# Wadsley-Roth FeNb11O29 as negative electrode material for lithium solid-state batteries

Dat Le Thanh, Dominique Foix, Emmanuelle Suard, Romain Berthelot, Long

Hoang Bao Nguyen

► To cite this version:

Dat Le Thanh, Dominique Foix, Emmanuelle Suard, Romain Berthelot, Long Hoang Bao Nguyen. Wadsley-Roth FeNb11O29 as negative electrode material for lithium solid-state batteries. Journal of Power Sources, In press, 666, pp.239120. 10.1016/j.jpowsour.2025.239120. hal-05426474

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Submitted on 19 Dec 2025

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# Wadsley-Roth FeNb₁₁O₂₉ as negative electrode material for lithium solid-state batteries

Dat Le Thanh<sup>a,b,c</sup><sup>ID</sup>, Dominique Foix<sup>c,d,e</sup>, Emmanuelle Suard<sup>a</sup><sup>ID</sup>, Romain Berthelot<sup>b,e</sup><sup>ID</sup>, Long Hoang Bao Nguyen<sup>b,e,*</sup><sup>ID</sup>

## HIGHLIGHTS

* Wadsley-Roth FeNb₁₁O₂₉ is implemented as negative electrode material in solid-state batteries.

* Intermediate VGCF content balances ionic transport and electronic conductivity.

* Kinetic limitation of Li⁺ insertion is mitigated by constant-voltage discharge steps.

* XPS reveals resistive Li₂S/oxide-rich interphase consistent with impedance growth.

## GRAPHICAL ABSTRACT

<figure><img src="image_4.png" /></figure>

## ARTICLE INFO

### Keywords:

FeNb₁₁O₂₉

Wadsley-Roth phase

Solid-state batteries

Electron-ion transport

Interfacial stability

## ABSTRACT

Li-ion solid-state batteries are promising next-generation rechargeable batteries thanks to their high energy density and enhanced safety, but their performance is strongly dependent on the choice of active materials, electrode microstructure and interfacial stability. In this work, the low-strain Wadsley-Roth phase FeNb₁₁O₂₉ is explored as a negative electrode for Li-ion solid-state batteries, benefiting from its open framework for fast Li⁺ transport. Composite electrodes are optimised by tuning the ratio of FeNb₁₁O₂₉, Li₆PS₅Cl solid electrolyte, and vapor grown carbon fibres. The optimal composition can deliver a high-rate capability, achieving ~210 and ~80 mAh·g⁻¹ at 0.19 and 1.9 mA·cm⁻², respectively, even at a high areal loading of ~10 mg·cm⁻², arising from a balanced electronic and ionic transport. Kinetic limitation of Li⁺ insertion at high rates is identified and can be partly mitigated by end-of-discharge constant-voltage steps. Post-mortem analyses confirm bulk stability of

<footnote><sup>a</sup> Institut Laue-Langevin (ILL), 38042, Grenoble, France</footnote>

<footnote><sup>b</sup> ICGM, Univ Montpellier, CNRS, ENSCM, Montpellier, France</footnote>

<footnote><sup>c</sup> Alistore-European Research Institute, CNRS FR 3104, 80039, Amiens, France</footnote>

<footnote><sup>d</sup> IPREM/ECP, Université de Pau, CNRS, 64053 Pau, France</footnote>

<footnote><sup>e</sup> Réseau sur le Stockage Électrochimique de l'Énergie (RS2E), CNRS, 80039, Amiens, France</footnote>

<footnote>* Corresponding author. ICGM, Univ Montpellier, CNRS, ENSCM, Montpellier, France.</footnote>

<footnote>E-mail address: long-hoang-bao.nguyen@umontpellier.fr (L.H.B. Nguyen).</footnote>
FeNb<sub>11</sub>O<sub>29</sub> and Li<sub>6</sub>PS<sub>5</sub>Cl but revealed an interfacial degradation, with a resistive Li<sub>2</sub>S/oxide-rich interface, consistent with impedance growth. These findings highlight the critical role of microstructural and interfacial optimisation in the electrode composites and establish FeNb<sub>11</sub>O<sub>29</sub> as a promising candidate for high-rate solid-state batteries.

# 1. Introduction

Li-ion solid-state batteries (Li-SSBs) have attracted significant attention in recent years due to their enhanced safety, durability, and high energy density [1–6]. Replacing flammable liquid electrolytes with solid-state electrolytes (SSEs) not only improves safety but also enables the use of high-capacity electrode materials, and broadens the operational temperature and current density range [2,7–10]. Various negative electrode materials have been studied for Li-SSBs, each with distinct advantages but also major limitations. Lithium metal stands out for its low potential and high specific capacity, yet its practical use is hindered by severe interfacial reactivity with SSEs, poor physical contact, and risks of dendrite growth due to inhomogeneous Li plating/stripping [11–14]. Graphite, in contrast, provides structural stability, and high electronic conductivity, but its relatively low capacity, sluggish kinetics, and risk of lithium plating at low potentials restrict its performance in high-rate solid-state systems [15–18]. Silicon electrode has recently been successfully integrated in Li-SSBs, offering high capacity, though electrode formulation process has to be carefully optimised to minimise volume changes and stress-induced contact loss that occur during cycling [19–23].

In the search for high-rate and stable negative electrode materials for Li-SSBs, several high-voltage, zero-strain materials have emerged as promising candidates, including spinel Li<sub>4</sub>Ti<sub>5</sub>O<sub>12</sub> and block phases of the Wadsley-Roth niobium-based oxide family [24–32]. Their relatively high operating potential (~1.5–1.6 V vs. Li<sup>+</sup>/Li) may enhance chemical compatibility with different solid electrolytes, mitigate interfacial side reactions, and reduce the risk of lithium plating, though at the expense of energy density. In addition, their low volume change during lithiation/delithiation may help limit mechanical stress, suppress particle cracking, and preserve interfacial contact, which are key factors for long-term cycling stability [24–31]. While Li<sub>4</sub>Ti<sub>5</sub>O<sub>12</sub> has been extensively developed using diverse improvement approaches, studies on Wadsley-Roth phases in Li-SSBs remain limited to a few compositions. For example, the performance of TiNb<sub>2</sub>O<sub>7</sub>-based electrodes has been improved by (i) optimising the electronic transport via the incorporation of carbon black and carbon nanotubes as complementary short- and long-range electron pathways [27], and (ii) employing Mo<sup>6+</sup> doping to address the low conductivity, both leading to enhanced performance [30]. Recently, interfacial stabilisation of Nb<sub>1.6</sub>Ti<sub>0.32</sub>W<sub>0.08</sub>O<sub>5–δ</sub> via *in situ* LiWS<sub>2</sub> formation enabled durable cycling and high rate cycling, delivering 175 mAh g<sup>-1</sup> at 0.2 A g<sup>-1</sup> (1C) and 25 °C for 1.8 mg cm<sup>-2</sup> loading [28].

These recent advances highlight the prospects of high-voltage negative electrode materials in Li-SSBs and motivates the investigation of other Wadsley-Roth block phases, many of which remained unexplored. Expanding the research within this family could diversify available materials and reveal new candidates for solid-state applications. Among them, FeNb<sub>11</sub>O<sub>29</sub> is a promising candidate with favourable properties: a relatively high operating potential, robust structure, fast ion transport in liquid electrolyte systems [33–42], and negligible volume change during cycling [36,39,41]. Compared with the parent compound Nb<sub>12</sub>O<sub>29</sub>, partial substitution of niobium by iron not only introduces a low-cost, earth-abundant resources but also provides an additional redox centre, raising the theoretical capacity of the material to ~400 mAh g<sup>-1</sup>. These properties make FeNb<sub>11</sub>O<sub>29</sub> an attractive alternative to conventional low-potential electrodes, particularly for high-power Li-SSB applications where interfacial and mechanical stability are essential. In this work, the feasibility of employing FeNb<sub>11</sub>O<sub>29</sub>

in Li-SSBs is evaluated for the first time to explore both its electrochemical performance and interfacial/transport limitations. This study first focuses on optimising FeNb<sub>11</sub>O<sub>29</sub>-based composite electrodes by tailoring the ratios of active material, solid-state electrolyte, and conductive additive to build a well-connecting ionic, electronic transport network. Then complementary studies of morphology, electronic conductivity, and lithium diffusion are conducted to gain insight into transport behaviour of the composite, while post-mortem analyses provide information on structural, morphological, interfacial chemistry and degradation mechanisms after cycling.

# 2. Experimental

## 2.1. Preparation of electrode composites

Monoclinic FeNb<sub>11</sub>O<sub>29</sub> was synthesised by solid-state reaction at 1100 °C in 4 h from stoichiometric mixture of Fe<sub>2</sub>O<sub>3</sub> (≥99 %, <5 µm, Honeywell) and Nb<sub>2</sub>O<sub>5</sub> (obtained by heating Nb powder (99.8 %, ~325 mesh, Alfa Aesar) at 900 °C for 24 h). Its phase purity was verified by X-ray diffraction, while its particle size and morphology were examined by SEM (Figs. S1 and S2). Electrode composites were prepared by manually mixing FeNb<sub>11</sub>O<sub>29</sub>, Li<sub>6</sub>PS<sub>5</sub>Cl argyrodite solid-state electrolyte (NEI Corporation) that has an ionic conductivity of ~2.7 mS cm<sup>-1</sup> as measured in this work, and vapor-grown carbon fibres (VGCF, Showa Denko VGCF-H) using a mortar and pestle inside an Ar-filled glovebox (O<sub>2</sub>, H<sub>2</sub>O < 0.1 ppm). Prior to mixing, FeNb<sub>11</sub>O<sub>29</sub> and VGCF were dried overnight at 200°C in a Büchi oven under dynamic vacuum to remove residual moisture. To study the influence of electrode compositions on the ionic/electronic conductivity and cycling performance, the weight ratio of FeNb<sub>11</sub>O<sub>29</sub> to Li<sub>6</sub>PS<sub>5</sub>Cl was fixed at 40:50, while the content of VGCF was systematically varied to have overall (FeNb<sub>11</sub>O<sub>29</sub> + Li<sub>6</sub>PS<sub>5</sub>Cl): VGCF weight ratios of 90:3, 90:5, 90:10, 90:20, and 90:30. These composites are denoted as VGCF-03, VGCF-05, VGCF-10, VGCF-20, VGCF-30, respectively.

## 2.2. Half-cell configuration setup

All electrochemical measurements were conducted inside a PEEK-lined, airtight press cell equipped with stainless steel plungers (Φ = 7 mm) as current collectors. The half-cell assembly process is schematically illustrated in Fig. 1.

35 mg of Li<sub>6</sub>PS<sub>5</sub>Cl was first evenly spread in the cell and uniaxially pressed between two plungers under 1 T (260 MPa) for 3 min to form a uniform separator layer. Then ~10 mg of FeNb<sub>11</sub>O<sub>29</sub> composite (active material loading ~10.9 mg cm<sup>-2</sup>) was uniformly placed on one side of the separator. This amount is required to fully cover the area of cell interior (Φ = 7 mm). To allow fair comparison between different composite compositions, the active material loading was kept constant, while the total electrode mass was adjusted accordingly. All steps were performed carefully to ensure the homogeneous distribution and full coverage of the powders. The separator and the composite layer were densified with a uniaxial press at 0.7 T (182 MPa) for 3 min. Finally, the counter electrode was prepared by stacking two indium discs (Φ = 6 mm, thickness = 0.127 mm, m ≈ 31.0 mg, Sigma-Aldrich 99.99 %) and one lithium disc (Φ = 5 mm, m ≈ 2.5 mg) in In/Li/In configuration, which was then gently pressed onto the cell plunger. This plunger was placed on the opposite side of the Li<sub>6</sub>PS<sub>5</sub>Cl separator, and the cell was subjected to final pressing step at 0.6 T (156 MPa) for 3 min. Before all tests, a resting duration of 24 h is necessary for Li-In alloy formation
and pressure equilibration inside the cell. To ensure comparison, the practical parameters, such as applied pressure and material quantities, were maintained consistent for all experiments.

## 2.3. Benchmarking the electrode composites

### 2.3.1. Rate capability tests

All galvanostatic cycling tests with potential limitations (GCPL) were conducted in half-cells within a voltage window of 0.4–2.4 V vs. Li⁺/Li⁻In, equivalent to 1.0–3.0 V vs. Li⁺/Li. The nominal capacity (1C) is defined as the exchange of 12 Li⁺ ions per hour per FeNb₁₁O₂₉ formula unit, corresponding to a specific capacity of 209 mAh·g⁻¹. Concerning rate capability tests, current rates of 1, 2, 5, 10, 20, and 50 Li⁺/h per formula unit were applied, followed by a return to 1 Li⁺/h current rate, with each rate held for 5 cycles. These conditions correspond to C-rates from C/12 to ~4.2C, with equivalent gravimetric and areal current densities summarised in Table 1.

For comparison, liquid electrolyte-based coin cells were assembled and tested under similar conditions. The cell used a conventional configuration of FeNb₁₁O₂₉: PVdF: Super P = 80:10:10 (wt%) coated on Cu current collector, paired with Li metal as the counter/reference electrode and Whatman glass fibre separators soaked in 1M LiPF₆ in EC: DMC (3/7 by volume) as electrolyte. The active material loading is around 10 mg cm⁻², similar to that of solid-state cells. Rate capability measurements were then carried out following the same protocol.

### 2.3.2. Electronic and ionic conductivity measurements

Two symmetric cell configurations were used to determine the electronic and ionic conductivity of the electrode composite (Fig. S3), following reported methods [43,44], with some modifications were made to adapt to the cell dimensions. For ion-blocking setup (stainless steel | composite | stainless steel), 25 mg of composite was placed in the cell and uniaxially pressed under 1 T for 3 min. The two stainless steel plungers play the role as the ion-blocking electrodes, allowing only electronic transport to be measured. Electron-blocking setup was used to measure ionic conductivity, in which a sandwich stack of Li⁻In | Li₆PS₅Cl | composite | Li₆PS₅Cl | Li⁻In was assembled: 25 mg of composite was pressed, followed by the addition and densification of 30 mg Li₆PS₅Cl on each side, one at a time. Finally, indium and lithium discs were placed on both sides and pressed at 0.6 T. Before the conductivity test, a resting duration of 24 h was applied.

Direct current (DC) polarisation with a stepwise voltage protocol was conducted using ion- or electron-blocking setups to determine the effective electronic and ionic conductivity of the composite. In each step, a constant voltage was applied, and the current response was recorded in 1 h until equilibrium. Voltage steps of 5–30 mV or 1–6 mV were selected to keep the current within the instrument’s resolution range. Conductivity was then calculated using the equation $σ = L/(R ×$

<caption>Table 1</caption>

Rate capability test expressed in terms of Li⁺ exchanged per hour per formula unit, C-rate, gravimetric (mA g⁻¹) and areal (mA cm⁻²) current densities.

<table><thead><tr><th>Li<sup>+</sup>/h</th><th>1</th><th>2</th><th>5</th><th>10</th><th>20</th><th>50</th></tr></thead><tbody><tr><th>C-rate</th><td>C/12</td><td>C/6</td><td>0.4C</td><td>0.8C</td><td>1.7C</td><td>4.2C</td></tr><tr><th>mA g<sup>-1</sup></th><td>17.4</td><td>34.8</td><td>87</td><td>174</td><td>348</td><td>870</td></tr><tr><th>mA cm<sup>-2</sup></th><td>0.19</td><td>0.38</td><td>0.95</td><td>1.9</td><td>3.8</td><td>9.5</td></tr></tbody></table>

S), where R is the steady-state resistance obtained from linear fitting of the stabilised current and the voltage step (R = ΔV/I), L and S are the pellet thickness and area.

## 2.4. Electrochemical tests of the optimised composite

### 2.4.1. Cycling performance

Galvanostatic cycling at the rate of C/6 (34.8 mA g⁻¹) was carried out for half-cells of the optimised composite for 50 cycles.

### 2.4.2. Asymmetric rate tests

Asymmetric rate tests were applied for half-cells containing optimised composite, in which either the charge or discharge current was fixed while the other was progressively increased. The purpose is to compare Li⁺ insertion and extraction kinetics in FeNb₁₁O₂₉. In the D-rate test, the charge current was fixed at C/12, while the discharge current was varied stepwise every five cycles from D/12 to 4.2D, similar to the conventional rate protocol. Conversely, in the C-rate test, the discharge current was fixed at D/12 while the charge current was varied in the same way.

A rate capability test was performed similar to the conventional rate protocol, with a constant-voltage step was added at the end of discharge to enhance lithium insertion. This voltage hold was maintained until the current dropped below C/40 (0.022 mA), or for a maximum of 5 h, whichever limit was reached first.

### 2.4.3. Cyclic voltammetry

Cyclic voltammetry was performed on a solid-state half-cell using the optimised composite, and a liquid electrolyte-based coin cell to investigate its redox behaviour and lithium-ion diffusion properties. Measurements were conducted within the same electrochemical window (1.0–3.0 V vs. Li⁺/Li) at increasing scan rates of 0.10, 0.15, 0.20, 0.25, and 0.30 mV s⁻¹. The resulting voltammograms were analysed qualitatively based on the characteristic anodic and cathodic peaks and quantitatively using the Randles–Sévčík equation [45,46] to estimate Li-ion diffusion coefficients.

<figure><img src="image_5.png" /><figcaption>Fig. 1. Schematic illustration of the half-cell assembly processes. Steps ①–③ show the assembly sequence with corresponding uniaxial pressures. The cell stack includes the Li₆PS₅Cl separator, FeNb₁₁O₂₉ composite layer, and In/Li/In counter electrode.</figcaption></figure>
## 2.5. Pre- and post-cycling characterisations

X-ray diffraction (XRD) were performed using a Panalytical X'pert Spinner diffractometer (equipped with Cu Kα₁,₂ X-ray source) in Bragg-Brentano θ-θ geometry, within the angular range of 10–70° and with a

step size of 0.033°. All the air-sensitive samples were protected using a Kapton® polyimide film covered on a silicon zero-background sample holder. The obtained XRD patterns were fitted by Le Bail method using FullProf suite [47], with the refined parameters including the zero shift, lattice parameters, and peak profile parameters. In order to recover the

<figure><img src="image_6.png" /><figcaption>**Fig. 2.** Voltage profiles and differential capacity (dQ/dV) curves of five FeNb₁₁O₂₉ composites and a liquid electrolyte coin cell at C/12 for the first (a, b) and second (c, d) cycle (1C = 2.28 mA cm⁻²), and their rate capability tests (e). Inset (d): polarisation from dQ/dV peak separations.</figcaption></figure>
FeNb<sub>11</sub>O<sub>29</sub> composite after cycling, the plunger on the FeNb<sub>11</sub>O<sub>29</sub> side was opened and FeNb<sub>11</sub>O<sub>29</sub> composite layer was extracted using a spatula.

**Scanning electron microscopy (SEM)** combined with energy-dispersive X-ray spectroscopy (EDS) analyses were performed using an FEI Quanta 200 FEG SEM microscope.

**X-ray photoelectron spectroscopy (XPS)** was carried out using a THERMO Escalab 250Xi spectrometer, using focused monochromatic Al Kα radiation (hν = 1486.6 eV) with an X-ray spot size of 650 µm for longer axis ellipse. The spectrometer was calibrated with reference binding energies (BE) of clean samples: Cu (932.6 eV), Ag (368.2 eV) and Au (84.0 eV). The samples were never in contact with air nor moisture from their preparation to their analysis since the XPS spectrometer is directly connected to an Ar-filled glove box. The high-energy resolution spectra were recorded with constant pass energy of 20 eV. The charge neutralisation was used for all the acquisitions. The pressure in the analysis chamber was around 2 × 10<sup>-7</sup> mbar during analyses. Short acquisition time spectra were recorded before each experiment to check that the samples were not subject to degradation during the X-ray irradiation. For the pristine material FeNb<sub>11</sub>O<sub>29</sub>, the binding energy scale was corrected using as reference the C 1s peak at 285.0 eV characteristic of the hydrocarbon species CH<sub>2</sub>–CH<sub>2</sub>. For the samples containing Li<sub>6</sub>PS<sub>5</sub>Cl, this correction was done with respect to the Cl 2p position at 199.0 eV. For data analysis, the spectra were mathematically fitted with Casa XPS software using a Shirley type background, and a combination of Gaussian and Lorentzian distributions for peaks fitting.

**Electrochemical impedance spectroscopy (EIS)** measurements were performed on a VSP impedance analyser (Biologic), with an excitation voltage of 10 mV in the frequency range of 1 MHz–100 mHz, and 10 points per decade were recorded. DRT analysis was performed with RelaxIS Online DRT (rhd instruments), using the radial-basis-function (RBF) discretisation approach described by Ciucci et al. [48].

# 3. Results and discussions

## 3.1. Optimisation of the electrode composite compositions

Galvanostatic cycling performance of several solid-state cells and liquid electrolyte coin cells is compared in Fig. 2a–d, showing their voltage profiles and differential capacity (dQ/dV) curves at C/12 for the first two cycles. Both systems exhibit similar shapes of charge-discharge curves with characteristic sloping features, consistent with the comparable redox peaks at ~2.4 V, 1.65 V, and 1.3 V corresponding to Fe<sup>3+/2+</sup>, Nb<sup>5+/4+</sup> and Nb<sup>4+/3+</sup> redox couples, respectively [34–42]. These results confirm the successful implementation of FeNb<sub>11</sub>O<sub>29</sub> in solid-state battery systems for the first time.

All composites exhibit a pronounced capacity drop from the first discharge to the first charge, reflected in their initial Coulombic efficiencies (iCEs) below 90 %, whereas 96 % is achieved with liquid electrolyte. By the second cycle, the Coulombic efficiencies (CEs) of the solid-state cells improve to the range of 89–93 %. The considerably higher capacity observed in discharge suggests two possible scenarios: (i) a portion of the discharge current is consumed by side reactions, such as SSE decomposition or interfacial degradation, or (ii) the charge process is unable to fully extract the lithium inserted during discharge. Both possibilities affect the long-term cycling stability and raise the question of chemical/electrochemical compatibility between FeNb<sub>11</sub>O<sub>29</sub> and Li<sub>6</sub>PS<sub>5</sub>Cl.

Regarding capacity, a clear improvement is observed as the VGCF content decreases. In the second cycle at C/12, the charge capacities of VGCF-30, VGCF-20, VGCF-10, and VGCF-05 were ~145, 170, 195, and 212 mAh g<sup>-1</sup>, respectively. This superior performance of VGCF-05 is also evident in the dQ/dV plot, showing a stronger Nb<sup>5+/4+</sup> peak and, more prominently, a broader feature below 1.5 V. The reduced capacity at higher VGCF content may result from (i) surface functional groups on VGCF that promote irreversible reactions with Li<sub>6</sub>PS<sub>5</sub>Cl [40], and (ii) the

poor densification of VGCF relative to FeNb<sub>11</sub>O<sub>29</sub> or Li<sub>6</sub>PS<sub>5</sub>Cl, which may increase porosity in the composite pellets. On the other hand, lowering the VGCF content too far (VGCF-03) causes the discharge capacity to drop again (~185 mAh g<sup>-1</sup>), likely due to insufficient electronic conductivity.

Cell polarisation can be quantitatively estimated from the voltage separation between the dQ/dV peaks during charge and discharge. Generally, lowering the VGCF content reduces polarisation, although this trend is not strictly monotonic. High-VGCF composites such as VGCF-30 and VGCF-20 exhibit higher polarisation (80–100 mV), whereas VGCF-10, VGCF-05, and VGCF-03 show lower values (50–60 mV), as shown in the inset of Fig. 2d. The polarisation trend is consistent with the electrochemical behaviour, as the composite with intermediate VGCF content (VGCF-05) has low polarisation and also delivers the highest reversible capacity.

To further evaluate the performance of FeNb<sub>11</sub>O<sub>29</sub> composites under practical cycling conditions, rate capability tests were carried out by increasing the current stepwise from C/12 to 4.2C, followed by a return to the initial rate (Fig. 2e). Across the entire current range, the VGCF-05 composite consistently delivers the best performance, achieving >220 mAh g<sup>-1</sup> at C/12 and retaining ~80 mAh g<sup>-1</sup> at 0.8C. Besides, reproducibility test on four VGCF-05-based cells shows highly consistent rate behaviour with small error margins, confirming the reliability of these tests (Fig. S4). However, capacity retention remains limited, with a gradual decline observed, especially at higher current rates.

By contrast, high-VGCF composites (VGCF-30 and VGCF-20) exhibit substantially lower capacities and more severe capacity losses with increasing current, likely due to reduced contact between FeNb<sub>11</sub>O<sub>29</sub> and Li<sub>6</sub>PS<sub>5</sub>Cl in the porous carbon network. The VGCF-10 composite exhibits intermediate performance, whereas VGCF-03 performs reasonably well at low current rates but its capacity fades rapidly at higher rates, likely due to insufficient electronic conductivity. When the current rate is returned C/12, all composites show partial capacity recovery, with VGCF-05 and VGCF-10 recovering most effectively.

Overall, the trends in capacity, CEs, polarisation and rate capability suggest that the best performance is achieved for the FeNb<sub>11</sub>O<sub>29</sub> composite with intermediate VGCF content. Specifically, VGCF-05 delivers the highest capacity and best rate performance, raising the question of how ion and electron transport affect the performance of solid-state cells.

To further investigate these transport properties, the best-performing (VGCF-05) and the VGCF-richest (VGCF-30) composites were examined using electron- and ion-blocking electrode setups and DC polarisation method (Fig. 3). For VGCF-05, the electronic conductivity is σ<sub>e</sub> ≈ 337 mS cm<sup>-1</sup>, about three orders of magnitude higher than the ionic value (σ<sub>ion</sub> ≈ 0.110 mS cm<sup>-1</sup>). Although modest, this ionic conductivity is sufficient to sustain Li<sup>+</sup> transport due to the well-dispersed FeNb<sub>11</sub>O<sub>29</sub> and Li<sub>6</sub>PS<sub>5</sub>Cl particles as observed in SEM (Fig. S7), consistent with its strong rate capability. In VGCF-30, the higher VGCF fraction yields an even greater σ<sub>e</sub> (~915 mS cm<sup>-1</sup>), but σ<sub>ion</sub> is slightly lower (0.097 mS cm<sup>-1</sup>), suggesting that excessive carbon disrupts the Li<sub>6</sub>PS<sub>5</sub>Cl network and increases tortuosity for Li<sup>+</sup> migration.

When compared with the rate capability of FeNb<sub>11</sub>O<sub>29</sub> in liquid electrolyte coin-cells at similar mass loading, solid-state cells exhibit inferior performance, as shown in Fig. S4. Using cyclic voltammetry, apparent lithium-ion diffusivities for solid-state cell are estimated to be ~4 × 10<sup>-9</sup> cm<sup>2</sup> s<sup>-1</sup>, one order of magnitude lower than those of liquid electrolyte-based cell (~2 × 10<sup>-8</sup> cm<sup>2</sup> s<sup>-1</sup>) (Fig. S5). This reduced diffusivity and weaker rate performance could arise from limited ionic conductivity of the solid-state electrolyte compared with the liquid, as well as interfacial or microstructural limitations inherent to the solid-state contacts. This is supported by a simple diffusion-limited current estimate (≈0.30 vs. 2.8 mA cm<sup>-2</sup> for solid-state and liquid systems, respectively), suggesting that pronounced capacity losses occur when the applied current exceeds these limits (Note S1).

Despite these limitations, FeNb<sub>11</sub>O<sub>29</sub> shows strong promises when
<figure><img src="image_7.png" /><figcaption>Fig. 3. DC polarisation measurements for electronic and ionic conductivities of VGCF-05 and VGCF-30 composites. Left: Applied stepwise voltages (red) and corresponding current responses (blue). Right: Linear fits of current vs. voltage to determine resistance and calculate conductivity. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>

compared with several high-voltage negative electrode materials in Li-SSBs reported previously [24–31]. Most reported cells typically used electrodes of low to moderate mass loadings (0.5–3 mg cm⁻²), fabricated by various formulations and methods (slurry casting, uniaxial pressing, or pulsed laser deposition), often achieved higher capacities than FeNb₁₁O₂₉-based cells at comparable or higher current densities (see Table S3). However, the electrodes in this study were uniaxially pressed with a much high active material loading (~10 mg cm⁻²) yet still delivers ~210 mAh g⁻¹ at C/12 and ~80 mAh g⁻¹ at 0.8C, which are competitive performance within the context of solid-state systems.

Overall, VGCF-05 provides the most effective balance of ionic and electronic transport, enabling both high capacity and good rate capability. This composite is therefore selected for further analysis.

## 3.2. Cycling performance and asymmetric rate test

Fig. 4a presents the cycling performance of the optimised composite, averaged over four cells, at a moderate rate of C/6. The cells deliver an initial charge capacity of ~127 mAh g⁻¹ with an iCE of ~82 %. From the 4th cycle, CE value remains consistently high at ~98–99.8 %, indicating largely reversible intercalation after some initial cycles. Capacity

remains above 100 mAh g⁻¹ for 30 cycles before gradually declining to 87 mAh g⁻¹ after 50 cycles, equivalent to about 68 % of the initial capacity. Despite its favourable transport characteristics, the composite still faces limitations during cycling, likely due to interfacial degradation, cumulative side reactions, or contact loss.

Asymmetric rate tests were conducted on half-cells containing VGCF-05 to compare the kinetics of Li⁺ insertion and extraction (Fig. 4b–f). At low currents (C/12 and C/6), both cases deliver similar capacities, indicating that lithium transport in FeNb₁₁O₂₉ is not kinetically limited under these conditions. However, from 0.4C onward, a clear asymmetry emerges: higher capacity is clearly observed for C-rate test in all cycles compared to the D-rate test. In the C-rate test, capacities remain relatively well retained, with ~50 mAh g⁻¹ (~20 % of the initial value) still accessible at 1.7C. In contrast, the D-rate test shows a rapid capacity fading beyond 0.4C, with almost no capacity remaining at 1.7C.

This asymmetry is further supported by the dQ/dV plots, where increasing current rates in both tests causes progressive broadening and suppression of the Nb⁵+/⁴⁺ and Nb⁴+/³⁺ peaks, along with gradual shifts in peak position. These distortions are more pronounced in the D-rate test, reflecting larger overpotentials during discharge and sluggish Li⁺ insertion kinetics under fast rates. Such behaviour is also observed in
<figure><img src="image_8.png" /><figcaption>Fig. 4. Cycling performance (a) and asymmetric rate tests of the optimised composite (b), with voltage profiles and differential capacity ($dQ/dV$) curves during C-rate (c–d) and D-rate tests (e–f).</figcaption></figure>

FeNb<sub>11</sub>O<sub>29</sub>-based liquid electrolyte cells (Fig. S6) and has also been reported for graphite electrodes [165].

Constant current–constant voltage (CC–CV) cycling was then applied to mitigate these kinetic limitations (Fig. 5). A constant-voltage (CV) hold was introduced at the end of each discharge step to ensure complete lithiation of FeNb<sub>11</sub>O<sub>29</sub>. At low rates (C/12–0.4C), most of the capacity is originated from the constant-current (CC) step, approaching values seen in liquid-electrolyte coin cells. Beyond 0.8C, however, the CC contribution drops sharply, and at 4.2C it becomes negligible, reflecting insufficient time for complete lithiation under high currents. The application of a CV step effectively relieves kinetic limitations and promotes deeper lithium insertion. As the current rate increases and the CC step becomes less effective, the CV step accounts for a progressively larger contribution to the total capacity (Fig. 5b). Importantly, capacity fully recovers when the current returns to C/12, confirming that the CV step enables reversible lithiation rather than driving parasitic reactions. Overall, these findings demonstrate that the intrinsic kinetic barrier to

lithium insertion at high current rates can be partially mitigated by incorporating a CV hold, thereby improving capacity utilisation and achieving more complete lithiation.

### 3.3. Pre- and post-cycling analyses

To investigate the evolution of the electrode during operation, pre- and post-cycling analyses were carried out to examine changes in phase composition, morphology, and the electronic states of the elements in the optimised composite. *Ex situ* XRD measurements were performed on the pristine and cycled composite. Le Bail refinements of their XRD patterns (Fig. 6) reveal more pronounced changes in the cell parameters of FeNb<sub>11</sub>O<sub>29</sub>, compared to little changes in those of Li<sub>6</sub>PS<sub>5</sub>Cl. Specifically, FeNb<sub>11</sub>O<sub>29</sub> shows slight but consistent expansion along the *a*, *b*, and *c* axes after cycling (Table S4), suggesting that a small amount of lithium remains trapped within the host framework at the end of charge. In contrast, the lattice parameter of Li<sub>6</sub>PS<sub>5</sub>Cl remains nearly unchanged,
<figure><img src="image_9.png" /><figcaption>Fig. 5. Voltage profiles of the VGCF-05 composite during the CC-CV test at various rates, with CV step applied at the end of discharge (a), and corresponding capacity contributions from CC (orange) and CV (green) steps (b). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>

indicating the structural stability of the bulk electrolyte during cycling. However, its interaction and chemical stability with other components of composite electrode, i.e., FeNb<sub>11</sub>O<sub>29</sub> and VGCF, need to be investigated with surface-sensitive techniques.

SEM and EDS were used to examine the microstructure and homogeneity of (i) pristine, (ii) densified, and (iii) cycled VGCF-05 composites. The top and middle rows show secondary electron (SE) images at high and low magnifications, revealing surface morphology and texture, while the bottom row shows backscatter electron (BSE) images that highlight heavier elements Nb/Fe in FeNb<sub>11</sub>O<sub>29</sub> (Fig. S7).

At high magnification, the fibrous VGCF structure is visible, with fibres ~0.1 µm in diameter and several micrometres (~4 µm) in length, in close contact with FeNb<sub>11</sub>O<sub>29</sub> and Li<sub>6</sub>PS<sub>5</sub>Cl particles. Medium-magnification SE images reveal similar surface texture in three samples, with VGCF fibres uniformly distributed in the matrix. In BSE images, bright regions attributed to FeNb<sub>11</sub>O<sub>29</sub> are evenly dispersed without significant agglomeration or phase segregation, indicating uniform mixing of three composite components. FeNb<sub>11</sub>O<sub>29</sub> morphology and particle size remain unchanged after cycling.

These SEM images are further supported by EDS mapping (Fig. S8), also showing homogeneous elemental distribution, suggesting that all

three samples exhibit a uniform phase dispersion and interconnected network, supporting efficient ionic and electronic transport. However, possible interfacial ageing during cycling may occur but is not evident in SEM.

To probe changes in electronic and chemical states at the interface, XPS was also performed on three samples: pristine FeNb<sub>11</sub>O<sub>29</sub> powder, compacted and cycled composites (Fig. 7 and Table S5).

For pristine FeNb<sub>11</sub>O<sub>29</sub>, the binding energy values confirm Nb<sup>5+</sup> in an oxide environment. Notably, iron is barely detectable on its surface, possibly due to its low stoichiometric amount. Nb 3d peak positions remain unchanged in compacted and cycled composites, suggesting little chemical degradation of FeNb<sub>11</sub>O<sub>29</sub> during cycling and a total reversible electrochemical behaviour of Nb on the electrode surface during cycling. Li<sub>6</sub>PS<sub>5</sub>Cl in compacted composite preserves its chemical integrity since peaks positions and atomic ratios are identical to pure argyrodite, which is not the case for Li<sub>6</sub>PS<sub>5</sub>Cl in the composite after cycling.

Notable changes are observed in the XPS spectra:

* C 1s: CH<sub>2</sub>, CO, and CO<sub>2</sub> contaminants are detected in all samples. VGCF signals appear in two composites, with lower intensity in the cycled sample.

* O 1s, Nb 3d: In the compacted sample, O-Nb and Nb 3d signals slightly decrease in intensity, indicating partial surface coverage of FeNb<sub>11</sub>O<sub>29</sub> particles by other components. After extended cycling, the O-Nb peak disappears and Nb 3d intensities become considerable low, suggesting the formation of a new surface layered.

* S 2p, P 2p, Cl 2p: These Li<sub>6</sub>PS<sub>5</sub>Cl-related peak positions remain unchanged across all samples. Elemental ratios in the compacted sample match those of the pristine state, confirming negligible changes during processing. In contrast, the cycled sample shows the emergence of new signals corresponding to the formation of sulfur oxide species (SO<sub>x</sub>), S<sup>0</sup> and Li<sub>2</sub>S during cycling, as previously reported [49].

These findings suggest the formation of a chemically modified surface layer, likely a solid-electrolyte interface (SEI), rich in sulfur oxides, S and Li<sub>2</sub>S. On the negative electrode, Li<sub>6</sub>PS<sub>5</sub>Cl has been reported to decompose into Li<sub>2</sub>S, LiCl and Li<sub>3</sub>P [28,49]. These products, especially Li<sub>2</sub>S, with poor ionic and electronic conductivities, has been shown to significantly increase interfacial resistance at the negative electrode material/Li<sub>6</sub>PS<sub>5</sub>Cl contact [50].

To monitor this effect, EIS was performed in the solid-state half-cells at different stages: (A) pristine, (B) after the rate test, and (C) after 50 cycles at C/6. Fig. 8 presents the Nyquist plots together with the distribution of relaxation times (DRT) analysis, which decomposes the impedance spectra into characteristic frequencies: each peak corresponds to a distinct process, and its area reflects the associated impedance.

Because of the system's complexity (many components and interfaces involved) and significant overlap in the relaxation times of the processes (Fig. 8b), a complete deconvolution of these contributions is not possible. Therefore, analyses focus on impedance magnitude and characteristic frequency at the semicircle maximum ($f_{peak}$), which defines the process's characteristic relaxation time, $\tau = 1/(2\pi f_{peak})$.

Nyquist plots exhibit two overlapping depressed semicircles, each modelled by a resistor (R) in parallel with a constant-phase element (CPE). These spectra can be fitted using the equivalent circuit in Fig. 8, which can be assigned to (i) resistance of the bulk SSE (R<sub>1</sub>), (ii) interfacial resistance and non-ideal capacitance (R<sub>2</sub>//CPE<sub>2</sub>), (iii) a charge-transfer resistance in parallel with the electrode double-layer capacitance (R<sub>3</sub>//CPE<sub>3</sub>), and (iv) diffusion inside SSE particles (CPE<sub>4</sub>) [26]. The CPE is characterised by its $\alpha$ parameter ($0 < \alpha \le 1$), which quantifies the deviation from the ideal capacitive behavior (when $\alpha = 1$), due to electrode inhomogeneity.

At high frequencies, the x-axis intercept of these spectra, which
<figure><img src="image_10.png" /><figcaption>Fig. 6. Le Bail-refined XRD patterns of pristine (top) and cycled (bottom) VCGF-05 composite.</figcaption></figure>

<figure><img src="image_11.png" /><figcaption>Fig. 7. XPS peak deconvolution of C 1s, O 1s, Nb 3d, Cl 2p, S 2p, and P 2p regions showing binding energies and chemical states in pristine FeNb<sub>11</sub>O<sub>29</sub>, compacted and cycled VGCF-05 composite.</figcaption></figure>
<figure><img src="image_12.png" /><figcaption>**Fig. 8.** EIS of solid-state half-cells: pristine (blue), after rate test (red), and after long cycling (green). (a) Nyquist plots with equivalent circuit (inset). Characteristic frequencies are highlighted. (b) DRT analysis, resolving the impedance into frequency-dependent processes. The spectra reveal progressive increases in ohmic, interfacial/charge-transfer resistance, and a shift of the dominant relaxation toward lower frequencies upon cycling. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)</figcaption></figure>

represents total Ohmic resistance ($R_1$), shifts progressively to higher $Z'$ value during cycling, reflecting a gradual increase in bulk ionic resistance. The initial spectrum (A) consists of two depressed semicircles, with $f_{peak}$ of ~400 Hz and 5 Hz, corresponding to a fast interfacial and slow charge-transfer processes, respectively. In DRT analysis, this slow process manifests as small peaks near 1 Hz in all three cases, with their peak areas increasing during cycling. As cycling progresses, the semicircles grow significantly in size, with $Z'$ rising from ~250 Ω (A) to above 700 Ω (B) and finally extending beyond 1500 Ω (C). This agrees with the substantial increase in the fitted resistance values shown in

**Table 2**

Impedance fitting results: $R_1$ (ohmic resistance), $R_2$ and $R_3$ (interfacial/charge-transfer processes) with corresponding $\alpha_{CPE}$ parameter and characteristic frequency (f) for each parallel R//CPE element.

<table><thead><tr><th></th><th>R<sub>1</sub></th><th>R<sub>2</sub>/α<sub>CPE2</sub>/f<sub>2</sub></th><th>R<sub>3</sub>/α<sub>CPE3</sub>/f<sub>3</sub></th></tr></thead><tbody><tr><td>Pristine</td><td>63(1) Ω</td><td>42(1) Ω/0.74(1)/728 Hz</td><td>109(1) Ω/0.43(1)/1.5 Hz</td></tr><tr><td>After rate test</td><td>128(2) Ω</td><td>228(9) Ω/0.26(2)/113 Hz</td><td>204(2) Ω/0.55(1)/45 Hz</td></tr><tr><td>After prolonged cycling</td><td>180(1) Ω</td><td>279(2) Ω/0.19(5)/144 Hz</td><td>1180(14) Ω/0.49(1)/4.7 Hz</td></tr></tbody></table>

**Table 2.** Among them, the resistance corresponding to the charge transfer process ($R_3$) increases significantly after the long-term cycling test. Furthermore, the peak frequency shifts from 400 Hz (A) to 72 Hz (B) and finally to 4 Hz (C), indicating a progressive transition toward a slower process with longer relaxation time. This trend is supported by the emergence of pronounced mid- and low-frequency features in the DRT plot. Such behaviour indicates increasing relaxation time, likely due to SEI formation. In all cases, the semicircles are highly compressed and CPE $\alpha$-values deviates significantly from the ideal value of 1, suggesting that several semicircles, as in a transmission line model, would be required to fit the interfacial resistance. Herein, the transmission line model is thus simplified by using a CPE to represent an average sum of all the interfacial resistances. The existence of several semicircles indicates that the system is inhomogeneous after cycling as the SEI formation does not occur uniformly in all particles.

Overall, the gradual increase of ohmic resistance ($R_1$) and the growing dominance of the slow process indicate increased $Li_6PS_5Cl$ ionic resistance and a substantial interfacial layer build-up. This SEI layer, rich in $Li_2S$, $S^0$, and $SO_x$ species as confirmed by XPS, becomes thickened during cycling and impedes $Li^+$ transport across the $FeNb_{11}O_{29}/Li_6PS_5Cl$ interface, thereby accelerating impedance growth and hindering charge-transfer kinetics.

## 4. Conclusions

This study systematically investigates the electrochemical performance and interfacial behaviour of $FeNb_{11}O_{29}$-based negative electrodes in Li-SSBs. By tailoring electrode compositions, the optimum balance between the electroactive material, $Li_6PS_5Cl$ SSE and VGCF conductive additive was identified to maximise its performance. Among the formulations tested, the optimised composite, VGCF-05, delivered the most favourable results with the highest reversible capacity, low polarisation, and satisfying rate capability. This optimal behaviour stems from a balanced electron-ion transport network, as supported by conductivity measurements. Asymmetric rate protocols revealed kinetic limitations of $FeNb_{11}O_{29}$ in Li-SSBs, particularly during lithium insertion at high rates. These could be partially mitigated by introducing a constant-voltage step during discharge, which enhanced lithium intercalation and improved capacity retention under fast-cycling conditions. Pre- and post-cycling characterisations further demonstrated the formation of an interfacial layer, rich in sulfur-containing species, suggesting chemical degradation at the electrode-electrolyte interface.

Overall, this work highlights the critical importance of microstructural optimisation in composite electrodes for solid-state batteries. While increasing VGCF content enhances electronic conductivity, its excessive amounts may hinder ionic transport or disrupt effective contact between the active material and the solid electrolyte, leading to higher polarisation and lower performance. Achieving a balanced electron-ion transport network is therefore essential for achieving high capacity, low overpotential, and long-term cycling stability. These findings provide valuable insights for the design of electrode composite architectures and highlight the strong promises of $FeNb_{11}O_{29}$ and related Wadsley-Roth phases as negative electrode materials in solid-state batteries. Future works can focus on surface coating to mitigate argyrodite decomposition (e.g. $LiWS_2$ [28]) or exploring more robust SSEs such as oxyhalide compounds (e.g. $LiTaOCl_4$ which is stable at low voltage [51]), to further improve durability and rate capability.

### CRediT authorship contribution statement

**Dat Le Thanh:** Writing – original draft, Formal analysis, Data curation. **Dominique Foix:** Formal analysis, Data curation. **Emma-nuelle Guard:** Writing – review & editing, Supervision, Methodology, Funding acquisition, Conceptualization. **Romain Berthelot:** Writing – review & editing, Supervision, Methodology, Funding acquisition, Conceptualization. **Long Hoang Bao Nguyen:** Writing – review &
editing, Supervision, Methodology, Conceptualization.

## Declaration of competing interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

## Acknowledgements

As a part of the DESTINY PhD program, this publication is acknowledged by funding from the European Union#8217; s Hori-zon2020 research and innovation program under the Marie Skłodowska-Curie Actions COFUND (Grant Agreement #945357). Funding from the French National Research Agency (STORE-EX Labex Project ANR-10-LABX-76-01) is gratefully acknowledged. The authors want to acknowledge Frédéric Fernandez (UAR 2041, PAC Chimie Balard, Montpellier) and Amandine Guiet (Institut des Molécules et Matériaux du Mans, UMR 6283 CNRS, France) for their help with the SEM measurement.

## Appendix A. Supplementary data

Supplementary data to this article can be found online at https://doi.org/10.1016/j.jpowsour.2025.239120.

## Data availability

Data will be made available on request.

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