# Domain Size Modulation of Clay–Nanocellulose Composites for Enhanced Na-Ion Transport

Sneha Mandal, Catherine Tom, Yuvraj Singh, Hassan Khoder, Pascale Launois, Erwan Paineau, Stephan L Grage, Mareen Schaller, Subbiah Alwarappan, Rakesh S Singh, et al.

► To cite this version:

Sneha Mandal, Catherine Tom, Yuvraj Singh, Hassan Khoder, Pascale Launois, et al. Domain Size Modulation of Clay–Nanocellulose Composites for Enhanced Na-Ion Transport. ACS Applied Energy Materials, 2025, 8 (9), pp.6071-6086. 10.1021/acsaem.5c00452. http://onlinelibrary.wiley.com/doi/abs/10.1021/acsaem.5c00452

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# Domain Size Modulation of Clay-Nanocellulose Composites for Enhanced Na-ion Transport

Sneha Mandal¹, Catherine Tom², Yuvraj Singh², Hassan Khoder³, Pascale Launois³, Erwan Paineau³, Stephan L. Grage⁴, Mareen Schaller⁵, Subbiah Alwarapan⁶, Rakesh S. Singh¹, Ravi Kumar Pujala² and Vijayamohanan K. Pillai¹

1. Department of Chemistry, Indian Institute of Science Education and Research (IISER), Tirupati-517619, India

2. Department of Physics, Indian Institute of Science Education and Research (IISER), Tirupati-517619, India

3. Université Paris-Saclay, CNRS, Laboratoire de Physique des Solides, 91405 Orsay, France

4. Karlsruhe Institute of Technology, Institute of Biological Interfaces IBG-2, P.O. Box 3640, 76021 Karlsruhe, Germany

5. Institute for Applied Materials – Energy Storage Systems, Karlsruhe Institute of Technology, Hermann-von-Helmholtz-Platz 1, 76344 Eggenstein-Leopoldshafen, Germany

6. Electrodes and Electrocatalysis Division, CSIR-Central Electrochemical Research Institute, Karaikudi- 630003, Tamil Nadu, India.

* Corresponding author- Vijayamohanan K. Pillai, vijay@iisertirupati.ac.in

## Abstract:

Rechargeable Batteries using alkali metal anodes garner significant interest as a transformative energy storage technology for grid and mobile applications. However, scanty resources of Lithium (Li) have pushed the research focus towards the realm of abundant alkali and alkaline earth metals such as Sodium (Na), Potassium (K), Magnesium (Mg), and Aluminium (Al). Moreover, safety issues relating to inflammability, dendritic growth, limited cycle life, and capacity fading in liquid electrolytes seriously hamper their extensive deployment. Here, we propose a novel, flexible, solid-polymer-composite electrolyte thin film (thickness ~ 20-25 µm) with enhanced sodium ion transport characteristics based on cellulose nanocrystals (CNC) and Laponite (Lap)/Montmorillonite (MMT) clay, which can be mass-produced at room temperature. This composite electrolyte has important flame retarding and mechanical properties coupled with excellent Na ion conductivity. From a critical analysis of the preliminary results using various compositions and optimized architecture of the composite electrolyte, an ionic conductivity of the order of $10^{-4}$ S cm$^{-1}$ at 50 °C is obtained with enhanced stability window. In parallel, a good electrochemical response was observed, indicating potential application in all-solid-state batteries (SSBs). We also
conducted molecular dynamics (MD) simulations of randomly packed CNC and MMT
composites to estimate the ionic diffusivity using the Einstein equation, and the estimated
conductivity is qualitatively aligned with the experimental results. Overall, these
nanocomposite films open game-changing perspectives for fabricating flexible power sources
from abundant resources.

<figure><img src="image_2.png" /></figure>

**Keywords-** Polymer composite electrolyte, clay, cellulose nanocrystals, sodium ion transport.

**1. Introduction**

The clean energy transition is a pivotal strategy in today's world to combat climate change, achieve sustainability, reduce over-dependence on fossil fuels, and, more importantly, accomplish a carbon-zero footprint. Consequently, deployment of more efficient energy storage devices have become an indispensable component of our daily lives, particularly with the widespread adoption of advanced technology to meet the escalating energy needs of the expanding population, which is projected to exceed 3 TWh annually by 2030.¹,² The incorporation of energy storage devices, such as batteries, supercapacitors, and fuel cells contribute to the synergy, modulating the flexibility of power systems, enabling the optimal utilization of variable energy sources such as Solar Photo-Voltaic (SPV) devices, with a projected capacity of over 5400 GW, and wind energy, with a global capacity expected to surpass 3500 GW by 2030.³,⁴ Significant efforts are underway to develop diverse battery materials and chemistries including solid electrolytes, alkali metal anodes, sulphur cathodes, and composite materials for various energy storage requirements across different sectors, from portable/wearable devices for consumer, computing, and communications (3C) to more demanding electric vehicles, and smart grids.⁵,⁶ In this context, solid-state batteries (SSBs) have emerged as a prominent area of research that can potentially address many of the
challenges associated with conventional Li-ion batteries (LIBs). By utilizing solids instead of liquid or gels as electrolytes, SSBs provide several advantages, including enhanced safety, high energy density, rapid charging, low cost, longer life span, and improved performance at extreme temperatures. Moreover, it can potentially reduce our dependence on rare or expensive elements and materials of high geopolitical or ethical risks, thereby enhancing sustainability and more energy-efficient future applications.⁷,⁸

The discovery of ionic conductivity in poly(ethylene oxide) (PEO) with different Na and K salts during the 1970s, along with the commercialization of materials like lithium phosphorous oxynitride (LiPON) that demonstrates stability against Li-metal anodes and extended cyclability with high coulombic efficiency in thin film batteries, has positioned solid electrolytes as safer and more energy-dense materials.⁹⁻¹¹ Despite sufficient mechanical strength of most composite polymer solid electrolytes, many still suffer from limited ionic conductivity (10⁻⁵ to 10⁻⁸ S cm⁻¹) at room temperature and interfacial instability.¹²⁻¹⁴ On the other hand, solid oxide electrolytes such as NASICON (Li₁.₅Al₀.₄Cr₀.₁Ge₁.₅(PO₄)₃)¹⁵, and garnet (Li₇La₃Zr₂O₁₂)¹⁶ show room temperature conductivity of 6.65 and 0.51 mS cm⁻¹ respectively with a wide electrochemical stability window. However, NASICONs are chemically unstable with Li metal anode and have low thermal resistance (~300 °C), while garnets require sparse lanthanides, and high sintering temperatures >1000 °C. This incurs additional costs and limits the compatibility with certain cathode active materials as the catholyte in an annealing process. Sulphide solid electrolytes like LISICON (Li₉.₅₄Si₁.₇₄P₁.₄₄S₁₁.₇Cl₀.₃)¹⁷, thio-LISICON (Li₃.₂₅Ge₀.₂₅P₀.₇₅S₄)¹⁸ and argyrodites (Li₅.₅PS₄.₅Cl₁.₅)¹⁹ display room temperature conductivity of 25, 2.2 and 12 mS cm⁻¹ respectively. However, their low electrochemical stability, results in high interfacial resistances both at the anode and cathode sides, thereby, causing degradation problems for high-rated applications. Furthermore, the stability of solid electrolytes with lithium metal and the prevention of dendrite formation remain significant challenges. Moreover, many metal oxides and sulphides do not possess suitable mechanical properties to accommodate the cyclic stresses that solid-state batteries undergo as a result of volumetric changes in the cathode during the charge-discharge processes. Therefore, developing more robust solid electrolytes is crucial for advancing energy storage systems.

In this context, natural clays stand out as promising materials due to their abundance and propensity for chemical modification to impart flexibility. These materials offer advantages such as porous structures, tuneable surface area, excellent thermal and mechanical stability,
and cost-effectiveness. Notably, their high ionic conductivity and hydrophilicity make them ideal for solid-state electrolytes.²⁰⁻²³ In this work, we describe a novel, flexible composite polymer electrolyte film with a thickness of 20-25 µm, based on a clay-nanocellulose structure that includes Laponite (Lap) or Montmorillonite (MMT) saturated with sodium ions in their interlayer spaces. The resulting composite exhibits ionic conductivity exceeding 10⁻⁵ S cm⁻¹ at room temperature and 10⁻⁴ S cm⁻¹ at temperatures above 50 °C. This design results in a high cation-exchange capacity and expanded interlayer spacing within the clay nanosheets. The negatively charged 1-D cellulose nanocrystals (CNCs) interact electrostatically (Zeta potential, ζ = -40 mV) with the positively charged edges of the 2-D clay nanosheets to form a layered structure, which reduces polymer crystallinity and facilitates ion transport within the interlayers and the amorphous polymer chains. The interaction between CNC and the clays varies depending on the clay dimension, with a significantly smaller platelet size in Laponite (25 nm) than in Montmorillonite (250 nm). The smaller size of Laponite platelets promotes greater connectivity with CNC and provides a larger surface area of positively charged edge sites to interact with the negatively charged CNC. This size difference enables effective networking among CNCs, enhancing the film's optical transparency, tensile strength, and thermal stability.²⁴,²⁵ These films are easy to process, air-stable, have a wide electrochemical stability window, and can be mass-produced at room temperature. Preliminary results using a coin-cell configuration provide promising results. However, more accurate and detailed experiments are needed for a prototype device capable of reasonable charge-discharge cycles with good energy and power density.

<figure><img src="image_3.png" /></figure>
Scheme 1: Illustration of the synthetic methodology for preparing free-standing, thin, flexible, and transparent composite polymer electrolytes of CNC-Laponite and CNC-Montmorillonite.

Scheme 1 represents the fabrication methodology for free-standing and flexible thin films using suitable compositions of CNC and Laponite or Montmorillonite in water. Figure S1(a and b) shows the reflectance curves for both composite polymer electrolyte thin films with varying concentrations of CNC-Lap and CNC-MMT. The results indicate that the films are transparent over the investigated wavelength range. Figure S1(c-n) displays characteristic images of these composite films of CNC-MMT and CNC-Lap.

<figure><img src="image_4.png" /></figure>
<figure><img src="image_5.png" /><figcaption>**Figure 1:** Temperature-dependent X-ray diffractograms (XRD) of CNC-MMT composite and CNC-Lap composite polymer electrolyte showing the phase changes (253-353 K), (a) 1C:2L, (b) 1C:1L, (c) 2C:1L, (d) 2.5C:1M, (e) 1C:1M, (f) 1C:2M. Raman Spectra of (g) CNC-Lap composite and (h) CNC-MMT composites.</figcaption></figure>

X-ray diffraction (XRD) and Raman spectroscopy unambiguously reveal that CNC molecules are intercalated within the clay layers, enhancing significantly the ionic conductivity and mechanical properties. Indeed, the MMT and Laponite nanoplatelets are organized effectively into macroscopic layered structures with varying degrees of CNC-MMT and CNC-Lap content. The nano-structural features and physicochemical interactions between CNC-MMT and CNC-Lap determine their multifunctional performance. Temperature-dependent XRD analysis (Figure 1a-f) shows a shift in the peaks of the composites beyond 298 K, reflecting an expansion of the crystal lattice, or structural reorganization associated with a phase change. The intensity also decreases with the temperature increase indicating disordering or partial amorphization of the composite. Due to the phase change, the lower angle peaks (inset) corresponding to 2θ – 2-3° show specific interlayer shifts with thermal expansion. Accordingly, Figure S2(a and b) and Table S1 show the (001) plane with an enhanced d-spacing from 13.16 to 15.55 Å and 12.3 to 14.38 Å for Lap and MMT respectively, as a result of the intercalation of CNC into the interlayer spacing of 2D clay nanosheets. This is also bolstered by a sudden increase in ionic conductivity (~298 K) observed in EIS and supported by the molecular dynamics simulations. Notably, this change in conductivity is monotonic
(vide infra Arrhenius plot), and this finding is further corroborated by the presence of peaks corresponding to amorphous nanocellulose in the clay-CNC composite, indicating a close interaction between the CNC molecules and the clay sheets, as observed in SAXS and solid-state NMR data (SS-NMR) (vide infra).

Raman spectroscopy, known for its high spatial resolution, is an effective tool for examining local structural variations and heterogeneities within composite materials. In crystalline substances, molecular vibrations split due to intermolecular interactions within the crystalline lattice. These split vibrations can be used to determine the properties of the crystalline phase, offering insights into structural aspects such as shape and orientation. Figure 1 (g and h) presents the Raman spectra of CNC-Lap and CNC-MMT composites. The spectra reveal characteristic peaks for both Laponite and MMT, maintaining the phyllosilicate structure of 2:1 ($T_d-O_h-T_d$) sheets.²⁶⁻³⁰ For MMT, peaks within the 50-200 cm⁻¹ range (144, 187, 214 cm⁻¹) correspond to the $O_h$ symmetry of $AlO_6$ or $MgO_5OH$ octahedral sheets. Peaks from 300-1200 cm⁻¹ are attributed to the $T_d$ symmetry of $SiO_4$ tetrahedral sheets. The peak at 3648 cm⁻¹ is assigned to OH vibrations, while the broad envelope of peaks below 3600 cm⁻¹ (centred at 2676 cm⁻¹) is attributed to OH vibrations in $H_2O$ molecules. The peak at 70 cm⁻¹ is also associated with $Na^+$ in the interlayer. Notably, there is a slight downshift in the peaks within the composites due to CNC's interaction with MMT.³¹

For Laponite, peaks in the 150-300 cm⁻¹ range (181, 291 cm⁻¹) correspond to $O_h$ sheets $MO_6$. The spectra exhibit three bands around 361, 423, and 463 cm⁻¹, corresponding to (Li, Mg)-O-Si bending. Sharp peaks at 677 and 1095 cm⁻¹ correspond to $\delta_{(Si-Si)O-OH}$ and Si-O stretch in the 2:1 layer, respectively. The broad peak in the 3200-3700 cm⁻¹ region includes various O-H stretching modes from metal-hydroxyl complexes and water adsorbed in different conformations: 3691 and 3712 cm⁻¹ (Mg/Li-OH), 3627 cm⁻¹ (Si-OH), 3470 cm⁻¹ (inner sheet M-OH), 3355 cm⁻¹ (amorphous H-OH, bulk water), and 3259 cm⁻¹ (in-phase H-OH with crystalline structure). The peak around 100 cm⁻¹ is attributed to $Na^+$ in the interlayer.³²
<figure><img src="image_6.png" /><figcaption>Figure 2: Attenuated total reflectance Fourier-transform infrared (ATR FT-IR) spectra of a) CNC-Laponite and (b) CNC-Montmorillonite. Survey spectrum of (c) 1C:2L and (d) 2.5C:1M.</figcaption></figure>

The interfacial characteristics were further analysed by Fourier transform infrared (FTIR) and X-ray photoelectron spectroscopies (XPS). In FT-IR, the rise of new peaks or the shift in original peaks could provide useful information about the physicochemical environment of
the functional groups within the nanocomposite. The changes may confirm specific
interactions between the polymer matrix and the nanoparticles, which helps to understand and
improve the nanocomposite design of materials.³³,³⁴ The FTIR spectra clearly show the
vibrational bands of (Figure 2a) CNC-Lap and (b) CNC-MMT. For MMT and Laponite, the
bending vibration band of Si-O at 998 and 960 cm⁻¹ and the stretching vibration band of O-
Si-O at 1119 and 434 cm⁻¹ indicate the layered structure of montmorillonite consisting of
silicon-oxygen tetrahedron respectively. The band at 3628 cm⁻¹ is assigned to the free-OH
stretching region. The broad band at 3453 cm⁻¹ could be ascribed to the overlapping
asymmetric and symmetric H-O-H stretching vibrations of H-bonded water. Further, OH
bending vibration is observed at 1638 and 1631 cm⁻¹ respectively. For CNC, peaks at 1202,
1158, 1059, 1033, and 898 cm⁻¹ could be assigned for the stretching vibration of C-O from
the pyranose ring, while the band at 3336 cm⁻¹ is attributed to the stretching vibration of O-H
of absorbed water and hydroxyl group.³⁵,³⁶ Interestingly, CNC-Lap and CNC-MMT show an
“isosbestic point” for all the concentrations demonstrating that no bonds are formed.³⁷
Although there is no shift in the peaks, the change in intensity of the composite peaks, might
be attributed to the interaction of CNC-Lap and CNC-MMT. This is possible due to the
intercalation of CNC in the interlayer spacing of Laponite and Montmorillonite, which is
also evident from XRD data.

The surface elemental composition and chemical status of CNC-MMT and CNC-Lap are
studied by XPS. The survey spectra (Figure 2c and d), show clear signals of C 1s, Na 1s, O
1s, Li 1s, Si 2p, and Mg 1s for CNC-Lap, while for CNC-MMT, C 1s, Al 2p, Si 2p, O 1s, Na
1s, Fe 2p, and Mg 1s are visible.³⁸⁻⁴⁰ Interestingly, the binding energy of the composite has
slightly shifted to lower values compared to the corresponding signals for pristine
montmorillonite and Laponite (Figure S2 (h, i, p) and Table S2). This shift can be attributed
to an increase in ionicity due to the interaction with CNC. Further, Figure S2(a-g) shows the
deconvoluted C 1s spectra for CNC-MMT with peaks at 285.4, 284.69, and 289.95 eV,
corresponding to C-C sp², C=O, and O-C=O bonds, respectively. The Mg 1s peak at 1302.93
eV corresponds to Mg(OH)₂, indicating the presence of non-exchangeable Mg in the intact
octahedral layer. The deconvoluted Na 1s spectra show two peaks at 1070.8 eV for Na-O
bond while, the peak at 1070 eV represents the Na⁺ intercalated into the carbon (C-Na). The
O 1s spectra at 530.5 eV correspond to the O-Mg/Si bond. The Si 2p spectrum shows a single
peak at 100.9 eV, corresponding to the Si-O₂ bond, while the deconvoluted Fe 2p spectra
show peaks at 727.2, 712.7, 722.7, and 709 eV, corresponding to Fe 2p₁/₂ and Fe 2p₃/₂ (FeO),
respectively.⁴¹ The Al 2p spectra at 73.48 and 72.8 eV correspond to Al-OH and Al-O bonds, respectively.⁴²

For CNC-Lap (Figure S2 i-o), the deconvoluted C 1s spectra exhibit peaks at 284.7 and 283 eV, corresponding to C=O and C-C sp² bonds, respectively. The Mg 1s peak at 1302.93 eV indicates the presence of Mg(OH)₂, signifying non-exchangeable Mg in the intact octahedral layer. The deconvoluted Na 1s spectra show two peaks at 1070.9 eV for Na-O bond while, the peak at 1069.9 eV represents the Na⁺ intercalated into the carbon (C-Na), which is also seen in ss-NMR.⁴³ The O 1s spectrum at 530.5 eV corresponds to the O-Mg/Si bond, while the Si 2p spectra show peaks at 101.11 and 101.93 eV, corresponding to Si-OH (2p₃/₂) and Si-O bonds respectively. The deconvoluted Li 1s spectra show peaks at 48.63 and 49.5 eV, corresponding to LiOH and Li₂O, respectively. Additionally, the peak at 62.5 eV indicates plasmon loss in Li.⁴⁴⁻⁴⁶
<figure><img src="image_7.png" /><figcaption>(c)</figcaption></figure>

<figure><img src="image_8.png" /><figcaption>(b)</figcaption></figure>

<figure><img src="image_9.png" /><figcaption>(d)</figcaption></figure>

<figure><img src="image_10.png" /></figure>

Figure 3: Solid-state <sup>23</sup>Na NMR spectra of Laponite CNC films (a) and Montmorillonite CNC films (b), as well as <sup>23</sup>Na NMR relaxation rates of Laponite CNC films (circles) and montmorillonite CNC film (squares) with different compositions of CNC and clay mineral as indicated. The <sup>23</sup>Na-NMR were acquired at 79.4 MHz Larmor frequency (corresponding to a <sup>1</sup>H Larmor frequency of 300.1 MHz) at 30 °C. Longitudinal relaxation rates R<sub>1</sub> were acquired as a function of temperature at <sup>23</sup>Na Larmor frequency of 79.4 MHz using inversion recovery. The static experiments were performed on strips of the film material, which were placed into a 5 mm quartz glass tube with the planes aligned along the magnetic field. <sup>23</sup>Na relaxation rates were measured at two different sample orientations, with the film normal aligned along (0°) or perpendicular (90°) to the magnetic field (d). The orientation dependent relaxation
rates were determined for two Montmorillonite samples (5C2M and 2C4M), wrapped in paper, at a field strength of 132.3 MHz (corresponding to 500MHz for 1H).

To get insights into the local environment of the sodium ions and their mobility within the composite materials, we acquired solid-state ²³Na NMR spectra of the Laponite CNC and montmorillonite CNC films (Figures 3a and b). For these measurements, strips of film were placed with their planes parallel to the magnetic field and measured statically, i.e., without magic angle spinning, at 79.4 MHz Larmor frequency at 30°C.

Accordingly, the ²³Na NMR spectra of the Laponite CNC samples exhibit a single resonance line near 0 ppm with a slightly asymmetric line-shape possessing an elongated slope on the right (Figure 3a). Interestingly, with increasing the CNC content, the line position shifted up-field by 7 ppm from 3.2 ppm for a CNC:Lap ratio of 2:4 to -3.8 ppm for a ratio of 4:2. Equally, an up-field shift is observed for the montmorillonite samples, where the signal shifted from 6.1 ppm for a CNC:MMT ratio of 2:4 to -3.4 ppm found for a ratio of 4:2 (Figure 3b). The series of different CNC amounts is extended to a CNC:MMT ratio of 5:2, where the signal is found to shift further to -5.0 ppm. In the montmorillonite samples, a second component or signs of a powder line-shape appear at low CNC content at around 21 ppm, which disappears with increasing CNC amount (marked with a red line in Figure 3b). The overall linewidth in the Laponite and montmorillonite samples is found to be rather different. Significantly narrower lines are found for the CNC-Lap samples with linewidths of around 1.0 kHz, whereas the montmorillonite samples possessed linewidths between 3.6 kHz for high CNC amounts and 4.7 kHz for low CNC amounts. The linewidths in the case of Laponite did not depend on the CNC content. In the case of Montmorillonite, the larger linewidth observed for low amount of CNC is probably associated with the second signal component observed here.

Overall, the ²³Na-NMR signal position near to 0 ppm, found for all samples, is close to the ²³Na chemical shift observed for sodium in CNC previously⁴⁷. For smectites, ²³Na chemical shifts between around 5 ppm to -30 ppm have been reported, depending on the oxygen coordination around the sodium ions⁴⁸. Coordination to oxygens of the smectite layers has been found to lead to shielding and a shift of the ²³Na signal to lower frequencies. Full coordination with oxygens of the smectite matrix has been attributed to signals around -30 ppm. Signals around 0 ppm have been described as originating from sodium, where the coordination by oxygen from the inorganic matrix has been replaced by water when the
samples have been measured at non-zero relative humidity⁴⁸. The sodium ion in the CNC
clay composite might hence be surrounded by a hydration shell or might be localized in the
CNC regions of the sample, where hydroxyl groups of the cellulose could coordinate the
sodium ions and lead to similar ²³Na chemical shift as found for hydrated sodium.

To further characterize the sodium mobility in the composite materials, we determined the ²³Na longitudinal relaxation rates R₁ as a function of temperature (Figure 3c). At first, we notice increased relaxation rates with increasing temperature. A maximum of R₁ is expected for the case that the rate of molecular mobility equals the Larmor frequency, hence the observed R₁ reflects a molecular mobility below the maximum, and increasing mobility by rising the temperature leads to higher R₁. A higher relaxation rate R₁ can hence be interpreted as a larger molecular mobility, if the chemical environment of sodium is the same.

The Montmorillonite-containing samples (Figure 3c, squares) possess higher relaxation rates
than the Laponite samples (Figure 3c, circles). Assuming that the chemical environment in
smectite derivatives is fairly similar in both cases, the higher rate for the CNC-MMT
indicates a higher sodium mobility in CNC-MMT than in CNC-Lap. Furthermore, in both
cases, the R₁ values increase with a lower amount of CNC, which could indicate that adding
CNC lowers sodium mobility. This agrees well with the activation energy calculated in the
Arrhenius Plot (vide infra).

To evaluate the directionality of the sodium motion, we determined the relaxation rates of
two samples with low and high Montmorillonite content (5C2M and 2C4M) placed in
different orientations in the magnetic field (Figure 3d). In both cases, a higher relaxation rate
was found when the sample was placed with the normal of the material films along the
magnetic field (0°). This difference in relaxation rates hints at different sodium mobility in
directions perpendicular and within the planes of the material. As field fluctuations
perpendicular to B⁰ are responsible for longitudinal relaxation, the higher R₁ value for the 0°
orientation can be interpreted as a higher mobility in the plane compared to perpendicular to
the planes.
<figure><img src="image_11.png" /><figcaption>**Figure 4:** Solid-state <sup>7</sup>Li NMR of Laponite CNC films acquired under static conditions (a), and under magic angle spinning (b), with different compositions of CNC and clay mineral as indicated. The static <sup>7</sup>Li-NMR spectra were acquired at 116.6 MHz (corresponding to a <sup>1</sup>H Larmor frequency of 300.1 MHz) at 30°C, with samples consisting of strips of film material placed with the plane along the magnetic field. The magic angle spinning <sup>7</sup>Li NMR spectra were acquired at 194.4 MHz (corresponding to a <sup>1</sup>H Larmor frequency of 500.1 MHz) at approximately 40°C using a spinning frequency of 25kHz, on grained flakes of the film material.</figcaption></figure>

In addition to sodium, we also characterized the mobility of lithium in the CNC-Lap samples using solid-state ⁷Li NMR (Figure 4), performed under static conditions (Figure 4a) using 25 kHz magic angle spinning (Figure 4b). A significant line broadening is observed in the static samples (Figure 4a), which is removed to a large extent under MAS (Figure 4b). MAS effectively removes anisotropic contributions from chemical shift and dipolar interactions, while second-order quadrupolar interactions are partially retained. The efficient line narrowing by MAS indicates the presence of chemical shift anisotropies and/or dipolar couplings, which are not averaged out by molecular or atomic mobility. Therefore, the lithium seems to be immobilized and forms a part of the mineral layers. However, with increasing the CNC amount, a slight shift of the ⁷Li NMR line position can be noticed, which could indicate a change in the local crystal structure or electronic structure, induced by the CNC.
<figure><img src="image_12.png" /><figcaption>**Figure 5:** Solid-state CP/MAS <sup>13</sup>C-NMR spectra of Laponite CNC films (a), and montmorillonite CNC films (b), with different compositions of CNC and clay material as indicated. For comparison, neat CNC material was measured as well (bottom spectra). The NMR spectra were recorded at a Larmor frequency of 125.8 MHz (corresponding to <sup>1</sup>H at 500.1 MHz) under magic angle spinning of 15 kHz using cross polarization from <sup>1</sup>H and <sup>1</sup>H-decoupling of 50 kHz. The assignment of the signals to the cellulose carbons C1 to C6 is indicated.</figcaption></figure>

To characterize the potential influence of the clay component on the CNC structure, we also acquired solid-state <sup>13</sup>C NMR spectra (Figure 5). Using magic angle spinning of 15 kHz, individual cellulose signals of the C1, C4, and C6 carbons could be resolved (see assignments in Figure 5). The signals of C2, C3, and C5 overlap and give rise to a line shape with two maxima at ~75 ppm. For comparison, we acquired a spectrum of neat CNC (Figure 5, bottom), which resembles the <sup>13</sup>C NMR spectra of the composite materials. Except for a change in the line shape of the signal at 75 ppm, no significant difference is noticed, neither between the spectra of the CNC-Lap samples (Figure 5a) and the CNC-MMT samples (Figure 5b) nor with varied CNC fractions. The signal at 75 ppm shows an intensity decrease of the upfield component and an increase of the downfield component with an increasing fraction of the clay component in the case of the CNC-MMT samples. These intensity changes might be due to changes in the linewidth of the contributions of C2, C3 and C5 to the line-shape, and could indicate an interaction of one or several carbons of this set with the inorganic component of the composite material.
<figure><img src="image_13.png" /><figcaption>Figure 6: (a) and (c) Amplitude; (b) and (d) frequency sweep of CNC/Lap binary and CNC/MMT binary colloidal mixtures respectively; flow curves of (e) CNC/Lap and (f) CNC/MMT binary colloidal mixtures. Solid and hollow symbols in (a), (b), (c), and (d) represent the G' and G'' respectively.</figcaption></figure>

A key advantage of dynamic networks is their ability to flow under appropriate conditions. Dynamic oscillatory shear tests are performed to investigate the viscoelastic properties exhibited by the samples in the Linear Visco Elastic (LVE) region. The evolution of G' (storage modulus) and G'' (loss modulus) as a function of shear strain is observed at a constant angular frequency of 10 rad/s and by varying the shear strain from 0.01 to 1000 %.
The amplitude sweep plot of the CNC/Lap composite system is shown in Figure 6a. In all the cases, G' dominates over G'' indicating the solid like behaviour. Even though G' values are higher for 1C:1L, the films made out of these gel samples are brittle and the ionic conductivity is less compared to 1C:2L, whose films are highly flexible and exhibit higher ionic conductivity. From the amplitude sweep plot, a constant strain amplitude is selected from the LVE region, and the evolution of G' and G'' is investigated as a function of the angular frequency (ω). The frequency sweep of CNC/Lap binary mixtures is plotted in Figure 6b where G' is observed to be one order of magnitude higher than that of G''. Further, G' is almost independent of angular frequency (ω), indicating that the samples exhibit a gel-like behaviour. The amplitude and frequency sweep plots of the CNC/MMT composite system are shown in Figures (6c and 6d) respectively. From both the graphs, it is evident that 5C:2M is the one with higher G' and the film made out of this sample is flexible and exhibits higher ionic conductivity. All the samples reveal a solid-like behaviour since G' dominates over G'' in all the cases. From the flow curves of CNC/Lap (Figure 6e) and CNC/MMT (Figure 6f) mixtures, it is clear that the samples exhibit a shear thinning behaviour as the viscosity decreases with an increase in shear rate suggesting entangled polymer chains.

The CNC-Lap composites exhibit solid-like behaviour with dynamic moduli (G' and G'') that are higher than those of CNC-MMT composites, confirming that CNC-Lap forms a more rigid network than CNC-MMT. The results from dynamic oscillatory shear tests align with these observations. CNC-Lap composites show stronger elastic behaviour than CNC-MMT composites, suggesting that CNC-Lap is a better candidate for applications requiring higher mechanical stability. These films also exhibit better dynamic moduli and flow nature than their pristine counterparts (Figure S4).

## 2.1 Morphological Characterisation:
<figure><img src="image_14.png" /><figcaption>Figure 7: Scanning Electron Microscopic (SEM) images of (a) 2.5C:1M and (c) 1C:2L (b) and (d) corresponding cross-sectional images respectively. High-resolution transmission electron microscopic (HR-TEM) images of (e) 1C:2L (f) 2.5C:1M, selected area electron diffraction (SAED) pattern of (g) 2.5C:1M, (h) 1C:2L. Two-dimensional X-ray scattering patterns of thin films oriented parallel to the X-ray beam. (i) 2CNC:4Laponite; (j) 2CNC:4MMT. The Miller indices related to clay particles are indicated in black while indices in red correspond to CNC. (k) The angular dependence of the 001 reflections of Laponite and montmorillonite.</figcaption></figure>

The distribution of clay within the CNC matrix and their interactions are studied using HR-TEM and SEM images of thin films of the composite polymer electrolyte (Figure 7 a-f). The image reveals a sheet-like morphology with a few layers. Figure 7 (b and d) shows a comparison of the cross-sectional SEM images of 2.5C:1M and 1C:2L. Both composites exhibit a layer-by-layer morphology, which is attributed to the 2D sheets of Laponite and MMT, with the 1D CNC intercalating between the layers of the 2D sheets. The sheets consequently display an average interlayer spacing of 10 Å. The HR-TEM images also show
a sheet-like morphology with 5-7 layers. The image in Figure 7(g-h) displays the selected area electron diffraction (SAED) patterns, which suggest diffuse ring patterns, indicating the amorphous nature of the composite polymer electrolytes.⁴⁹,⁵⁰ The polarised optical microscopic (POM) images (Figure S3) shown before and after heating the composite polymer electrolytes reveal morphological rearrangement of the sheets, enhancing the ionic conductivity at higher temperatures (90 °C).

We then investigated the film texture by X-ray scattering experiments with their normal perpendicular to the incident X-ray beam. In this case, the intensities are angularly modulated on the detector, showing preferential orientations of both cellulose nanocrystals and clay nanoplatelets, as illustrated in Figure 7(i and j) for films containing Laponite and MMT, respectively. The orientation of a clay platelet can be identified by the angle θ between its normal and the normal to the film plane. The Hermans order parameter $\langle P_2 \rangle = \frac{1}{2}(3\langle \cos^2 \beta \rangle - 1)$ is used to characterize the mean orientation between a film in which all clay platelets are perfectly parallel ($\langle P_2 \rangle = 1$) and randomly oriented platelets ($\langle P_2 \rangle = 0$). We have used the methodology presented by Dabat et al.,⁵¹ to calculate the Hermans parameters from the angular variation of the 001 Bragg peaks of Laponite or montmorillonite (Figure 7k). $\langle P_2 \rangle$ values for the different films studied are reported in Table 1. The Hermans order parameter varies between 0.73 and 0.86, indicating very good orientation of the clay particles in the films. In addition, Laponite clay particles seem to be better oriented than montmorillonite, regardless of the sample composition. However, we do not notice any tangible effect of the cellulose content on improving the mean orientation of clay platelets.

<table><caption>**Table 1: Value of the Hermans order parameter $\langle P_2 \rangle$ for the different CNC/clay composite films.**</caption><thead><tr><th>Sample composition</th><th>P<sub>2</sub> (±0.01)</th></tr></thead><tbody><tr><td>2C:3M</td><td>0.76</td></tr><tr><td>2C:4M</td><td>0.81</td></tr><tr><td>3C:2M</td><td>0.73</td></tr><tr><td>3C:3M</td><td>0.74</td></tr><tr><td>4C:2M</td><td>0.80</td></tr><tr><td>2C:3L</td><td>0.84</td></tr><tr><td>2C:4L</td><td>0.86</td></tr></tbody></table>
3C:2L 0.80
3C:3L 0.80
4C:2L 0.81

<figure><img src="image_15.png" /><figcaption>Figure 8: Thermogravimetric analysis (TGA) of the composite polymer electrolyte (a) CNC-Lap and (b) CNC-MMT films.</figcaption></figure>

The thermal stability of composite polymer electrolytes CNC-Lap and CNC-MMT is investigated using TGA, revealing that both thin films remain relatively stable up to 300 °C (Figure 8). The Laponite curve exhibits two thermal events: the first, a release of water molecules (6.7% mass loss) between 30 and 150 °C, and the second, a collapse of the layer structure due to de-hydroxylation between 250 and 700 °C.⁵² For MMT, the initial degradation occurs at a similar temperature due to the loss of lattice water outside the coordination spheres. The second weight loss, occurring between 100 °C and 200 °C, is due to the release of hydrogen-bonded and coordinated water molecules, leading to a collapsed structure. The third degradation phase, between 200 °C and 250 °C, is attributed to the de-hydroxylation of structural water within the clay.⁵³ The degradation of both composite electrolytes around 300 °C is primarily due to the decomposition of surface carboxylated cellulose chains. The interaction between CNC, MMT and Laponite enhances the thermal stability of the nanocomposite thin films, making them suitable as solid-state electrolytes in sodium solid-state batteries.

## 2.2 Electrochemical characterisation:
<figure><img src="image_16.png" /><figcaption>Figure 9: Nyquist plots for (a) CNC-MMT, (b) CNC-Lap; Arrhenius plots (293-363 K) for (c) CNC-MMT, (d) CNC-Lap.</figcaption></figure>

Electrochemical impedance measurement is used to investigate the variation of conductivity in order to gain insights into ionic transport in composite electrolytes. Accordingly, Figures 9(a and b) display the Nyquist plot for CNC-MMT and CNC-Lap composite polymer electrolytes respectively. At higher frequencies, semicircular arcs are observed, which can be attributed to the bulk electrical properties of the material. However, at low frequencies, a linear regime indicative of diffusion-limited Warburg impedances is noticed, and more importantly, this frequency range is strongly influenced by the temperature and the composition⁵⁴. The presence of a single semicircle is characteristic of a single time constant, perhaps due to a geometrical capacitance in parallel with one or more resistive components. Interestingly, conductivity shows cross-over behaviour with temperature and exhibits an Arrhenius temperature dependence, as shown in Figures 9(c and d). This behaviour is possibly due to the swelling or reorganization of the clay (phase change) at different temperatures. The Arrhenius equation describes ionic conductivity.

$$ \sigma = A e^{-\frac{E_a}{RT}} \qquad (1) $$
where σ is total ionic conductivity, A is Arrhenius constant, E_a is the activation energy, R is gas constant and T is the temperature in Kelvin. Accordingly, Figure 9a, illustrates that the conductivity of CNC-Lap composites increases with clay concentration, peaking at a 1C:2L ratio with a conductivity of 4.7 × 10⁻⁵ S cm⁻¹. For CNC-MMT (Figure 9b), conductivity increases with CNC concentration, reaching a maximum at 2.5C:1M ratio with a conductivity of approximately 9.2 × 10⁻⁵ S cm⁻¹ at 25°C. Further, these composites exhibit conductivities of 1.08 × 10⁻⁴ S cm⁻¹ and 1.42 × 10⁻⁴ S cm⁻¹ respectively at 90°C, despite the phase change from room temperature to high temperature resulting in a bi-linear behaviour. Several others have also reported this type of variation in similar polymer composites. For instance, a similar Arrhenius behaviour is observed in PEO-LS solid polymer electrolytes by Li et al,⁵⁵ Laponite particles grafted with imidazolium ionic liquid on a PEO matrix by Jin et al,⁵⁶ and cellulose nanofibrils aerogel infiltrated with PEO and bis(trifluoromethane)sulfonimide lithium salt (LiTFSI) prepared by Qin et al.⁵⁷ More interestingly, the conductivity of our films is retained even after heating (90 °C) and subsequent cooling to 25 °C (Figure S5) suggesting no major reorganisation of the domains. From these results, it is evident that the impedance is strongly dependent on the composition and temperature as indicated by the ionic conductivity and activation energy (E_a) variation in Table 2.

<figure><img src="image_17.png" /></figure>

<figure><img src="image_18.png" /></figure>

<figure><img src="image_19.png" /></figure>

<figure><img src="image_20.png" /></figure>
**Figure 10: Storage modulus vs. concentration for (a) CNC-MMT, (b) CNC-Lap;**
**Conductivity vs. concentration of CNC plots for (c) CNC-MMT (d) CNC-Lap;**
**composite polymer electrolytes with different weight ratios of CNC-MMT/Lap.**

Dynamic percolation theory describes the ionic motion within polymeric chains of composite polymer electrolytes, focusing on the interactions between clays and cellulose nanocrystals (CNC), as well as their mechanical properties.⁵⁸,⁵⁹ Composite solid electrolytes exhibit two percolation thresholds: (1) the long-range connectivity of inorganic particles (contact mode) and (2) the long-range network formation of an interfacial layer. The interfacial layer represents the region where the polymer directly interacts with the inorganic particle (e.g., clay). This region exhibits unique properties influenced by chain orientation and confinement effects, distinguishing it from the bulk polymer and the bulk inorganic particle. Additionally, the interfacial region is believed to extend beyond the diameter of the inorganic particle.⁶⁰,⁶¹ Contact percolation occurs at high loadings of 33 vol% (clay: polymer) and interfacial percolation can occur even at 4 vol% (clay: polymer).⁶²,⁶³ Previous studies suggest lithium-ion transport in a composite is intrinsically linked to polymer-inorganic interactions and confinement effects.⁶⁴⁻⁶⁶ Consequently, the clay content significantly influences transport properties in composite polymer electrolytes (Figure 10). More specifically, since CNC has a negatively charged surface, while the clay nanosheets have a negatively charged surface and positively charged edges and Na⁺ ions in their interlayer spacing (single-ion conductor),⁶⁷⁻⁶⁹ which influences the electrostatic interactions between Lap/MMT and CNC surfaces and hydrogen bonding between the hydroxyl groups on their surfaces should promote the formation of complex self-assembled structures. Thus, with an increase in clay concentration from (1.5:1) weight ratio for Laponite and vice-versa for MMT, there is a subsequent increase in conductivity and elasticity of these films (Figure 10c-d). Adding clay to CNC increases the electrostatic interactions and decreases the crystallinity of the CNC matrix. The sodium ions interact with the hydroxyl groups of the CNC, as observed in ss-NMR and XPS. Higher CNC concentrations minimise conductivity in CNC-Lap, whereas increases in CNC-MMT composites. Such behaviour stems from the smaller particle size of Laponite (an order of magnitude smaller than MMT), which facilitates better networking with rod-shaped CNCs and provides more positively charged edge sites to interact with the negatively charged CNCs. Based on wettability measurements, the contact angle between the composite polymer electrolyte and water indicates enhanced hydrophilicity. This suggests that the surface properties of the composite are well-suited for energy applications requiring improved
interaction with CNC environments, potentially aiding in efficient ion transport and compatibility. The difference in particle size influences conductivity, viscosity, morphology, and diffusivity in these systems (Figure S6 and Table S3).

## 2.3 Ionic diffusivity and conductivity estimated using molecular dynamics simulations

We performed all-atom molecular dynamics (MD) simulations to investigate ionic diffusivity within the CNC-MMT composite (see Materials and Methods section in the Supplementary Information for details) at a weight ratio of 2.5:1 (CNC:MMT), which exhibited the highest conductivity among the composites studied across the experimental temperature range. A representative snapshot of the system is shown in Fig. 11(a). The Na⁺ self-diffusion constant (D) was computed from the slope of the Na⁺ ions' mean mean-square displacement (MSD) vs time (t) plot as, $D = \frac{1}{6} \frac{d(\text{MSD})}{dt}$. The ionic conductivity (σ) was then estimated using the Nernst-Einstein equation:⁷⁰ $\sigma = \frac{e^2}{Vk_BT} (Z_-^2 N_- D_- + Z_+^2 N_+ D_+)$, where e is the electronic charge, V is the system volume, and T is the temperature (in K), k<sub>B</sub> is the Boltzmann constant. Here, Z<sub>-</sub> and Z<sub>+</sub> are the valencies, N<sub>-</sub> and N<sub>+</sub> are the number of negative and positive ions, respectively, and D<sub>-</sub> and D<sub>+</sub> are their corresponding diffusion constants. In the CNC-MMT composite, the MMT surface carries a negative charge, and since MMT segments remain nearly frozen on the simulation timescale of 500 ns, we set the D<sub>-</sub> to zero.

<figure><img src="image_21.png" /><figcaption>**Figure 11:** (a) A representative snapshot of the randomly packed CNC-MMT composite. The variation of the Na⁺ self-diffusion constant *D* (b) and ionic conductivity *σ* (c) with inverse temperature.</figcaption></figure>

Figures 11(b) and 11(c) show the Na⁺ diffusivity and conductivity, respectively, as functions of inverse temperature. Both *D* and *σ* increase nearly monotonically with temperature, although a slight decrease in these transport coefficients is observed near 340 K. However,
given the error bars from independent simulations, we could not conclusively establish this trend, and warrants further investigation. Furthermore, we observe a crossover in the temperature dependence of both *D* and *σ* near 340 K. The estimated ionic conductivity from the Nernst-Einstein equation (see Fig. 11(c)) differs from experimental values by approximately an order of magnitude. Similar discrepancies have been reported in computational studies of lithium-ion conductivity in solid electrolytes.⁷¹,⁷²

This deviation may arise due to: (i) an incomplete representation of the composite's molecular-scale structure, as we modeled it as a randomly packed CNC-MMT matrix with the desired composition, which may not fully replicate the experimental environment, or (ii) the choice of force-field parameters adopted for modeling of the CNC-MMT composite. However, despite these plausible limitations, within the framework of our *in silico* modelling, we were able to capture the overall temperature-dependent trend of *σ*.

<table><caption>**Table 2: Kinetic properties of the CNC-Lap and CNC-MMT composite polymer electrolyte.**</caption><thead><tr><th>Composition</th><th>R<sub>s</sub> (Ω)</th><th>R<sub>ct</sub> (Ω)</th><th>σ (10<sup>-5</sup> S cm<sup>-1</sup>) @ 25 °C</th><th>σ (S cm<sup>-1</sup>) @ 90 °C</th><th>E<sub>a</sub> (eV)</th></tr></thead><tbody><tr><td>1C:2L</td><td>43</td><td>64.7</td><td>4.7</td><td>1.2 x 10<sup>-4</sup></td><td>0.14</td></tr><tr><td>1C:1.5L</td><td>45</td><td>64.5</td><td>3.4</td><td>1.4 x 10<sup>-4</sup></td><td>0.16</td></tr><tr><td>1C:1L</td><td>180.1</td><td>244.6</td><td>2.5</td><td>2.6 x 10<sup>-5</sup></td><td>0.15</td></tr><tr><td>1.5C:1L</td><td>47.7</td><td>70.7</td><td>1.7</td><td>7.14 x 10<sup>-5</sup></td><td>0.15</td></tr><tr><td>2C:1L</td><td>118.3</td><td>155.5</td><td>1.1</td><td>4.4 x 10<sup>-5</sup></td><td>0.14</td></tr><tr><td>1C:1M</td><td>59</td><td>316</td><td>7.63</td><td>1.1 x 10<sup>-4</sup></td><td>0.08</td></tr><tr><td>1C:1.5M</td><td>47.8</td><td>382</td><td>7.5</td><td>1.1 x 10<sup>-4</sup></td><td>0.045</td></tr><tr><td>1C:2M</td><td>108</td><td>680</td><td>7.6</td><td>1.05 x 10<sup>-4</sup></td><td>0.07</td></tr><tr><td>1.5C:1M</td><td>69.9</td><td>259</td><td>8.3</td><td>1.02 x 10<sup>-4</sup></td><td>0.05</td></tr><tr><td>2C:1M</td><td>21.4</td><td>98</td><td>8.4</td><td>1.1 x 10<sup>-4</sup></td><td>0.06</td></tr><tr><td>2.5C:1M</td><td>19.6</td><td>59.6</td><td>9.2</td><td>1.4 x 10<sup>-4</sup></td><td>0.04</td></tr></tbody></table>

## 2. Conclusions:
This study presents a novel approach for developing clay-based solid electrolytes with enhanced ionic conductivity and mechanical stability. ATR-FTIR, XRD, XPS, SAXS, SS-NMR, and Raman analysis prove that CNCs in the clay interlayer improve the ionic conductivity, structural integrity, and safety thereby preventing electrical shortages. After solvent evaporation from the hydrogel, the resulting transparent thin films exhibit smooth surfaces, indicating homogeneity and no significant aggregation within the membrane, which enhances interfacial interactions between CNC and clay. Molecular dynamics simulations align with the experimental results further corroborating the enhanced Na⁺ transport to facilitate the fabrication of all solid-state batteries. The composite polymer electrolytes are freestanding, flexible, and can bend without breaking, demonstrating good mechanical properties for applications in flexible electronics, addressing a significant challenge for inorganic solid electrolytes, with implications for creating safer, more efficient, and more sustainable solid-state batteries. These findings contribute to the broader goal of advancing clean energy technologies and moving towards a carbon-zero footprint future.

**Associated Contents:** Experimental procedures, supporting figures and tables are present in Supplementary Information.

**Conflict of Interest:** The authors declare no competing financial interest.

**Acknowledgements:**

The authors thank Antoine Boury for his help during the SEM acquisition. H. K., P. L., and E. P. acknowledge the financial support of the French ANR agency (ANR) under grant agreement ANR-20-CE05-0028 (OCCAMH2). S. L. G. acknowledges financial support from the Helmholtz Association (NACIP program) and the instrument grant INST 121384/58-1 from Deutsche Forschungsgemeinschaft (DFG). R. S. S. acknowledges financial support from DST-SERB (Grant No. SRG/2020/001415 and CRG/2023/002975). Y. S. thanks to the computing facility at the IISER Tirupati and PARAM Brahma at IISER Pune. R. K. P. thanks the CNRS for financing his research visit to the LPS. R.K.P. acknowledges the Department of Science and Technology for an INSPIRE Faculty Award Grant [DST/INSPIRE/04/2016/002370] and the Core Research Grants [CRG/2020/006281, CRG/2021/004759] from SERB, Government of India. V. K. P. would like to thank SERB for the J. C. Bose Fellowship (JCB/2020/000018). S. M. thanks SAIF, IIT Bombay for the HR-TEM measurements, and Shri M. N. Singh for his assistance in the temperature-dependent XRD measurements at BL-12 (Indus-2), RRCAT, Indore.
References

1. Solid-State Battery Roadmap 2035+, *Fraunhofer ISI*, (2022).

2. Trahey, L., Brushett, F. R., Balsara, N. P., Ceder, G., Cheng, L., Chiang, Y.-M., Hahn, N. T., Ingram, B. J., Minteer, S. D., Moore, J. S., Mueller, K. T., Nazar, L.F., Persson, K. A., Siegel, D. J., Xu, K., Zavadil, K. R., Srinivasan, V. and Crabtree, G. W. Energy storage emerging: A perspective from the Joint Center for Energy Storage Research. *PNAS*, 117, 23; 12550-12557 (2020).

3. Renewables 2024, Analysis and Forecast to 2030, *IEA*, 2024.

4. Stenclik, D., Denholm, P., Chalamala, B. R. The role of energy storage for renewable integration. *IEEE Power and Energy*, 15, 31–39 (2017).

5. Denholm, P., Jorgenson, J., Hummon, M., Jenkin, T., and Palchak, D. The Value of Energy Storage for Grid Applications (National Renewable Energy Laboratory, 2013).

6. Janek, J., Zeier, W.G. Challenges in speeding up solid-state battery development. *Nat Energy*, 8; 230–240 (2023).

7. Challenges for and Pathways toward Li-Metal-Based All-Solid-State Batteries. *ACS Energy Lett.*, 6, 4, 1399–1404 (2021).

8. Janek, J., Zeier, W. A solid future for battery development. *Nat Energy*, 1, 16141 (2016).

9. Meng, Y. S., Venkat, S., and Xu, K. Designing better electrolytes. *Science*, 378, 6624; (2022).

10. Brandell, D., Mindemark, J., Hernández, G. Polymer-Based Solid-State Batteries. https://doi.org/10.1515/9781501521140-202.

11. Anton Van der Ven. Solid electrolytes redefine ion conduction. *Science*, 382, 6670; 513-514 (2023).

12. Balaish, M., Gonzalez-Rosillo, J.C., Kim, K.J. Zhu, Y., Hood, Z. D., and Rupp, J. L. M. Processing thin but robust electrolytes for solid-state batteries. *Nat Energy*, 6, 227–239 (2021).

13. Shannon, R. D., Taylor, B. E., English, A. D., & Berzins, T. New Li solid electrolytes. *Electrochimica Acta*, 22, 7; 783–796 (1977).

14. Bates, J. B., Dudney, N. J., Gruzalski, G. R., Zuhr, R. A., Choudhury, A., Luck, C. F., & Robertson, J. D. Fabrication and characterization of amorphous lithium electrolyte thin films and rechargeable thin-film batteries. *Journal of Power Sources*, 43, 1-3; 103–110 (1993).
15. Kato, Y., Hori, S., Saito, T., Suzuki K, Hirayama, M., Mitsui, A., Yonemura, M., Iba, H. and Kanno, R. High-power all-solid-state batteries using sulfide superionic conductors. *Nat Energy* 1, 16030 (2016).

16. Murugan, R., Thangadurai, V. and Weppner, W. Fast lithium ion conduction in garnet-type $Li_7La_3Zr_2O_{12}$. *Angew. Chem. Int. Ed.*, 46, 41; 7778–7781 (2007).

17. Ye, L., Gil-González, E., & Li, X. $Li_9.54Si_{1.74}(P_1-xSb)x_1.44S_{11.7Cl}O_{3.0.3}$: A functionally stable sulfide solid electrolyte in air for solid-state batteries. *Electrochemistry Communications*, 128, 107058 (2021).

18. Kanno, R., & Murayama, M. Lithium Ionic Conductor Thio-LISICON: The $Li[sub 2]S-GeS[sub 2]-P[sub 2]S[sub 5]$ System. *Journal of The Electrochemical Society*, 148, 7; A742 (2001).

19. Nazar, L., Adeli, P., Bazak, J. D., Park, K.-H., Kochetkov, I., Huq, A., & Goward, G. Boosting solid-state diffusivity and conductivity in lithium superionic argyrodites by halide substitution. *Angew. Chem. Int. Ed.*, 58, 8681–8686 (2019).

20. Li, H., Lv, T., Sun, H., Qian, G., Li, N., Yao, Y., Chen, T. Ultrastretchable and superior healable supercapacitors based on a double cross-linked hydrogel electrolyte. *Nat Commun* **10**, 536 (2019).

21. Zhao, J.; Chen, D.; Boateng, Bismark; Zeng, Guangfeng; Han, Yupei; Zhen, C.; Goodenough, J. B.; Weidong, H. Atomic interlamellar ion path in polymeric separator enables long-life and dendrite-free anode in lithium ion batteries. *Journal of Power Sources*, 451, 227773 (2020).

22. Lan, Y., Chen, D. Fabrication of nano-sized attapulgite-based aerogels as anode material for lithium ion batteries. *J. Mater Sci.*, **53**, 2054–2064 (2018).

23. Costenaro, D., Bisioa, C., Carniato, F., Gattia, G., Oswald, F., Meyer, T., B., Marchese, L. Size effect of synthetic saponite-clay in quasi-solid electrolyte for dye-sensitized solar cells. *Solar Energy Materials & Solar Cells*, 117, 9–14 (2013).

24. Bryant, S. J., Calabrese, V., da Silva M. A., Hossain, K. M. Z., Scott, J. L., Edler, K. J. Rheological modification of partially oxidised cellulose nanofibril gels with inorganic clays. *PLoS ONE* 16(7): e0252660 (2021).

25. Wu, W., Dong, Z., He, J., Yu, J., Zhang J. Transparent cellulose/Laponite nanocomposite films. *J Mater Sci.*, 51:4125–4133 (2016).

26. Colomban, P. Analysis of strain and stress in ceramic, polymer and metal matrix composites by Raman spectroscopy. *Adv. Eng. Mater.* 4, 535–542 (2002).
27. Paillet, M., Parret, R., Sauvajola, J.-L., and Colomban P. Graphene and related 2D materials: An overview of the Raman studies. J. Raman Spectrosc., 49, 8-12 (2018).

28. Wang, A., Freeman J. J., and Jolliff B. L. Understanding the Raman spectral features of phyllosilicates. *J. Raman Spectrosc.*, 46, 829-845 (2015).

29. Frost, R. L., and Rintoul L. Lattice vibrations of montmorillonite: an FT Raman and X-ray diffraction study. *Applied Clay Science*, 11; 171 – 183 (1996).

30. Wang, Y.-Q., Xu, H., Cao, B., Ma, J. and Yu, Z.-W. In-Situ Species Analysis of a Lithium-Ion Battery Electrolyte Containing LiTFSI and Propylene Carbonate. *J. Phys. Chem. Lett.*, 15, 5047-5055 (2024).

31. Pujala R. K., and Bohidar H. B. Slow dynamics, hydration and heterogeneity in Laponite dispersions. *Soft Matter*, 9, 2003-2010 (2013).

32. Iturrioz-Rodriguez, N., Martin-Rodriguez, R., Renero-Lecuna, C., Aguado, F., Gonzalez-Legarreta, L., Gonzalez, J., Fanarraga, M. L.; Perdigón, A. C. Free-labeled nanoclay intracellular uptake tracking by confocal Raman imaging. *Applied Surface Science*, 537, 1478-1480 (2021).

33. Farmer, V. C., Russell, J. D. The infra-red spectra of layer silicates. *Spectrochimica Acta*, 20, 7; 1149-1173 (1964).

34. Müller, C., Pejcic, B., Esteban, L., Piane, C. D., Raven, M., and Mizaikoff, B. Infrared Attenuated Total Reflectance Spectroscopy: An Innovative Strategy for Analyzing Mineral Components in Energy Relevant Systems. *Sci Rep.*, 4, 6764 (2014).

35. Srasra, E., Bergaya, F. and Fripiat, J. J. Infrared spectroscopy study of tetrahedral and octahedral substitutions in an interstratified illite-smectite clay. *Clays Clay Miner.*, 42, 237-241 (1994).

36. Madejova, J. FTIR techniques in clay mineral studies. *Vib. Spectrosc.*, 31, 1-10 (2003).

37. Wang, Y.-Q., Xu, H., Cao, B., Ma, J., and Yu, Z.-W. In Situ Species Analysis of a Lithium-Ion Battery Electrolyte Containing LiTFSI and Propylene Carbonate. *J. Phys. Chem. Lett.*, 15, 5047-5055 (2024).

38. Belgacem, M. N., Czeremuszkin, G., Sapieha, S., Gandini, A. Surface characterization of cellulose fibres by XPS and inverse gas chromatography. *Cellulose*, 2, 3; 145-157 (1995).

39. Khili, F., Borges, J., Almeida, P. L., Boukherroub, R., and Omrani, A. D. Extraction of Cellulose Nanocrystals with Structure I and II and Their Applications for Reduction
of Graphene Oxide and Nanocomposite Elaboration. *Waste Biomass Valor*, 10, 1913–1927 (2019).

40. Raza, M., Abu-Jdayil, B., Banat F. and Al-Marzouqi, A. H. Isolation and Characterization of Cellulose Nanocrystals from Date Palm Waste. *ACS Omega*, 7, 29, 25366–25379 (2022).

41. Grosvenor, A. P., Kobe, B. A., Biesinger, M. C. and McIntyre N. S. Investigation of multiplet splitting of Fe 2p XPS spectra and bonding in iron compounds. *Surf. Interface Anal.*, 36; 1564–1574 (2004).

42. Barr T. L., Seal, S., He, H., Klinowski, J. X-ray photoelectron spectroscopic studies of kaolinite and montmorillonite. *Vacuum*, 46, 12; 1391–1395 (1995).

43. Andrello, C., Gouder, T., Favergeon, L., Desgranges, L., Tereshina-Chitrova, E., Havela, L., Konings, R. J. M., and Eloirdi, R. In-situ high resolution photoelectron spectroscopy study on interaction of sodium with UO₂+X film (0 ≤ x ≤ 1). *Journal of Nuclear Materials*, 545, 152646 (2021).

44. Seyama H. and Soma, M. X-Ray Photoelectron Spectroscopic Study of the Effect of Heating on Montmorillonite Containing Sodium and Potassium Cations. *Clays and Clay Minerals*, 34, 6; 672 - 676 (1986).

45. Pollak, R. A., Ley, L., McFeely, F. R., Kowalczyk, S. P., and Shirley, D. A. Characteristic energy loss structure of solids from x-ray photoemission spectra. *Journal of Electron Spectroscopy and Related Phenomena*, 3(5), 381–398 (1974).

46. Xu, P., Erdem, T., and Eiser, E. A simple approach to prepare self-assembled, nacre-inspired clay/polymer nanocomposites. *Soft Matter*, 16, 5497-5505 (2020).

47. Ohashi, R.; Michal, C. A.; Hamad, W. Y.; Nguyen, T.-D.; Mizuno, M.; MacLachlan, M. J. Solid-state 23Na NMR spectroscopy studies of ordered and disordered cellulose nanocrystal films. *Solid State Nuclear Magnetic Resonance*, 97, 31-39 (2019).

48. Nanda, R., Bowers, G. M., Loganathan, N., Burton, S. D., and Kirkpatrick, R. J. Temperature dependent structure and dynamics in smectite interlayers: $^{23}$Na MAS NMR spectroscopy of Na-hectorite. *RSC Adv.*, 9, 12755-12765 (2019).

49. Alhassan, S. M., Qutubuddin, S. and Schiraldi, D. A. Graphene Arrested in Laponite–Water Colloidal Glass. *Langmuir*, 28 (8), 4009-4015 (2012).

50. Beermann, T., and Brockamp, O. Structure analysis of montmorillonite crystallites by convergent-beam electron diffraction. *Clay Minerals*, 40(01), 1–13 (2005).
51. Dabat, T., Hubert, F., Paineau, E., Launois, P., Laforest, C., Grégoire, B., Dazas, B., Tertre, E., Delville, A., and Ferrage, E. A general orientation distribution function for clay-rich media. *Nature Communications*, 10, 1; 5456 (2019).

52. Silva, J. M., Barud, H. S., Meneguin, A. B., Constantino, V. R. L., and Ribeiro, S. J. L. Inorganic-organic bio-nanocomposite films based on Laponite and Cellulose Nanofibers (CNF). *Applied Clay Science*, 168; 428-43 (2019).

53. Bastow, T., Hardin, S. G., and Turney, T. W. The formation of β'-sialon from a montmorillonite-polyacrylonitrile composite by carbothermal reduction: an NMR, TGA, XRD and EM study. *Journal of Materials Science*, 26(6), 1443-1453 (1991).

54. Vivier, V. and Orazem, M. E. Impedance Analysis of Electrochemical Systems, *Chem. Rev.*, 122, 11131-11168 (2022).

55. Li, J., Li, F., Li, D., Cheng, D., Wang, Z., Liu, X., Wang, H., Zeng, X., Huang, Y. and Xu, H. Negatively Charged Laponite Sheets Enhanced Solid Polymer Electrolytes for Long-Cycling Lithium-Metal Batteries. *ACS Appl. Mater. Interfaces*, 15, 4044-405 (2023).

56. Jin, B., Wang, D., He, Y., Mao, J., Kang, Y., Wan, C., Xia, W., Kim, J., and Yamauchi, Y. Composite polymer electrolytes with ionic liquid grafted-Laponite for dendrite-free all-solid-state lithium metal batteries. *Chem. Sci.*, 14, 7956 (2023).

57. Qin, H., Fu, K., Zhang, Y., Ye, Y., Song, M., Kuang, Y., Jang, S.-H., Jiang, F., and Cui, L. Flexible nanocellulose enhanced $Li^+$ conducting membrane for solid polymer electrolyte, *Energy Storage Materials*, 28, 293-299 (2020).

58. Nitzan A., Granek R., and Ratner M. Mechanical properties of dynamically disordered networks. *Journal of Non-Crystalline Solids*, 131-133, 1018-1021 (1991).

59. Dixit, M. B., Zaman, W., Hortance, N., Vujic, S., Harkey, B., Shen, F., Tsai, W.-Y., Andrade, V. D., Chen, X. C., Balke, N. and Hatzell, K. B. Nanoscale Mapping of Extrinsic Interfaces in Hybrid Solid Electrolytes. *Joule*, 4, 207-221, (2020).

60. Hamming, L. M., Qiao, R., Messersmith, P. B. and Brinson, L. C. Effects of Dispersion and Interfacial Modification on the Macroscale Properties of $TiO_2$ Polymer Matrix Nanocomposites. *Composites science and technology*, 69, 1880-1886 (2009).

61. Sebastian, M. T. and Jantunen, H. The Effects of $TiO_2$ Particle Size on the Properties of PTFE/$TiO_2$ Composites. *International Journal of Applied Ceramic Technology*, 7, 415-434 (2010).
62. Zamana, W., Hortance, N., Dixit, M. B., Andrade, V. D., and Hatzell, K.B. Visualizing percolation and ion transport in hybrid solid electrolytes for Li metal batteries. *J. Mater. Chem. A*, **7**, 23914-23921 (2019).

63. Dieterich, W., Dürr, O., Pendzig, P., Bunde, A., Nitzan, A. Percolation concepts in solid state ionics. *Physica A*, **266**, 229-237 (1999).

64. Winter M., The solid electrolyte interphase—the most important and the least understood solid electrolyte in rechargeable Li batteries. *Z. Phys. Chem.*, **223**, 1395–1406 (2009).

65. Wang, L., Menakath, A., Han, F., Wang, Y., Zavalij, P. Y., Gaskell, K. J., Borodin, O., Iuga, D., Brown, S. P., Wang, C., Xu, K., Eichhorn B. W. Identifying the components of the solid–electrolyte interphase in Li-ion batteries. *Nat. Chem.* **11**, 789–796 (2019).

66. Zhu, Y., Hood, Z. D., Paik, H., Groszewicz, P. B., Emge, S. P., Sayed, F. N., Sun, C., Balaish, M., Ehre, D., Miara, L. J., Frenkel, A. I., Lubomirsky, I., Clare, P. G., and Rupp, J. L. M. Highly disordered amorphous Li-battery electrolytes. *Matter*, **7**, 2; 500-522 (2024).

67. Balnois, E.; Durand-Vidal, S.; Levitz, P. Probing the Morphology of Laponite Clay Colloids by Atomic Force Microscopy. *Langmuir*, **19**, 6633–6637 (2003).

68. Saunders, J. M.; Goodwin, J. W.; Richardson, R. M.; Vincent, B. A Small-Angle X-Ray Scattering Study of the Structure of Aqueous Laponite Dispersions. *J. Phys. Chem. B*, **103**, 9211–9218 (1999).

69. Liu, D.; Chen, X.; Yue, Y.; Chen, M.; Wu, Q. Structure and Rheology of Nanocrystalline Cellulose. *Carbohydr. Polym.*, **84**, 316–322 (2011).

70. Mehrer, H. *Diffusion in solids: fundamentals, methods, materials, diffusion-controlled processes*, volume 155. Springer Science & Business Media, (2007).

71. Webb, M. A., Jung, Y., Pesko, D. M., Savoie, B. M., Yamamoto, U., Coates, G. W., Balsara, N. P., Wang, Z.-G., and Miller III, T. F. Systematic computational and experimental investigation of lithium-ion transport mechanisms in polyester-based polymer electrolytes. *ACS central science*, **1**(4):198–205, (2015).

72. Baktash, A., Reid, J. C., Roman, T. and Searles, D. J. Diffusion of lithium ions in lithium-argyrodite solid-state electrolytes. *npj Computational Materials*, **6**(1):162, (2020).