , (g) point C , and (h) point\n1 1 1 1 2 2 2\nD . Defects with very high ΔH are not shown in the figures. The calculated equilibrium E (E )\nis denoted by vertical black dash lines.\nThirdly, according to the technical concerns No.5 of the reviewer (The detail will be\ndiscussed in the response to the concerns No.5), we have considered the band edge correction\nand subsequently obtained the final results, which are depicted in Figs. 3a−h.\nWe gratefully thank the reviewer again for his/her astute observation of this issue, which\nhas allowed us to recognize and rectify the oversight. We have revised the manuscript and the\nsupplementary information accordingly.\nRevisions in the manuscript:\n“Among the calculated nine intrinsic point defects, V , Br/Cl V , V , and Br /Cl are\nBr/Cl i i Zn Pip Zn Zn\nthe dominant defects, which determine the electrical properties of PipZnBr and PipZnCl .\n4 4\nUnder the Zn-rich and Br/Cl-poor condition (Point A , and A ), the primary donor defect is\n1 2\nhalogen vacancy (V /V ), while the domain acceptor defects are V , V , and Br/Cl,\nBr Cl Pip Zn i i\nindicating complex defect chemistry in this condition. The equilibrium Fermi levels (E₋F,e) for\nPipZnBr₄ and PipZnCl₄ at Points A /A are 3.61 eV and 4.27 eV above the VBM, respectively\n1 2\n(Figs. 3a and 3e). As the chemical potential gradually shifts from Br-poor to Br-rich (from\npoint A to D ), the E of PipZnBr moves correspondingly from near the VBM towards the\n1 1 F,e 4\nCBM (Figs. 3a−d). Similarly, for PipZnCl , as the chemical potential changes from Cl-poor to\n2 2 F,e\nno matter how the E of PipZnBr and PipZnCl changes, it is far away from the band edge,\nF,e 4 4\nand the corresponding hole density (N ) or electron density (N ) is very low. (Supplementary\nh e\nTable 6 and Supplementary Table 7.) These results further show the insulating nature of these\ntwo materials. It should be noted that the dominant donor defects are always the halogen\nvacancy defects (V for PipZnBr and V for PipZnCl ), while the dominant acceptor defects\nBr 4 Cl 4\nchange with different chemical potential. The dominant halogen vacancy defects suggest that\nboth materials are susceptible to breaking the original zinc-halogen tetrahedral structure due to\nthese defects, which may provide more space for zinc ion conduction.”\nFig. 2 | a, c Primitive cells of (a) PipZnBr and (c) PipZnCl . b, d Chemical phase diagrams of\n4 4\n(b) PipZnBr and (d) PipZnCl .\n4 4\nFig. 3 | a−h Calculated formation enthalpies (ΔH) of intrinsic defects in (a−d) PipZnBr and\n4\n(e−h) PipZnCl as a function of Fermi level (E ) at different chemical potential points: (a) point\nA , (b) point B , (c) point C , (d) point D , (e) point A (f) point B , (g) point C , (h) point D .\n1 1 1 1 2 2 2 2\nDefects with very high ΔH are not shown in the figures. The calculated equilibrium E (E ) is\nF F,e\ndenoted by vertical black dash lines.\nRevisions in the supplementary information:\nSupplementary Fig. 13 | Calculated charge-state transition levels for the calculated\nintrinsic defects in PipZnBr . The red lines on the left are acceptor-like defects while the blue\n4\nlines on the right are donor-like defects. Zn , which does not display characteristics of either\nPip\nacceptor-like or donor-like defects, is not illustrated in the figure.\nSupplementary Fig. 14 | Calculated charge-state transition levels for the calculated\nintrinsic defects in PipZnCl . The red lines on the left are acceptor-like defects while the blue\n4\nlines on the right are donor-like defects. Zn , which does not display characteristics of either\nPip\nacceptor-like or donor-like defects, is not illustrated in the figure.\nSupplementary Table 3 | Crystal system, space group, and GGA-PBE calculated total energies\n(E) of PipZnBr , PipZnCl and their possible secondary phases.\nt 4 4,\nFormula Crystal system Space group Reference E (eV/f.u.)\nt\nPipZnBr Orthorhombic P2 2 2 (19) 15 −110.655\n4 1 1 1\nPipZnCl Orthorhombic P2 2 2 (19) 16 −112.985\n4 1 1 1\nZn Hexagonal P6 /mmc (194) 17 −1.105\n2 1\nZnCl Tetragonal I42d (122) 19 −8.232\n4 2 12\nPipBr Monoclinic C/2c (15) 21 −102.911\nPip Zn Br\ndiagram of PipZnBr (point A , point B , point C , and point D ).\n4 1 1 1 1\nChemical potential\nΔμ (eV) Δμ (eV) Δμ (eV)\nPip Zn Br\npoints\nA −4.53 0 −1.69\nPip Zn Cl\ndiagram of PipZnCl (point A , point B , point C , and point D ).\n4 2 2 2 2\nChemical potential\nΔμ (eV) Δμ (eV) Δμ (eV)\nPip Zn Cl\npoints\nA −4.53 0 −2.11\n21. Bujak M. Efficient Diffusion-controlled ligand exchange crystal growth of isostructural\ninorganic–organic halogenidorhodates (III): The missing hexaiodidorhodate (III) anion.\nCryst. Growth Des., 15, 1295–1302 (2015).\n22. Dehghanpour S, Mahmoudkhani A H, Langer V. Redetermination of piperazine\ndihydrochloride monohydrate. Struct. Rep., 57, o610–o611 (2001).\n3. Related to point 2 above, the formation energy of C, H, N vacancies can be used to test if the\nassumption is valid. In addition, what about C/H/N interstitials? Those can be charged defects\nand move the equilibrium Fermi energy.\nResponse: We greatly appreciate the reviewer’s valuable comments and suggestions. Following\nthe reviewer’s suggestion, we calculated the vacancies and interstitial defects of C, H, and N to\ninvestigate their effects on the defect equilibrium Fermi level. The calculation results of\nvacancy and interstitial defects (C, H, N) are as follows:\n(1) As shown in Figs. R6 (before band edge correction) and R7 (after band edge\ncorrection), we have calculated the ΔH of intrinsic defects in PipZnBr and PipZnCl as a\n4 4\nfunction of E at Points A −D / A −D including the C, H, and N vacancies (V , V , and V ).\nF 1 1 2 2, C H N\nThe results show that for both PipZnBr and PipZnCl at different chemical potential points,\n4 4\nthe formation enthalpies of V , V , and V are higher than those of the dominant defects.\nC H N\nTherefore, the V , V , and V defects do not affect the defect properties of PipZnBr and\nC H N 4\nPipZnCl .\nincluding the C, H, and N vacanices: (a) point A , (b) point B , (c) point C , (d) point D , (e)\n1 1 1 1\npoint A (f) point B , (g) point C , (h) point D . The calculated equilibrium E (E ) is denoted\n2 2 2 2 F F,e\nby vertical black dash lines. The color dash lines successively represent the ΔH of C, H, and N\nvacancy defects.\nFig. R7 | a−h Calculated formation enthalpies (ΔH) of intrinsic defects in PipZnBr and\nincluding the C, H, and N vacanices: (a) point A , (b) point B , (c) point C , (d) point D , (e)\n1 1 1 1\npoint A (f) point B , (g) point C , (h) point D . The calculated equilibrium E (E ) is denoted\n2 2 2 2 F F,e\nby vertical black dash lines. The color dash lines successively represent the ΔH of C, H, and N\nvacancy defects.\n(2) We further calculated ΔH of intrinsic defects in PipZnBr and PipZnCl as a function\n4 4\nof E at points A −D / A −D , including C, H, and N interstitial defects (C, H, and N), as\nF 1 1 2 2 i i i\nshown in Figs. R8 (before band edge correction) and R9 (after band edge correction). The\nresults similarly indicate that for both PipZnBr and PipZnCl at the tested chemical potential\n4 4\npoints, the formation enthalpies of C, H, and N are higher than those of the dominant defects.\ni i i\nGiven their significantly higher ΔH, C, H, and N defects consequently do not influence the\ni i i\ndefect properties of PipZnBr₄ and PipZnCl₄.\nFig. R8 | a−h Calculated formation enthalpies (ΔH) of intrinsic defects in PipZnBr and\nincluding the C, H, and N interstitial defects: (a) point A , (b) point B , (c) point C , (d) point\n1 1 1\nD , (e) point A (f) point B , (g) point C , (h) point D . The calculated equilibrium E (E ) is\n1 2 2 2 2 F F,e\ndenoted by vertical black dash lines. The color dash lines successively represent the ΔH of C,\nH, and N interstitial defects.\nFig. R9 | a−h Calculated formation enthalpies (ΔH) of intrinsic defects in PipZnBr and\nincluding the C, H, and N interstitial defects: (a) point A , (b) point B , (c) point C , (d) point\n1 1 1\nD , (e) point A (f) point B , (g) point C , (h) point D . The calculated equilibrium E (E ) is\n1 2 2 2 2 F F,e\ndenoted by vertical black dash lines. The color dash lines successively represent the ΔH of C,\nH, and N interstitial defects.\nIn PipZnBr and PipZnCl , the piperazine used as the initial raw material is an organic\n4 4\nwhole. Upon formation of a compound, the Pip2+ entity exhibits a +2 valence state.\nConsequently, the elements of C, H, and N exhibit a more covalent nature, rendering their\nprecise valence state somewhat ambiguous. As a result, the high ΔH of C, H, and N vacancy\nand interstitial defects is reasonable. Given this understanding, coupled with the\naforementioned calculation results, the validity of our current calculation has been further\nconfirmed. We express our gratitude to the reviewers for their verification methods and valuable\nsuggestions.\n4. Since the band edge/gap correction is not included in the defect formation enthalpy, the\nuncertainty of the predicted values is unknown and thus not strong enough to fully support the\ndiscussion. The authors need to either add the correction to the results or discuss the limitation.\nResponse: We sincerely thank the reviewers for their insightful feedback. In the original\nmanuscript, we inadvertently overlooked band edge correction. Consequently, the reported\ndefect properties were influenced by the inherent band gap underestimation associated with the\nGGA method. We appreciate the reviewer's valuable suggestion to address this methodological\nlimitation. Based on this suggestion, we have incorporated band edge correction to determine\nthe final defect formation energies and transition levels.\nOur prior study (Chem. Eng. J. 468 (2023) 143818) has demonstrated that the standard\nHSE method calculated band gaps (PipZnBr and PipZnCl ) provide very good agreement with\n4 4\nexperimental results. As a result, we apply the band edge correction based on the standard HSE\nmethod in the revised version. As illustrated in Fig. R10, the GGA-PBE functional\nsystematically overestimates the valence band maximum (VBM) while underestimating the\nconduction band minimum (CBM). Specifically, the VBM is overestimated by 0.89 eV with\nCBM underestimated by 0.41 eV in PipZbBr , whereas for PipZnCl , the VBM overestimation\n4 4\nis 1.02 eV and CBM underestimation is 0.44 eV. To address this problem, we applied the band\nedge correction based on the previous work (Adv. Funct. Mater. 25 (2015), 6733–6743), and\nobtained the final results as shown in Figs. 3a−h.\nFig. R10 | a−b Schematic illustration of the alignment between band edges obtained with the\nGGA-PBE method and the HSE method.\nWe have revised the manuscript, methods, and supplementary information accordingly.\nRevisions in the manuscript:\nFig. 3 | a−h Calculated formation enthalpies (ΔH) of intrinsic defects in (a−d) PipZnBr and\n4\n(e−h) PipZnCl as a function of Fermi level (E ) at different chemical potential points: (a) point\nA , (b) point B , (c) point C , (d) point D , (e) point A (f) point B , (g) point C , (h) point D .\n1 1 1 1 2 2 2 2\nDefects with very high ΔH are not shown in the figures. The calculated equilibrium E (E ) is\nF F,e\ndenoted by vertical black dash lines.\nRevisions in methods of the manuscript:\n“…including the band-filling correction, the potential alignment correction, and the image\ncharge correction61,62. To address the band edge problem using the GGA-PBE method59,63,\ncorrections were applied for both the valence band maximum (VBM) overestimation and the\nconduction band minimum (CBM) underestimation. For PipZnBr , the corrections are ΔVBM\n4\n= 0.89 eV and ΔCBM = 0.41 eV, while for PipZnCl , they are ΔVBM = 1.02 eV and ΔCBM =\n4\n0.44 eV. These corrections are incorporated by adding ΔVBM (or ΔCBM) to the formation\nenergy of deep acceptor (or donor) defects calculated using the HSE hybrid functional method.\nFor shallow defects, the formation energy is adjusted to maintain its relative position to the\nedge. The charge transition level and the equilibrium Fermi level (E ) were determined using\nF,e\nmethods identical to those employed in our previous research64.”\nReference added and changed in methods of the manuscript:\n63. Yuan, Z. et al. Engineering solar cell absorbers by exploring the band alignment and defect\ndisparity: the case of Cu‐and Ag‐based kesterite compounds. Adv. Funct. Mater. 25, 6733–\n6743 (2015).\n64. Hu, S., Xia, B., Yan, Y. & Xiao, Z. Lead chloride perovskites for p-type transparent\nconductors: A critical theoretical reevaluation. Phys. Rev. Materials 4, 115201 (2020).\nRevisions in the supplementary information:\nSupplementary Fig. 13 | Calculated charge-state transition levels for the calculated\nintrinsic defects in PipZnBr . The red lines on the left are acceptor-like defects while the blue\n4\nlines on the right are donor-like defects. Zn , which does not display characteristics of either\nPip\nacceptor-like or donor-like defects, is not illustrated in the figure.\nSupplementary Fig. 14 | Calculated charge-state transition levels for the calculated\nintrinsic defects in PipZnCl . The red lines on the left are acceptor-like defects while the blue\n4\nlines on the right are donor-like defects. Zn , which does not display characteristics of either\nPip\nacceptor-like or donor-like defects, is not illustrated in the figure.\n5. One of the main conclusions from computational results is that PipZnBr /PipZnCl can be\n4 4\ngood insulator based on the large band gap and defect calculations. While the results of\nequilibrium Fermi energy being at least about 1 eV away from band edges can support the\nargument, a more quantitative analysis on the number of electron/hole carriers as a function of\nsynthesis temperature can be a stronger support.\nResponse: We greatly appreciate the reviewer for this excellent and insightful suggestion. The\nreviewer rightly points out that a quantitative analysis of charge carrier concentration can\nprovide stronger evidence for the insulating nature of PipZnBr₄ and PipZnCl₄. We fully agree\nwith this insightful suggestion. Following this suggestion, we added detailed calculations of the\nelectron (n) and hole (p) carrier concentrations as a function of temperature for both PipZnBr₄\nand PipZnCl₄. These calculations were carried out within the temperature range from 300 K to\nThese results reveal that both PipZnBr4 and PipZnCl4 exhibit ultra-low carrier concentrations\nat the equilibrium Fermi level over the temperature range spanning from ambient to synthesis\nconditions, reinforcing their intrinsic insulating properties.\nWe have revised the manuscript and the supplementary information accordingly.\nRevisions in the manuscript:\n“Impressively, no matter how the E of PipZnBr and PipZnCl changes, it is far away from\nF,e 4 4\nthe band edge, and the corresponding hole density (N ) or electron density (N ) is very low.\nh e\n(Supplementary Table 6 and Supplementary Table 7.)”\nRevisions in the supplementary information:\nSupplementary Table 6 | Calculated carrier concentrations as a function of synthesis\ntemperature at different chemical potential points of PipZnBr (point A , point B , point C , and\n4 1 1 1\npoint D ). Hole density (N ), electron density (N ).\n1 h e\nTemperature (K) E (eV) N (cm−3) / N (cm−3)\nF,e h e\nPoint A 3.61 N = 2.17 × 10−10\n1 e\nPoint B 3.20 N = 1.34 × 10−17\n1 e\n1 h\nPoint D 1.44 N = 1.28 × 10−6\n1 h\nPoint A 3.61 N = 2.17 × 10−10\n1 e\nPoint B 3.20 N = 1.40 × 10−17\n1 e\n1 h\nPoint D 1.44 N = 1.18 × 10−6\n1 h\nPoint A 3.61 N = 2.17 × 10−10\n1 e\n1 e\nPoint C 1.74 N = 6.99 × 10−12\n1 h\nPoint D 1.44 N = 1.14 × 10−6\n1 h\nPoint A 3.61 N = 2.26 × 10−10\n1 e\nPoint B 3.20 N = 1.45 × 10−17\n1 e\n1 h\nPoint D 1.44 N = 1.05 × 10−6\n1 h\nPoint A 3.61 N = 2.26 × 10−10\n1 e\nPoint B 3.20 N = 1.45 × 10−17\n1 e\n1 h\nPoint D 1.44 N = 1.01 × 10−6\n1 h\nNote: At each chemical potential point, we only give the higher carrier concentration in N or\nh\nN , because the concentration of the other carrier is smaller than the given concentration.\ne\nSupplementary Table 7 | Calculated carrier concentrations as a function of synthesis\ntemperature at different chemical potential points of PipZnCl (point A , point B , point C , and\n4 2 2 2\npoint D ).\nF,e h e\nPoint A 4.27 N = 1.84 × 10−15\n2 e\nPoint B 3.74 N = 1.05 × 10−24\n2 e\n2 h\nPoint D 1.60 N = 2.11 ×10−9\n2 h\nPoint A 4.27 N = 1.84 × 10−15\n2 e\nPoint B 3.74 N = 1.05 × 10−24\n2 e\n2 h\nPoint D 1.61 N = 2.02 ×10−9\n2 h\nPoint A 4.27 N = 1.84 × 10−15\n2 e\nPoint B 3.74 N = 1.10 × 10−24\n2 e\n2 h\nPoint D 1.61 N = 1.87 ×10−9\n2 h\nPoint A 4.27 N = 1.84 × 10−15\n2 e\nPoint B 3.74 N = 1.10 × 10−24\n2 e\n2 h\nPoint D 1.61 N = 1.80 ×10−9\n2 h\nPoint A 4.27 N = 1.84 × 10−15\n2 e\nPoint B 3.74 N = 1.14 × 10−24\n2 e\n2 h\nPoint D 1.61 N = 1.66 ×10−9\n2 h\nNote: At each chemical potential point, we only give the higher carrier concentration in N or\nh\nN , because the concentration of the other carrier is smaller than the given concentration.\ne\n6. For NEB calculations, how are the initial path determined? In other words, how can we know\nthis is the minimum energy path among all the possible pathways.\nResponse: We sincerely appreciate the insightful and constructive feedback provided by the\nreviewers. The reviewer’s thoughtful comments and suggestions have offered a significantly\nmore comprehensive view of our manuscript, greatly enhancing the depth and rigor of our work.\nIn our original manuscript, we only showed the Zn2+ migration pathways of PipZnBr and\n4\nthe reviewer’s suggestion, in the revised version, we have added the migration pathways of\nPipZnBr and PipZnCl along the b-axis and c-axis of the crystals (Supplementary Figs. 15\n4 4\nand 16). The results show that both PipZnBr and PipZnCl have the minimum energy path for\n4 4\nZn2+ migration in the a-axis direction.\nWe have revised the manuscript and the supplementary information accordingly.\nRevisions in the manuscript:\n“To evaluate the Zn2+ mobility in PipZnBr and PipZnCl , the Zn ion pathways for a Zn atom\n4 4\nconfined with two materials were investigated by the climbing image nudged elastic band (Cl-\nNEB) method30. Exploration of migration pathways along all three axis directions revealed that\nthe lowest energy barriers occur along the a-direction (Figs. 3i−l, Supplementary Figs. 15 and\n16). As shown in Figs. 3i and 3k, the migration of the Zn-ion along the a-axis occurs almost in\nthe direction of the calculated supercells, specifically through the gap between the [ZnX ]2−\n4\nunits and the organic cations. The corresponding energy barriers for Zn ion diffusion along the\na-direction in PipZnBr and PipZnCl are 0.37 eV (Fig. 3j) and 0.39 eV (Fig. 3l), respectively,\n4 4\nwhich means Zn ions can easily migrate in the two materials. The migration results of Zn2+ in\nPipZnBr and PipZnCl indicate that both materials exhibit excellent Zn-ion conductivity. The\n4 4\nabove defect properties and the Zn2+ migration calculation results suggest that both materials\nhave good insulation properties and excellent Zn-ion conductivity, which shows great potential\nas superior solid-state zinc-ion electrolytes.”\nRevisions in the supplementary information:\nSupplementary Fig. 15 | Zn2+ migrations of PipZnBr and PipZnCl along b-axis. a, c The\n4 4\nZn2+-migration schematic diagrams of (a) PipZnBr and (b) PipZnCl . b, d The corresponding\n4 4\nZn2+-migration paths of (c) PipZnBr in a and (d) PipZnCl in c, respectively.\n4 4\nSupplementary Fig. 16 | Zn2+ migrations of PipZnBr and PipZnCl along c-axis. a, c The\n4 4\nZn2+-migration schematic diagrams of (a) PipZnBr and (b) PipZnCl . b, d The corresponding\n4 4\nZn2+-migration paths of (c) PipZnBr in a and (d) PipZnCl in c, respectively.\n4 4\n7. “The broken tetrahedra leaves a low migration energy barrier for Zn ions in the material,…”\nIt is unclear whether broken tetrahedra is included in the NEB calculations. If not, the statement\nis not fully supported\nResponse: We sincerely appreciate the reviewer's insightful comment regarding the relationship\nbetween broken tetrahedra and Zn-ion migration energy barriers. The reviewer correctly notes\nthat our original statement required stronger computational support, as our NEB calculations\nwere performed on representative diffusion pathways in the overall structure rather than\nspecifically isolating configurations with broken tetrahedra. To ensure full rigor, we have\nrevised the relevant statements in the revised manuscript.\nRevisions in the manuscript:\n“The above defect properties and the Zn2+ migration calculation results suggest that both\nmaterials have good insulation properties and excellent Zn-ion conductivity, which shows great\npotential as superior solid-state zinc-ion electrolytes.”\n8. For XRD results, do the authors identify competing phases? This can provide a hint whether\nthe material is near the phase boundary, i.e. close to energy of the A,B,C points shown in Fig\n2b.\nResponse: We sincerely thank the reviewer for raising this important comment. Following the\nsuggestion, we have conducted a reconfirmation of the theoretically calculated phase diagram\nand then identified the competitive phase based on the XRD results.\n(1) Based on the response to the reviewer’s comment #2 (Technical concerns), we have\nadded the database structures of PipBr and PipCl in the revised phase diagrams. The\n2 2\ncalculated phase diagrams show that ZnBr /PipBr , ZnCl /PipCl are important competing\n2 2 2 2\nphases for PipZnBr and PipZnCl , respectively (Figs. 2b, d). Our subsequent XRD of the\n4 4\ncompeting phase is based on the experimental structure of these competing phases.\nFig. 2 | b, d Chemical phase diagrams of (b) PipZnBr and (d) PipZnCl .\n4 4\n(2) Then, we have identified the competitive phase by using experimental structures from\nICSD/CCDC databases, as shown in Fig. R11. The results confirm that both of the synthesized\nmaterials are the target structures, but XRD is closer to ZnBr /ZnCl . Considering our synthesis\n2 2\nsteps, which require the addition of a more than sufficient amount of ZnBr /ZnCl to inhibit the\n2 2\nformation of decomposition products, it is inferred that the final range of material synthesis is\nnear point BC (B C for PipZnBr and B C for PipZnCl ) on the phase diagram.\n1 1 4 2 2 4\nFig. R11 | XRD patterns in this work and ICSD/CCDC databases. a PipZnBr . b PipZnCl .\n4 4\n9. For the measured stability window, how does this compare to prediction (Fig 2b)? Related to\npoint 2 above, this seems the indicate the breakdown of the assumption that treats pip as one\ncomponent.\nResponse: We sincerely thank the reviewer for this insightful comment.\nThe measured stability window (Fig. 4f) in our manuscript refers to the electrochemical\nstability window, which is the range of electrode potentials at which a material does not undergo\noxidation or reduction. In contrast, our calculated phase diagram (Fig. 2d) reflects the\nthermodynamic stability window of the material. We would like to clarify that these two\nstability windows are two different concepts, so they cannot be compared.\nAs a consequence, the fact that the two stability windows (measured stability window in\nFig. 4f and the calculated phase diagram in Fig. 2d) are not the same does not indicate that the\nhypothesis of PIp is invalid. In our manuscript, treating Pip as one component is a fundamental\nstrategy due to the following aspects\n(1) Following the reviewer's valuable suggestion in comment #3 (Technical concerns),\nwe have further validated the rationality of treating pip as an integrated unit. We performed\nadditional calculations explicitly considering C, N, and H vacancies and their interstitial defects.\nThese calculation results confirm that treating Pip as a component remains a valid and justified\nsimplification.\n(2) Based on the insightful suggestions of the reviewers in comment #8 (Technical\nconcerns), we have confirmed the viability of Pip as a whole unit from an experimental\nsynthesis perspective. In our target structures and the competing phases, Pip is always present\nas a unit. Additionally, in our experimental synthesis process, Pip is also added as a whole raw\nmaterial, confirming that Pip was introduced and behaves as a single component throughout\nmaterial formation.\nTherefore, we maintain that treating the Pip as a single component is a reasonable and\nwell-founded assumption. We deeply appreciate the reviewer again for their constructive\ncritique.\n10. Why is the measured activation barrier (~0.11 eV) much lower than the calculated migration\nbarrier?\nResponse: We sincerely thank the reviewer for this valuable comment. We agree that this\ndifference warrants careful discussion, and we appreciate the opportunity to elaborate on the\npotential reasons behind it. This difference may arise from fundamental distinctions between\nexperimental and theoretical methods, as well as material-specific factors.\nFirstly, the calculated migration energy barrier is estimated using an idealized crystal\nlattice under periodic boundary conditions. This approach is standard for isolating the\nfundamental migration mechanism and its intrinsic energy barrier. However, in real materials,\nespecially at elevated temperatures where diffusion occurs, the presence of multiple vacancies,\nvacancy clusters, or other extended defects is common. Secondly, the theoretical migration\nbarrier was calculated by using the first-principles method at 0K. In contrast, the experimental\nactivation energy is measured at finite temperatures. The temperature-dependent entropy and\nother factors are also one of the reasons for the inconsistency between theoretical and\nexperimental values. Consequently, while the calculated barrier serves as a qualitative indicator\nof a material’s ion migration capability, it should not be directly compared to experimental\nactivation energies in absolute values.\nIn addition, this phenomenon aligns with reports in solid-state lithium-ion electrolyte\nliterature. For example, Li Sc Zr Cl exhibits a calculated migration barrier of 0.47 eV\n2.75 0.75 0.25 6\nbut an experimental activation energy of 0.33 eV (ACS Energy Letters, 2024, 9(3): 1043-1052),\nfurther supporting the discussed origins of such discrepancies.\n11. Here is a general question since I see some connection between this work and hybrid\nperovskite. Can the authors comment on the stability of PipZnBr against water/moisture?\nstability of PipZnBr . To directly evaluate water/moisture resistance, we exposed a fresh\n2 hours. After that, we observed no visual degradation or dissolution, confirming retained\nmorphological integrity (Fig. R12). Furthermore, we verified structural stability through XRD\nanalysis, which revealed identical patterns to the pristine sample with no detectable hydrated\nphases (Fig. R13). These findings collectively demonstrate good stability of PipZnBr₄ against\nwater/moisture.\nFig. R12 | Optical microscopy image of PipZnBr₄ after 2-hour exposure to ambient air.\n10 20 30 40 50 60\nFig. R13 | XRD patterns of PipZnBr₄ after 2-hour exposure to ambient air (red) and\nsimulated XRD patterns based on the cifs file of CCDC 284070 (black).\n).u.a(\nytisnetnI\nAfter 2 hours\nPipZnBr -CCDC: 284070\n4\n2θ (deg)\nReadability/Presentation:\n1. In abstract, the importance of Zn-based battery is unclear\nResponse: We sincerely thank the reviewer for this valuable comment. We agree that explicitly\nstating the significance of Zn-based batteries in the abstract enhances the motivation for this\nwork. To address this, we have now added the following sentence to the beginning of the\nabstract, immediately after the first sentence.\nRevisions in the manuscript:\n“… advancement of batteries. While zinc-ion batteries (ZIBs) offer compelling advantages in\nsafety, cost, and environmental sustainability, their performance is constrained by conventional\naqueous electrolytes. All-solid-state zinc-ion electrolytes, …”\n2. In abstract, the statement, “low energy barriers between organic groups and halogen\ntetrahedra”, should be more precise. I understand this statement only after seeing Fig 3 i/k. I\noriginally thought this means Zn diffuse between organic groups and halogen tetrahedra.\nResponse: We greatly appreciate the reviewer for your careful reading of our manuscript and\nthe constructive comment regarding the phrasing in the abstract. We fully agree that the original\nstatement should be more precise. To address this, we have revised the sentence to clarify that\nlow energy barriers refer to the low migration energy barrier of Zn ions in the crystal structure.\nRevisions in the manuscript:\n“… Both solid-state electrolytes exhibit good insulating properties and low zinc-migration\nenergy barriers. …”\n3. In abstract, minor grammar issue for the last sentence “… a new approach to designing ….”\nResponse: We thank the reviewer for pointing out the potential concern regarding the phrase\n“... a new approach to designing ...”. We appreciate your careful reading. To ensure maximum\nclarity, we are happy to revise the sentence. We have implemented this change in the revised\nmanuscript.\nRevisions in the manuscript:\n“This work provides a new approach to the design of metal halide-based solid electrolytes for\nZIBs.”\n4. One general comment, the readership of nature communication is not limited to chemist.\nSome of the statements might be straightforward to chemist, but further elaboration can be\nhelpful for a broader audience.\nResponse: We sincerely thank the reviewer for this insightful comment. We fully agree that it\nis important to improve the readability of our work for interdisciplinary readers. In the revised\nmanuscript, we have modified certain technical phrases to enhance their accessibility, ensuring\nthat a broader audience can comprehend them while preserving scientific accuracy.\nRevisions in the manuscript:\n“Metal halide materials, a large family of materials composed of metal and halide ions (like\nchloride or bromide), have been widely used in various fields14–17.”\n“ When the coordination environments of Li and Zn ions with halogens are either 4-coordinated\nor 6-coordinated (Li/Zn surrounded by 4 or 6 halide atoms), their ionic radii are notably similar.”\n“This similarity implies that they occupy similar space in the crystal. To understand why these\nstructurally similar materials behave differently as electrolytes, a comparative analysis of the\ntetrahedral and octahedral structures of Li and Zn at the orbital level was performed to elucidate\ntheir similarities and differences. As shown in Supplementary Figs. 2−5, when Li and Zn\natoms bond with halogen atoms to form tetrahedral or octahedral structures, the nature of their\ninteraction differs fundamentally. The electron orbitals of Li and halogen atoms exhibit clear\nseparation, with minimal overlap between their electron clouds (Supplementary Figs. 2 and\n3). This result means Li and halogen atoms retain a non-bonding state in their elemental orbitals,\nwhich suggests that the predominant binding mode between Li ions and halogens is ionic bonds.\nConversely, when zinc ions coordinate with halogens to form tetrahedral or octahedral\nstructures, their interaction exhibits stronger orbital overlap (Supplementary Figs. 4 and 5).\nThis overlap electron cloud sharing generates distinct bonding and antibonding states, which\nmeans a covalent character in the Zn-halogen bond. (Supplementary Figs. 4 and 5).”\n“For PipZnBr , the Heyd–Scuseria–Ernzerhof (HSE) calculated band structure shows a large\n4\ndirect band gap of (5.37 eV) at the Γ point as shown in Fig. 2e, implying a poor electrical\nconductivity. The density of states (DOS) analysis (Fig. 2f) reveals that the valence band\nmaximum (VBM) is dominated by nonbonding states of Br 4p orbitals, whereas the adjacent\nvalence band (VB) region consists of antibonding states from Zn 3d and Br 4p hybridization.”\n5. In introduction, “Zinc ion possesses a high valence charge and mediocre electronegativity,\nmaking the study of solid electrolytes that utilize zinc ion transport as the primary conduction\nmechanism particularly challenging” the connection is not fully clear.\nResponse: We thank the reviewer for highlighting this lack of clarity. To address this lack of\nclarity and strengthen the logical flow, we have revised the relevant sentences in the revised\nmanuscript of the introduction section.\nRevisions in the manuscript:\n“Zn²⁺ ions present significant challenges for solid-state migration due to two key properties:\ntheir high charge valence (+2) and relatively low electronegativity (1.65). The high charge\ndensity associated with the +2 state induces strong electrostatic interactions with the anion\nlattice, resulting in high migration barriers. Simultaneously, the low electronegativity promotes\nstronger bonding with anions (e.g., O2⁻), which further restricts mobility. This dual impediment\nmakes zinc ion conduction particularly difficult to achieve in solid electrolytes. Recently,\ninorganic-based solid-state zinc-ion electrolytes have been developed…”\n6. In introduction, “These two materials show good electronic properties…” The statement is\nvague\nResponse: We thank the reviewer for pointing out the vagueness in the statement. We fully\nagree that the phrase “good electronic properties” was imprecise in the original manuscript. We\nhave revised the sentence to clarify that the materials possess excellent insulating properties.\nRevisions in the manuscript:\n“These two materials exhibit good insulating properties and low zinc-ion migration ...”\n7. For Fig. S2 and S3, the meaning of red and blue dots is not defined.\nResponse: We appreciate the reviewer’s careful attention to detail. We apologize for this lack\nof clarity in the figure captions. In the Supplementary Figs. 2−5, the red and blue dots signify\nspin-up and spin-down electrons, respectively. In addition, we have increased the caption\nrelated to the atomic and orbital electron density. These definitions have been prominently\nadded to the respective figure captions in the revised supplementary material.\nRevisions in the supplementary information:\nSupplementary Fig. 2 | Qualitative interaction diagram based on orbital analysis of\n[LiX ]3− (X = Halogen atom) tetrahedral cluster. a Li atom; b [LiX ]3− tetrahedral cluster; c\n4 4\n[X ]0 tetrahedral cluster. The red and blue dots in a, b, and c signify spin-up and spin-down\n4\nelectrons, respectively. The dashed green boxes inserted in b show the electron clouds of the\ncorresponding orbital of the [LiX ]3− tetrahedral cluster. Within the green boxes, bluish-gray\n4\natoms represent halogen atoms (X), green atoms signify Li atoms, and red/blue colors are orbital\nelectron density.\nSupplementary Fig. 3 | Qualitative interaction diagram based on orbital analysis of\n[LiX ]5− (X = Halogen atom) octahedral cluster. a Li atom; b [LiX ]5− octahedral cluster; c\n6 6\n[X ]0 octahedral cluster. The conventions for depicting spin electrons, atoms, and orbital\n6\nelectron densities in this figure are identical to those in Supplementary Fig. 2.\nSupplementary Fig. 4 | Qualitative interaction diagram based on orbital analysis of\n[ZnX ]2− (X = Halogen atom) tetrahedral cluster. a Zn atom; b [ZnX ]2− tetrahedral cluster;\n4 4\nc [X ]0 tetrahedral cluster. The red and blue dots in a, b, and c signify spin-up and spin-down\n4\nelectrons, respectively. The dashed green boxes inserted in b show the electron clouds of the\ncorresponding orbital of the [ZnX ]2− tetrahedral cluster. Within the green boxes, bluish-gray\n4\natoms represent halogen atoms (X), silvery gray atoms signify Zn atoms, and red/blue colors\nare orbital electron density.\nSupplementary Fig. 5 | Qualitative interaction diagram based on orbital analysis of\n[ZnX ]4− (X = Halogen atom) octahedral cluster. a Zn atom; b [ZnX ]4− octahedral cluster; c\n6 6\n[X ]0 octahedral cluster. The conventions for depicting spin electrons, atoms, and orbital\n6\nelectron densities in this figure are identical to those in Supplementary Fig. 4.\n8. “The localized charge of ionic bonds suggests that Li-based halides can conduct lithium\nions” elaboration is needed.\nResponse: We appreciate the reviewer’s insightful comment. In the original manuscript, this\nsentence preceded the valence bond analysis and was acknowledged to lack sufficient rigor.\nConsequently, the sentence has been removed in the revised version. Moreover, following the\nreviewer’s comment #9 (Readability/Presentation), we have added the COBI analysis,\nsubsequently generating the final results. To avoid duplication, detailed revisions will be shown\nin the subsequent reply to comment #9 (Readability/Presentation).\n9. On the related topic to the point 8 above, Crystal Orbital Bond Index (COBI) analysis could\nbe a good way to quantify the covalency. (optional)\nResponse: We thank the reviewer for the valuable suggestion regarding the use of the Crystal\nOrbital Bond Index (COBI) to quantify bond covalency. As recommended, we conducted a\ncomprehensive COBI analysis by using Li−Cl and Zn−Cl as representative cases. The COBI\nresults provide rigorous quantitative evidence characterizing the distinct bonding nature\nbetween Zn-Cl and Li-Cl interactions:\n1. Li−Cl exhibits consistently low COBI values (0.15 for 4-coordinated LiCl and 0.09 for 6-\ncoordinated LiCl), confirming ionic-dominated bonding with localized charges, which\nfacilitates Li⁺ ion liberation.\n2. Zn−Cl shows significantly higher COBI values (0.22 for 4-coordinated ZnCl and 0.18 for\n2\n6-coordinated ZnCl ), demonstrating relatively covalent character with delocalized\n2\nelectrons, thereby constraining Zn²⁺ migration.\nWe have revised the manuscript and the supplementary information accordingly.\nRevisions in the manuscript:\n“To quantify the interaction, an in-depth analysis of the crystal orbital bond index (COBI)21 was\nconducted for Li−Cl and Zn−Cl as representative cases. As illustrated in Supplementary Fig.\n6, the COBI values for 4-coordinated and 6-coordinated LiCl are 0.15 and 0.09, respectively,\nsuggesting a strong ionic character. In contrast, the COBI values for 4-coordinated and 6-\ncoordinated ZnCl are 0.22 and 0.18 (Supplementary Fig. 7), indicating a stronger covalent\n2\ninteraction between Zn and the halogens. These results suggest that the Li-based halides\ncharacterized by localized ionic bonding, which could facilitate lithium ion liberation, render\nthem potential candidates for solid-state lithium electrolytes. In contrast, the relatively strong\ncovalent bonding interactions in Zn-based halides may hinder Zn²⁺ ion migration. Consequently,\ntraditional Zn-based halides …”\nRevisions in methods of the manuscript:\n“Crystal orbital bond index (COBI) analysis. The structural models of 4-coordinated/6-\ncoordinated LiCl sourced from the Materials Project database51. The optimization calculations\nof the structures were performed using the projector-augmented wave (PAW)52 method, as\nincorporated in the Vienna Ab inito Simulation Package (VASP)53 code. The sutoff energy was\nset to 500 eV. Γ-centered k-meshes, with a k-spacing of 0.2 Å–1, were employed for the sampling\nof the Brillouin zones. The generalized gradient approximation (GGA) Perdew-Burke-\nErnzerhof (PBE)54 was used as the exchange-correlation functional. The energy convergence\ncriterion for COBI analysis is 10–7 eV. The COBI values were calculated by using the Lobster55\ncode.”\n“Density functional theory (DFT) calculations for electronic structures. DFT calculations\nwere performed using the VASP53 code. The cut-off energy, k-points, and exchange-correlation\nfunctional settings for all the structures (PipZnBr , PipZnCl , and their possible secondary\n4 4\nphases) optimizations were the same as those used in the structure optimization for COBI\nanalysis. …”\nReference added in the manuscript and methods:\n21. M üller, P. C., Ertural, C., Hempelmann, J., & Dronskowski, R. Crystal orbital bond index:\ncovalent bond orders in solids. J. Phys. Chem. C, 125, 7959–7970 (2021).\naccelerating materials innovation. APL Mater. 1, 011002 (2013).\nbonding analysis from projector‐augmented‐wave‐based density‐functional theory. J.\nComput. Chem. 41, 1931–1940 (2020).\nRevisions in the supplementary information:\nSupplementary Fig. 6 | Crystal structures and COBI results of different coordinated LiCl.\na 4-coordinated LiCl. b 6-coordinated LiCl. The crystal structures are obtained from the\nMaterials Project database2.\nSupplementary Fig. 7 | Crystal structures and COBI results of different coordinated\nZnCl . a 4-coordinated ZnCl . b 6-coordinated ZnCl . The crystal structures are obtained from\n2 2 2\nthe Materials Project database2. The 6-coordinated ZnCl is obtained through 6-coordinated\n2. Jain, A., et al. Commentary: The Materials Project: A materials genome approach to\naccelerating materials innovation. APL Mater. 1, 011002 (2013).\n10. How the authors created the interaction diagram is not fully clear in current format.\nResponse: We sincerely thank the reviewer for highlighting this point. We agree that the\ndescription of how the interaction diagram was created could be clearer in the original\nmanuscript. The interaction diagrams in Supplementary Figs. 2−5 include the following\nanalytical steps:\n(1) Symmetry analysis: Molecular orbitals for the tetrahedral and octahedral structures\nwere calculated using Gaussian 16, and the results can provide the fundamental symmetry\ninformation (e.g., irreducible representations: a , a , t , e …) of the orbitals.\n1 1g 1g g\n(2) Orbital electron density visualization: The relevant molecular orbitals were then\nvisualized using Multiwfn software to compute and export electron density isosurface data from\nthe Gaussian output files. These data were rendered into 3D representations of the orbital lobes\nusing VMD software.\n(3) Bonding Interaction Analysis: The bonding interactions depicted in the diagram were\ninferred through detailed analysis that combined the orbital symmetry information obtained\nfrom Gaussian (Step 1) with the spatial distribution, phase, and overlap characteristics of the\norbital lobes visualized via Multiwfn/VMD (Step 2). The bonding patterns (e.g., antibonding\nand nobonding interactions) shown in the interaction diagrams were identified and illustrated\nthrough analysis of the symmetry matching, extent of orbital overlap, and relative energies of\npertinent frontier orbitals\nAs a consequence, the labels and dashed lines indicating bonding interactions in the\ndiagram represent the conclusions based on this combined symmetry and orbital shape analysis.\nWe have revised the methods in the revised manuscript to provide a clearer description of this\nprocess.\nRevisions in methods of the manuscript:\n“ Orbital analysis and interaction diagram. The orbital analysis of tetrahedral and octahedral\nstructures was performed with the Gaussian16 package44. The Lee–Yang–Parr (B3LYP)45\ncorrelation functional at def2-TZVP46 level was used for the optimization and self-consistent\nfield (SCF) calculations, with DFT-D3 (Becke-Johnson)47,48 method for the van der Waals\ninteractions. The visualization of the orbital results was conducted using Multiwfn49 and\nVMD50 software. The irreducible representations such as a , a , t , e , were also obtained from\n1 1g 1g g\nthe Gaussian 16 calculation results. The bonding interactions illustrated in the interaction\ndiagrams were deduced based on the orbital symmetry information and the spatial\ndistribution/overlap characteristics of the visualized orbitals.”\n11. For many figures, color scheme for each atom is currently missing.\nResponse: We sincerely thank the reviewer for pointing out the missing color schemes for\natoms. We fully agree that clear atomic color coding is essential for interpretation. In the revised\nversion, we have carefully added the missing color schemes for atoms to all affected figures.\nThese updated figures now explicitly define the color representing each atom type, significantly\nenhancing their clarity. The specific modifications are as follows:\n(1) In Fig. 1, we have added color schemes and supplemented essential caption annotations;\nRevisions in the manuscript:\nFig. 1| Comparison of properties of Li/Zn-based halide materials and composition\nengineering design of zinc-ion solid electrolytes. a, b Extranuclear electron orbital\nconfigurations of (a) Li+ and (b) Zn2+. Solid blue spheres represent occupied electrons, while\nhollow black spheres represent unoccupied electrons. c, d Metal-halogen tetrahedron (c) MX\n4\nand octahedron (d) MX (M = Li or Zn; X = Halogen atom). e, f Energy profiles of Li or Zn ion\n6\nmigration pathways in (e) β-Li ZrCl and (f) Cs ZnCl . The insets are the crystal structures of\n2 6 2 4\n(e) β-Li ZrCl and (f) Cs ZnCl . g Schematic of composition engineering design of zinc-ion\n2 6 2 4\nsolid electrolytes. In the schematic, A represents an inorganic atom, and R represents organic\ngroups. h Structure diagrams of anion unit [ZnX ]2− (X = Halogen atom) and some possible\n4\norganic cationic clusters.\n(2) We have added missing atomic annotations in Figs. 2a and 2c;\nRevisions in the manuscript:\nFig. 2 | a, c Primitive cells of (a) PipZnBr and (c) PipZnCl . b, d Chemical phase diagrams of\n4 4\n(b) PipZnBr and (d) PipZnCl .\n4 4\n(3) We have added the color scheme in Supplementary Figs. 2−5 and Supplementary Fig.\n8.\nRevisions in the supplementary information:\nSupplementary Fig. 2 | Qualitative interaction diagram based on orbital analysis of\n[LiX ]3− (X = Halogen atom) tetrahedral cluster. a Li atom; b [LiX ]3− tetrahedral cluster; c\n4 4\n[X ]0 tetrahedral cluster. The red and blue dots in a, b, and c signify spin-up and spin-down\n4\nelectrons, respectively. The dashed green boxes inserted in b show the electron clouds of the\ncorresponding orbital of the [LiX ]3− tetrahedral cluster. Within the green boxes, bluish-gray\n4\natoms represent halogen atoms (X), green atoms signify Li atoms, and red/blue colors are orbital\nelectron density.\nSupplementary Fig. 3 | Qualitative interaction diagram based on orbital analysis of\n[LiX ]5− (X = Halogen atom) octahedral cluster. a Li atom; b [LiX ]5− octahedral cluster; c\n6 6\n[X ]0 octahedral cluster. The conventions for depicting spin electrons, atoms, and orbital\n6\nelectron densities in this figure are identical to those in Supplementary Fig. 2.\nSupplementary Fig. 4 | Qualitative interaction diagram based on orbital analysis of\n[ZnX ]2− (X = Halogen atom) tetrahedral cluster. a Zn atom; b [ZnX ]2− tetrahedral cluster;\n4 4\nc [X ]0 tetrahedral cluster. The red and blue dots in a, b, and c signify spin-up and spin-down\n4\nelectrons, respectively. The dashed green boxes inserted in b show the electron clouds of the\ncorresponding orbital of the [ZnX ]2− tetrahedral cluster. Within the green boxes, bluish-gray\n4\natoms represent halogen atoms (X), silvery gray atoms signify Zn atoms, and red/blue colors\nare orbital electron density.\nSupplementary Fig. 5 | Qualitative interaction diagram based on orbital analysis of\n[ZnX ]4− (X = Halogen atom) octahedral cluster. a Zn atom; b [ZnX ]4− octahedral cluster; c\n6 6\n[X ]0 octahedral cluster. The conventions for depicting spin electrons, atoms, and orbital\n6\nelectron densities in this figure are identical to those in Supplementary Fig. 4.\nSupplementary Fig. 10 | Structure of possible organic molecules and their corresponding\ncations. From left to right are piperazine, 3-aminopyrrolidine, homopiperazine, N-\nmethylpiperazine, and 3-aminopyridine. The Brown, blue, and white spheres depict carbon (C),\nnitrogen (N), and hydrogen atoms (H), respectively.\n12. On line 171, “The valence band maximum (VBM) mainly consists of the nonbonding states\nof Br 4p orbitals” This statement can only be drawn with further details discussed later in the\nparagraph.\nResponse: We appreciate the reviewer's thoughtful observation regarding the description of\nnonbonding states at the VBM in PipZnBr₄. We agree that definitive orbital characterization\nrequires integrated evidence. This insightful comment can help us clarify the evidential\nhierarchy in describing orbital characteristics. In the revised version, DOS analysis serves as\nthe primary basis for identifying nonbonding/antibonding states (Fig. 2f), while the charge\ndensity plots (Fig. 2g, h) provide complementary spatial visualization rather than standalone\nproof. To better reflect this, we have reordered the text to foreground the DOS evidence.\nRevisions in the manuscript:\n“The density of states (DOS) analysis (Fig. 2f) reveals that the valence band maximum (VBM)\nis dominated by nonbonding states of Br 4p orbitals, whereas the adjacent valence band (VB)\nregion consists of antibonding states from Zn 3d and Br 4p hybridization. The conduction band\nminimum (CBM) primarily consists of the antibonding states of Zn 4s and Br 4p orbitals, as\nwell as partial p orbitals of organic cations (Fig. 2f). Additionally, the local charge density can\nprovide a more intuitive representation of the orbital contributions to support this interpretation.\nAs shown in Fig. 2g, …”\n13. “For PipZnCl , due to the electronegativity of the Cl atom is more negative than that of the\ncase results in a larger band gap (6.31 eV) for PipZnCl than that of PipZnBr (5.37 eV)” This\n4 4\nis a correct statement but might not be intuitive.\nResponse: We sincerely thank the reviewer for their suggestion regarding the intuitiveness of\nour explanation. To enhance clarity, we have revised the explanation regarding the larger band\ngap in PipZnCl₄ (6.31 eV) compared to PipZnBr₄ (5.37 eV) by emphasizing the direct impact\nof electronegativity differences.\nThe revised manuscript contains the following modifications to the new sentences.\nRevisions in the manuscript:\n“For PipZnCl , the higher electronegativity of Cl relative to Br lowers the energy of Cl 3p\n4\norbitals, which results in stronger Zn 4s–Cl 3p orbital hybridization. This effect enhances the\nsplitting between bonding and anti-bonding states, significantly elevating the energy of the anti-\nbonding orbital at the CBM. Consequently, PipZnCl₄ exhibits a larger band gap (6.31 eV)\ncompared to PipZnBr₄ (5.37 eV), as depicted in Fig. 2i.”\n14. Likely a typo in figure 2j, it should be Cl instead Br\nResponse: We sincerely thank the reviewer for their meticulous attention to detail. The label\n\"Br\" in Fig. 2j is indeed a typo and should be \"Cl\". We apologize for this oversight. This error\nhas been corrected in the revised manuscript. We have also taken this opportunity to carefully\nre-examine all chemical structures and labels throughout the manuscript to ensure no similar\nerrors remain.\nRevisions in the manuscript:\nFig. 2 | e−l, Calculated (e, i) band structure, (f, j) total and projected DOSs, local charge density\nat (g, k) CBM and (h, l) VBM of PipZnBr and PipZnCl , respectively. The zero in band\n4 4\nstructures and DOSs is denoted as the VBM. The insets in f and j represent projected DOSs\nwhich have been enlarged by 10 times for clearer indication of their orbital contributions. The\nisosurfaces of local charge density are all set to 0.001 e Bohr−3\n.\n15. “The Fourier transform spectra of the extended XAFS (FT-EXAFS) curves in R-space\nshow that the bond length of Zn−Br in PipZnBr4 is shorter than that in pure ZnBr2 (Fig. 6b),\nwhich suggests a relatively weaker Zn−Br bond in PipZnBr4.” Like a typo since Zn-Br is longer\nin Fig. 6b.\nResponse: We greatly thank the reviewer for the careful examination of the associated\ninterpretation for Fig. 6b. We apologize for this mistake, and we have corrected the related\nsentence in our revised manuscript. We have also re-examined all relevant discussions in the\nmanuscript to preclude the possibility of similar errors.\nRevisions in the manuscript:\n“The Fourier transform spectra of the extended XAFS (FT-EXAFS) curves in R-space show\nthat the bond length of Zn−Br in PipZnBr is longer than that in pure ZnBr (Fig. 6b), which\n4 2\nsuggests a relatively weaker Zn−Br bond in PipZnBr .”\n4\n16. In the Discussion, there is a typo “low activation energy (0.01 eV), ...” since the value is\ndifferent from ones in the Results section.\nResponse: We deeply appreciate the reviewer for pointing out this typo. In light of several\ncomments regarding the activation energy received from reviewers, we have recalculated and\ncorrected the previously calculated activation energy. The revised value is now established as\n0.13 eV, and this modification has been appropriately reflected in the revised version.\nRevisions in the manuscript:\n“…, low activation energy (0.13 eV), wide electrochemical window (~2.28 V vs. Zn2+/Zn),\nand high Zn2+ transfer number (0.61).”"}, {"round": 2, "reviewer_comments": {"reviewer_1": "I appreciate the authors provided further results to support the zinc ion conduction in the propsed metal halide hybrids. For example, the Zn-I2 solid state battery using chloride-type electrolyte has also been demonstrated under the small areal capacity of 0.1 mAh cm-2. However, I am still a little confusing with the Zn2+ ion conduction mechanism presented in this work. First, it is not clear how the authors set up conduction sites to calculate the migration pathway and energy barrier as shown in Figure 3i-l. In addition, the calculated migration energy barrier (0.37 eV) is much higher than the activation energy of 0.13 eV as shown in Arrhenius plots. This deviation is too large to be convincing. Finally, to exclude the possibility of halide ion conduction, the authors should use CsSnCl3 (10.1002/aesr.202400198) as the blocking layer not the PP separator. I think it is better to clarify the Zn2+ ion conduction mechanism rather than showing the battery cycling under limited areal capacity.", "reviewer_2": "I would like to take the time to thank the authors' efforts and they addressed all my concerns. The manuscript in its current form can now be considered for publication in nature communications."}, "author_response": "Response to Reviewers’ Comments\nDear Reviewers：\nThank you once again for your continued evaluation of our manuscript entitled\n“Zinc-based metal halide electrolytes for all-solid-state zinc-ion batteries” (NCOMMS-\n25-16926-B). We are grateful for your insightful comments and suggestions, which\nhave been instrumental in further improving our work. In response to the remaining\npoints raised during this second round of review, we have provided additional\ntheoretical and experimental data, as well as the necessary textual amendments to solve\nthe concerns point-by-point as follows. All changes have been highlighted in yellow in\nthe main manuscript and the supplementary materials. New or modified figures are\nindicated with a red box for easy identification.\nResponse to reviewers’ comments\nReviewer #1\nI appreciate the authors provided further results to support the zinc ion conduction in\nthe proposed metal halide hybrids. For example, the Zn-I solid state battery using\n2\nchloride-type electrolyte has also been demonstrated under the small areal capacity of\n0.1 mAh cm−2. However, I am still a little confusing with the Zn2+ ion conduction\nmechanism presented in this work.\nResponse: We sincerely thank Reviewer #1 for the thoughtful review and for\nacknowledging the additional results we provided to demonstrate zinc-ion conduction\nin the proposed metal halide hybrids. We are also grateful to Reviewer #1 for the\nconstructive feedback regarding the Zn²⁺ ion conduction mechanism. We are trying to\nsolve these concerns as follows.\nFirst, it is not clear how the authors set up conduction sites to calculate the migration\npathway and energy barrier as shown in Figure 3i-l.\nResponse: We greatly thank the reviewer for pointing out this concern. To alleviate the\nconcern of the reviewer, we provide a more detailed explanation of the settings for the\nmigration sites and paths in Figure 3i−l.\nFirstly, we performed a series of density functional theory (DFT) calculations to\nidentify the most stable interstitial sites for an additional Zn2+ ion within the PipZnBr\n4\nand PipZnCl structures. The lowest-energy site was selected as the initial state (starting\n4\npoint) for the migration simulation. Secondly, a symmetry-equivalent site along the a-\naxis was chosen as the final state. Given the symmetry of the sites, the energy levels of\nboth the initial and final states remain essentially equivalent. Finally, the intermediate\nimages along the migration pathway were initially generated using the built-in utility\nscript of the climbing-image nudged elastic band (Cl-NEB) method. These images were\nthen carefully inspected and manually adjusted where necessary to ensure a smooth and\nchemically intuitive path before the barrier calculation was performed.\nTo ensure a comprehensive understanding of the Zn²⁺ migration and to validate the\nresults, we applied an analogous methodology to identify stable sites and calculate\nmigration pathways along the b- and c-axis directions (Supplementary Figs. 15 and\n16 ). We believe this systematic approach robustly defines the conduction sites and\nprovides reliable energy barriers for Zn²⁺ migration in the studied crystals.\nWe have revised the manuscript accordingly.\nRevisions in the methods of the manuscript:\n“Zn-ion migration energy barrier calculations. The initial sites of the Zn2+ migration\ncalculations were determined through numerous DFT energy calculations.\nSubsequently, the final position was identified as the symmetrical position in the\ncorresponding direction of the supercell. For the migration energy barrier calculations\nalong the a, b, and c directions, supercells of dimensions 2×1×1, 1×2×1, and 1×1×2\nwere employed, respectively. The diffusion pathways and barrier energies…”\nIn addition, the calculated migration energy barrier (0.37 eV) is much higher than the\nactivation energy of 0.13 eV as shown in Arrhenius plots. This deviation is too large to\nbe convincing.\nResponse: We sincerely thank the reviewer for this valuable comment. We agree that\nthere is an apparent difference between the calculated migration energy barrier and the\nexperimental activation energy, and we appreciate the opportunity to elaborate on the\npotential reasons behind it. This difference may arise from fundamental distinctions\nbetween experimental and theoretical methods, as well as material-specific factors.\nFirstly, it is important to note that the theoretical migration barrier was obtained from\nfirst-principles calculations at 0 K within an idealized periodic crystal model. This\napproach is designed to isolate the fundamental migration mechanism and its intrinsic\nenergy barrier. However, in real materials, especially at elevated temperatures where\ndiffusion occurs, the presence of multiple vacancies, vacancy clusters, or other\nextended defects is common.\nSecondly, the experimental activation energy is measured at finite temperatures,\nwhile the static calculation does not incorporate vibrational free energy or anharmonic\neffects. So the temperature-dependent entropy and other factors are one of the reasons\nfor the inconsistency between theoretical and experimental values.\nEntropic contributions and temperature-dependent lattice dynamics often lead to a\nlower apparent activation energy in experiments than the theoretically predicted barrier\nheight, which represents solely the enthalpic component under frozen-ion conditions.\nIn addition, this phenomenon aligns with reports in solid-state lithium-ion electrolyte\nliterature. For example, Li Sc Zr Cl exhibits a calculated migration barrier of\n2.75 0.75 0.25 6\n0.47 eV but an experimental activation energy of 0.33 eV (ACS Energy Letters, 2024,\n9(3): 1043-1052), further supporting the discussed origins of such discrepancies.\nConsequently, while the calculated barrier serves as a qualitative indicator of a\nmaterial’s ion migration capability, it should not be directly compared to experimental\nactivation energies in absolute values.\nFinally, to exclude the possibility of halide ion conduction, the authors should use\nCsSnCl (10.1002/aesr.202400198) as the blocking layer not the PP separator. I think it\n3\nis better to clarify the Zn2+ ion conduction mechanism rather than showing the battery\ncycling under limited areal capacity.\nResponse: We sincerely appreciate the reviewer’s insightful suggestion. Following the\nsuggestions, we made our best efforts to synthesize CsSnBr and CsSnCl samples\n3 3\naccording to Reference 10.1002/aesr.202400198, with the corresponding X-ray\ndiffraction pattern shown in Supplementary Fig. 48a, b. The Reference has also been\nadded to the reference list. The synthesized PipZnBr and PipZnCl using CsSnBr and\n4 4 3\nCsSnCl samples as the blocking layer were pressed into the 16 mm diameter pellets\n3\nwith the thickness of ~487 and ~498 μm as the solid-state zinc-ion electrolyte\n(Supplementary Fig. 48c, d). These blocking layers can inhibit Zn2+ ion migration\nwhile allowing the transit of halide ions in PipZnBr and PipZnCl electrolytes. To\n4 4\ninvestigate the migration of halide ions, we evaluated the resistance of\nSS||PipZnBr ||CsSnBr ||SS and SS||PipZnCl ||CsSnCl ||SS symmetrical cells at room\n4 3 4 3\ntemperature. Compared with the Nyquist plot of SS||PipZnBr ||SS (Supplementary Fig.\n4\n22f), the resistance profiles of SS||PipZnBr ||CsSnBr ||SS exhibit increased electrolyte\n4 3\nresistance and charge transfer resistance, indicating greater difficulty in Br− ion\nmigration (Fig. R1). The SS||PipZnCl ||CsSnCl ||SS cell displays a similar phenomenon:\n4 3\nalthough partial halogen ion migration occurs, the migration process is considerably\nsluggish.\nTo verify the contribution of Zn2+ and halide ion (i.e.,Br− or Cl−) migration in the\nbattery, we performed Zn||Zn symmetrical cells and Zn||I full cells tests. Firstly, the\n2\ncyclic voltammetry (CV) curves of Zn||PipZnBr ||CsSnBr ||I and\n4 3 2\nZn||PipZnCl ||CsSnCl ||I cells were tested. After blocking Zn2+ conduction, the CV\n4 3 2\ncurves are difficult to determine the oxidation-reduction peak position (Supplementary\nFig. 49a, b), and the repeatability is also very poor. Furthermore, we also assessed the\ncycling performance of Zn||PipZnBr ||CsSnBr ||Zn and Zn||PipZnCl ||CsSnCl ||Zn cells.\n4 3 4 3\nAs shown in Supplementary Fig. 49c, after blocking Zn2+ conduction, the\nZn||PipZnBr ||CsSnBr ||Zn cell can not work at 0.1 mA cm−2 and 0.1 mAh cm−2 at 40 ℃.\n4 3\nSimilarly, the Zn||PipZnCl ||Zn cell by applying a CsSnCl interlayer is also difficult to\n4 3\nrun with the normal charge/discharge curve (Supplementary Fig. 49d). Not\nsurprisingly, the Zn||PipZnBr ||CsSnBr ||Zn and Zn||PipZnCl ||CsSnCl ||Zn full cells are\n4 3 4 3\nalmost unable to operate at 0.2 mA cm−2 at 40 ℃ (Fig. R2). Based on the above\nelectrochemical test results, although halide ions could migrate in PipZnBr and\n4\na dominant role. We thank the reviewer again for your insightful comments. Next,\ninspired by the fact that halide ions (i.e.,Br− or Cl−) could also be transported, we are\nmotivated to explore the development of halogen batteries for enhanced utilization.\nSupplementary Fig. 22 | f Nyquist plot of SS||PipZnBr ||SS cell at 25 ℃.\nCsSnBr and CsSnCl samples as the blocking layer on the electrolyte side. a\n3 3\nNyquist plot of SS||PipZnBr ||CsSnBr ||SS cell at 25 ℃. b Nyquist plot of\n4 3\nSS||PipZnCl ||CsSnCl ||SS cell 25 ℃.\n4 3\nFig. R2 | Electrochemical characterization of Zn||I cell with halide-based\n2\nelectrolyte using CsSnBr and CsSnCl samples as the blocking layer. a Cycling\n3 3\nperformance of Zn||PipZnBr ||CsSnBr ||I cell at 0.2 mA cm−2. b Cycling performance\n4 3 2\nof Zn||PipZnCl ||CsSnCl ||I cell at 0.2 mA cm−2.\n4 3 2\nRevisions in the manuscript:\n“…Finally, to exclude the possibility of halide ion conduction, CsSnBr and CsSnCl\n3 3\nsamples as the blocking layer were applied in the solid-state PipZnBr and PipZnCl\n4 4\nelectrolytes (Supplementary Fig. 48). It is found that the CV curces of\nZn||PipZnCl ||CsSnCl ||I and Zn||PipZnCl ||CsSnCl ||I full cell have no oxidation-\n4 3 2 4 3 2\nreduction peak and Zn||Zn symmetrical cells are almost unable to operate\n(Supplementary Fig. 49). These results demonstrate that although halide ions could\nmigrate in PipZnBr and PipZnCl electrolytes, Zn2+ ion conduction in the halide-based\n4 4\nzinc-ion batteries play a dominant contribution.”\nRevisions in the supplementary information:\nSupplementary Fig. 48 | a XRD spectra for CsSnBr powder. b XRD spectra for\n3 4\nelectrolyte using CsSnBr sample as the blocking layer. c Optical photograph of the\n3\nthickness (~498 μm) of PipZnCl electrolyte using CsSnCl sample as the blocking\n4 3\nlayer.\nSupplementary Fig. 49 | Electrochemical performance of zinc-ion cell with halide-\nbased electrolyte using CsSnBr and CsSnCl samples as the blocking layer. a CV\n3 3\ncurves of the initial three cycles of Zn||PipZnBr ||CsSnBr ||I cell at 0.1 mV s−1. b CV\n4 3 2\ncurves of the initial three cycles of Zn||PipZnCl ||CsSnCl ||I cell at 0.1 mV s−1. c\n4 3 2\nCycling performance of Zn||PipZnBr ||CsSnBr ||Zn cell at 0.1 mA cm−2 and 0.1 mAh\n4 3\ncm−2. d Cycling performance of Zn|| PipZnCl ||CsSnCl ||Zn cell at 0.1 mA cm−2 and 0.1\n4 3\nmAh cm−2.\nReviewer #2:\nI would like to take the time to thank the authors' efforts and they addressed all my\nconcerns. The manuscript in its current form can now be considered for publication in\nnature communications.\nResponse: We are deeply grateful to Reviewer #2 for their positive feedback and for\nacknowledging our revisions. We would also like to reiterate our appreciation for their\ninsightful comments during the previous round, which were invaluable in improving\nthe quality of our work. We greatly appreciate the time and expertise they have\ncontributed to the review process."}, {"round": 3, "reviewer_comments": {"reviewer_1": "The authors have provided clear evidence to show the Zn2+ conduction in their synthesized solid electrolytes. I am happy to see the publication of this work."}, "author_response": ""}]

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