Interface-driven enhancement of Nafion thin-film conductivity via controlled aminothiol modification

Abstract

The performance of electrochemical devices is critically influenced by the interfacial confinement of ionomer chains at the catalyst interface. In sub-μm-thick ionomer films at the catalyst interface, confinement of water and ionomer chains limits proton conductivity, hindering electrochemical reactions at the cathode and compromising fuel cell performance. Here, we investigated the effect of interfacial engineering on Nafion thin-film conductivity by covalently modifying gold (Au) electrodes with cysteamine (Cys), an aminothiol linker. Nafion films (∼60 nm thick) spin-coated on Cys-modified Au exhibited up to 4-fold increase in in-plane proton conductivity (σ_IP) at 85% RH compared to films on bare Au, with insignificant changes in water uptake. Molecular dynamics (MD) simulations revealed that the Cys modification of Au induced a more surface-parallel orientation of the Nafion backbone near the substrate. This promoted highly ordered –SO₃H groups and in-plane proton-conducting pathways while slightly lifting Nafion chains away from the hard wall-like Au surface. Attaching to the –NH₂ terminal of Cys, but hanging slightly up from Au, rendered a milder confinement to Nafion chains and surrounding water molecules. Consequently, the in-plane self-diffusion coefficients of water and hydronium ions increased, and the storage modulus of the films decreased, supporting more water-ionomer mobility and a milder confinement. While the more surface-parallel nature of ionomer chains with lower confinement on Cys improved in-plane proton conductivity of Nafion films, the out-of-plane conductivity did not change significantly for the same reason. These findings demonstrated the importance of controlling interfacial chemistry and linker design to gain control over ionomer chain orientation, confinement, and directional ionic conductivity, offering a promising strategy to enhance ion transport in thin ionomer films for next-generation electrodes and electrochemical devices.

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Introduction

Electrochemical cells are at the core of clean energy technologies. Their performance, however, is often limited by ion and gas transport through the catalyst binder, an ultrathin ionomer layer that binds catalyst particles to electrode surfaces. In proton exchange membrane fuel cell (PEMFC) cathodes, for instance, sluggish ion transport at the ionomer–catalyst interface hinders the oxygen reduction reaction (ORR). Sub-μm-thick ionomer films like Nafion, when deposited on electrodes, exhibit proton conductivity 1–2 orders of magnitude lower than their bulk, free-standing membrane counterparts (typically tens of μm thick).^1–4 This reduced conductivity delays proton delivery to catalyst active sites, compromising ORR kinetics, power density, and overall device efficiency. This challenge underscores the need for focused design strategies targeting electrodes and ionomer-electrode interfaces. However, historically, research emphasized bulk membrane design where interfacial effects are minimal. Only recently has attention shifted to the behavior of ionomers under thin-film confinement, where ionomer-substrate interactions become dominant.

Geometric and interfacial confinements are primarily held responsible for the weak proton conductivity in thin ionomer films. As the film thickness approaches only a few times the ionomer's radius of gyration, polymer chain mobility^5,6 and diffusion⁷ become increasingly restricted, a phenomenon known as geometric confinement. Interfacial confinement, on the other hand, arises from strong interactions among water molecules, the ionomer's sulfonic acid groups (–SO₃H), and the substrate, further constraining chain dynamics and disrupting the nanoscale phase separation necessary for efficient proton conduction.^5,6,8–16

Perfluorosulfonic acid-based ionomer, Nafion, conducts protons through hydrated channels formed by –SO₃H-containing side chains and water molecules.^17–23 While –SO₃H and H₂O naturally interact to form these conductive domains, they also strongly interact with substrates in Nafion thin films.^{3,5,6,10–12,14–16,22,24–29} For example, a silica-based substrate with a native oxide layer can undergo interfacial interaction with water and –SO₃H groups of ionomers via H-bonding.^3,5,6,9 Experimental and computational efforts consistently indicate that as Nafion film thickness approaches the polymer's radius of gyration or the length of the longer axis of self-assembled chain bundles in dispersion, the Nafion backbones can be compelled to align parallel to the substrate.^25,30–32 Kreuer et al.²³ simulated planar polymer-water sheets that fit well with experimental GISAXS data, indicating near-surface-parallel chain orientation of Nafion in thin films. Birefringence measurements,^10,25 ionomer peak analysis as a function of azimuthal angle (2D GISAXS),³³ and neutron reflectometry (NR)^34–37 further confirmed surface-parallel alignment of Nafion chains in sub-micron thick films, especially when film thickness drops below 100 nm.

This surface-parallel backbone orientation immediately next to the substrate naturally forces the side chains of Nafion to extend perpendicularly toward the substrate and makes its –SO₃H groups interact more with the substrate. The water molecules accumulating next to the substrate^34,38 also participate in this interactive process, together causing ionomer chain pinning to the substrate. This observation extends beyond oxide substrates to metallic ones (Au, Pt), having close relevance in electrochemical studies and practical applications. Both Au and Pt are hydrophilic,^3,10,39 making them interactive with –SO₃H groups.^26,27,40 Birefringence studies,²⁵ indicated surface-parallel orientation of Nafion backbones on these metals, while DFT calculations predicted ordered sulfonate layers on Pt.⁴¹ Interactions between sulfonate groups and metal surfaces (both Au, Pt)^26,40 and ionomer chain adsorption⁴² have been reported via sum frequency generation spectroscopy,²⁶ surface enhanced Raman spectroscopy,⁴⁰ electrochemical QCM and AFM imaging.⁴²

If severe chain pinning happens near the substrate, the water-ionomer chains near the substrate should lose mobility, fail to spontaneously phase segregate, and show weak proton conduction near the substrate. This is exactly what we observed in depth-specific proton conduction profiles of Nafion films, obtained via confocal laser scanning microscopy (CLSM) imaging.² In Nafion films on glass substrates, up to half the film thickness near the substrate exhibits extremely weak proton conduction, which gradually improves toward the air interface.² Even under humidified conditions, when water accumulates near the substrate (as per NR^34,38), this near-substrate region remains poorly conductive. Importantly, the interfacial water-ionomer interaction was found to stiffen ionomer films,^5,6 and the widths of the stiffened and low-proton-conducting zones were similar in Nafion films.² These demonstrate that ionomer chain pinning at the ionomer-electrode interface significantly impairs interfacial proton conduction in Nafion thin films. Such adsorption of –SO₃H groups on the Pt surface not only weakens proton conduction but also blocks O₂ from reaching Pt catalyst active sites, impeding ORR under fuel cell operating conditions.⁴³ It is thus critical to optimize interfacial chemistry to alleviate interfacial chain pinning and boost thin-film proton conductivity and ORR activity.

By far, strategies to modify electrode-ionomer interfaces, include tailoring catalyst^37,44,45 and carbon support⁴⁶ structure and chemistry, altering substrate hydrophilicity^45,47–49 or zeta potential,⁴⁵ doping substrate with atoms (like N),^50–52 or incorporating/immobilizing molecules, such as ionic liquids (ILs),^49,53–56 IL-based block copolymers,^57–59 strategically designed ionomers for interconnected ion channels,^1,9,60,61 or engineered peptides^62–64 near electrodes. Structure and chemical makeup of catalysts^37,44,45 and carbon supports⁴⁶ have influenced catalyst activity, agglomeration, coverage, and stability, water-ionomer distribution, and proton and gas transport, thereby. IL-containing sulfonated block copolymers⁵⁷ or IL-modified catalyst pores⁶⁵ reduced Nafion coverage of Pt, which improved membrane electrode assembly (MEA)-level performance of PEMFCs. Certain ionomers were able to reduce chain crowding near the substrate interface, causing reduced chain pinning and improved proton conduction near the substrate interface as well as across the film.^60,61 Simulations also showed that, as opposed to staying in contact with Pt, if the –SO₃H groups are slightly distant from the Pt (at the edge of Helmholtz plane), it can alleviate the blocking effect on ORR.⁴³

While these findings strongly suggested weakening –SO₃H/–SO₃⁻ adsorption on substrate as a key to improving ion conductivity in confined ionomer films and ORR activity in catalyst layers, it is still not clear which interfacial physical and chemical makeup is optimal to achieve so. Our previous work¹⁰ attempted to address this gap by engineering SiO₂ substrates and Au electrodes with 3-aminopropyltriethoxysilane (APTES) and understanding its impact on thin film proton conductivity and other properties. APTES introduced a spiderweb-like barrier near the substrate, disrupting chain alignment and reducing surface pinning of Nafion. This was confirmed by decreased fluorine and sulfur signals near the substrate (via SEM-EDX), reduced film stiffening, and enhanced through-thickness proton conduction.¹⁰ CLSM imaging further showed that APTES narrowed down the low-conductivity region next to the substrate and improved the proton conduction across Nafion films.¹⁰ These findings suggested that a physical, small obstacle next to the substrate can go a long way by reducing chain pinning, chain confinement, and improving proton conduction behavior at the ionomer-electrode interface as well as across the film. Having said that, a big role will be played in these processes by the nature and length of the linker, but a systematic understanding of their influence has a long way to go.

To address these needs, this work explored the use of cysteamine (Cys) as a covalent linker on Au electrodes to engineer the ionomer-electrode interface (Fig. 1). Cys features a short ethyl side chain with terminal amino (–NH₂) and thiol (–SH) groups, enabling strong Au–thiol bonding and surface functionalization without harsh treatments. Au is widely used for electrochemical characterization, and since Pt, the practical electrode material, also binds thiols,^66–68 insights gained here on the role of Cys modification on thin-film ion conduction behavior can translate to real device systems. Compared to APTES, forming a relatively thick (∼9–12 nm) spiderweb-like structure while making Nafion chains less anchored to the substrate, Cys forms a compact, thinner (∼0.7 nm), and more uniform monolayer next to the electrode that spatially organizes Nafion chains slightly away from the substrate, but in a more organized manner while minimizing direct ionomer-electrode interaction. This study aimed to understand how tailoring Au electrodes with Cys influences interfacial Nafion chain pinning, orientation, distribution, and ultimately proton conduction and water self-diffusion in confined thin-film environments. We combined experimental and computational approaches to reveal these interfacial processes caused by the Cys-based interfacial engineering, demonstrating an increase in in-plane proton conductivity of Nafion thin films.

Fig. 1 Schematic illustrating (left) a Nafion film on bare Au, (right) Cys modification of the Au substrate/electrode followed by Nafion thin-film deposition. The perfluorinated backbones of Nafion chains in the films are represented by orange–yellow chains, while the red balls represent terminal sulfonic acid (–SO₃H) groups on the side chains appended to Nafion backbones.

Materials and methods

Materials

A 20 wt% Nafion solution (EW ∼ 1100, IEC 0.909 meq g⁻¹) and cysteamine (Cys) were purchased from Millipore Sigma (Milwaukee, WI). Ethanol (anhydrous, denatured) was obtained from Decon Labs, Inc. (King of Prussia, PA). Silicon wafers coated with native silicon dioxide (SiO₂) (thickness of native oxide layer ∼1.72–1.79 nm) were purchased from Wafer Pro (San Jose, CA). Au interdigitated electrodes (IDEs), used for in-plane proton conductivity measurements, were obtained from Revtek Inc. (Torrance, CA). Au-coated crystals (5 MHz, Ti/Au/Ti/, polished, 25.4 mm (1 inch) diameter) for quartz crystal microbalance (QCM) measurements were purchased from Inficon (Syracuse, NY).

Sample preparation

Au deposition on silicon wafers via magnetron sputtering. Titanium/gold (Ti/Au) multilayer films were deposited on single-sided polished silicon (100) wafers using a magnetron sputtering system (AJA International, Inc.). The deposition was conducted at a low Ar pressure of 5 mTorr. The system also features a sample load-lock to minimize exposure of the main chamber to atmospheric contaminants. Substrates were mounted face-down on holders attached to a rotating propeller located along the central axis of the main deposition chamber. The propeller rotated at approximately 45 Hz to ensure uniform film coverage. Sputtering rates were controlled by adjusting the sputtering power while maintaining the constant Ar pressure. The substrate-to-target distance and gun tilt were typically set to 30 mm and 2.5 inches, respectively, as indicated by the system dials. For Au deposition, a DC power supply was used with a power setting of 75 W, yielding a deposition rate of approximately 0.97 Å s⁻¹. For Ti, a separate DC source was operated at 120 W, with a deposition rate of 0.39 Å s⁻¹. These sputtering conditions yielded a ∼10 nm-thick Ti-adhesion layer, followed by a 50 nm-thick Au layer on top. These Au-coated Si-wafers were then used for subsequent Cys modification, Nafion film coating, and film characterization using X-ray photoelectron spectroscopy (XPS), scanning electron microscopy with energy-dispersive X-ray spectroscopy (SEM-EDX), ellipsometry, contact resonance atomic force microscopy (CR-AFM), and water contact angle measurements.

Cysteamine (Cys) self-assembly on Au-substrates. Au substrates (Au-coated silicon wafers, IDEs, QCM crystals) were modified with 10 mM Cys in a light-protected environment for 6 h, following previous reports.^69–72 Briefly, the modification process involved cleaning the Au substrates with acetone and ethanol, followed by air-drying and treatment with UV-ozone for 20 min. The cleaned Au substrates were immersed in a solution of 10 mM Cys in ethanol in glass vials. This enables string interaction between thiol (–SH) groups of Cys and Au, and the Au–S formation leads to immobilization of Cys on the Au surface. The entire modification process was conducted under dark conditions for 6 h. This protocol was designed to ensure a controlled and light-protected environment throughout the Cys modification process, aiming to achieve a consistent and reproducible modification of the Au surface. After 6 h incubation, substrates were carefully removed, rinsed thoroughly with ethanol to eliminate any unbound Cys, and subsequently dried using N₂. The samples were stored in a desiccated chamber for further ionomer deposition and/or analysis. The Cys immobilization on Au was verified via X-ray photoelectron spectroscopy (XPS), scanning electron microscopy with energy-dispersive X-ray spectroscopy (SEM-EDX), QCM, and ellipsometry measurements.

Ionomer thin film deposition. To produce Nafion thin films on unmodified and Cys-modified Au, a 20 wt% Nafion stock solution was diluted with ethanol, then vortexed for 1 min, and ultrasonicated for 15 min. This yielded 1–5 wt% Nafion solutions, which were spin-coated at 3000 rpm for 40 s on Au and Cys–Au substrates to get Nafion films with a thickness ranging between 15–120 nm. Before spin-coating, all the substrates were cleaned with ethanol and UV-Ozone-treated for 20 min. The Nafion films spin-coated on Au and Cys–Au were then dried for 3 h at 42 °C, annealed at 100 °C for 7 h, and cooled down to r.t. overnight inside a vacuum oven (Model #1415, VWR, Radnor, PA).

SEM-EDX measurements. Elemental mapping to confirm Cys functionalization on Au was done using SEM-EDX. For this imaging, FEI Helios 660 NanoLab field emission SEM with EDX detector was used with an acceleration voltage of 20 kV in immersion mode. EDAX TEAM software was utilized for chemical mapping and elemental data acquisition. To enhance image contrast and protect the samples from electron beam damage, an Au layer was coated on top of the substrates using a Cressington Au sputter coater. All measurements were conducted under high vacuum conditions.

Water contact angle (WCA) measurements. WCAs of bare Au, Cys-modified Au substrates, and Nafion films on those substrates were determined using a Ramé-Hart Model 590 F4 Series Goniometer and Tensiometer. A 5 μL water droplet was applied using the sessile drop method, and a digital photograph was taken for precise WCA measurement.

In-plane proton conductivity measurements. To determine thin-film impedance and proton conductivity of Nafion films on Au and Cys–Au, Solartron 1260a Impedance/Gain-Phase analyzer (Solartron Analytical, Leicester, England) was used with a 2-probe technique as reported earlier.^1,10,73 For in-plane proton conductivity measurements using electrochemical impedance spectroscopy (EIS), Au interdigitated electrodes (IDEs) were fabricated on silicon wafers (110), featuring a thermally grown SiO₂ layer (200 nm) as an electrical insulator. Each IDE consisted of 150 gold teeth, each being 8 μm wide, spaced 40 μm apart, with an overlapping length of 4 mm. The IDEs were modified with Cys and spin-coated with Nafion as mentioned earlier in the sample preparation section. To ensure optimal cleanliness of the contact pads, the segments of the films covering them were removed using a small brush saturated with ethanol. Subsequently, the contact pads were left to air-dry. Following this, the Nafion thin film-coated and bare/Cys-modified interdigitated electrodes (IDEs) were promptly placed on a temperature stage set at 22 °C inside an environmental chamber. The Au tips of the probes were positioned carefully onto the IDE contact pads. The relative humidity (% RH) of the environmental chamber was controlled using silica gel for ∼0–5% RH, K-acetate for ∼25% RH, K₂CO₃ for ∼50% RH, NaCl for ∼75% RH, and Na₂SO₄ for ∼90% RH. By passing dry air (0.5 ft³ h⁻¹–235 cm³ min⁻¹) through a bottle/container containing the silica gel or saturated salt solution (immersed in a water bath at 23 °C), air with a specific % RH was generated, which was then passed through the environmental chamber. A humidity sensor was strategically placed at the chamber outlet to continually monitor and ensure the attainment and maintenance of the targeted % RH level within the chamber. At each humidity condition, the samples were stabilized for 1 h inside the environmental chamber prior to EIS measurement. Within this 1 h period, the impedance of the samples was measured periodically (at about every 14 min, from 1–10 kHz frequency range) to ensure sample stability at a specific % RH condition. Once stabilized, the impedance response of the Nafion film was recorded over a frequency range from 10 MHz to 1 Hz at an AC potential of 100 mV (choice of AC voltage is discussed in the discussion of Fig. S5a). The in-plane impedance data were analyzed using ZView software from Scribner Associates (Southern Pines, NC), which employed an appropriate equivalent circuit model (Fig. S5b) to fit the impedance spectra (Fig. 5a–c and S5c) and determine the film resistance (R_f) in the in-plane direction. The selection of components for the equivalent circuit model and the fitting process were based on our previous work.^1,10 Subsequently, the values of film resistance (R_f), film thickness (d_film), and the parameters of the interdigitated electrodes (IDEs)—including the spacing between teeth (S_teeth), overlapping length of each tooth (l), and number of teeth (N)—were utilized in eqn (1) to calculate the in-plane proton conductivity (σ_IP) of the films.

Please see the SI for representative Bode plots (Fig. S6a and b), residual plots (Fig. S6c–f), and χ² values (obtained upon fitting the impedance data, Table S3).

The protocols for measuring and fitting out-of-plane impedance data to obtain out-of-plane proton conductivity (σ_OP) are also discussed in the SI (Fig. S7 and eqn (S3)).

Water uptake measurement. To measure the effect of Cys on water sorption within Nafion films, Au-coated quartz crystal microbalance (QCM) crystals were modified first with Cys (as described earlier), then Nafion films were deposited on top. Water uptake of Nafion films on unmodified and Cys-modified Au crystals was measured using a QCM system (Stanford Research Systems (Sunnyvale, CA)). Nafion film-coated Au crystals were placed on a crystal holder within a custom-designed plastic humidity chamber. The humidity inside the chamber was maintained using appropriate saturated salt solutions. All QCL samples were equilibrated at each % RH for at least 1 h.

The dry mass of Nafion film on a QCM crystal (m_{dry Nafion film}) was determined at ∼5% RH by analyzing the corresponding frequency change (Δf = f_{dry Nafion film on crystal} − f_{bare crystal}).^7,74 To calculate water uptake or hydration number (λ_w) of Nafion films, the Sauerbrey equation was used. The choice of the Sauerbrey equation is appropriate for the film thickness range (∼15–100 nm thick) we studied, since within this thickness range, dissipative loss remains insignificant.^75,76 In the Sauerbrey equation, the bare crystal frequency was used as ∼5 MHz, active crystal area (A) as 1.27 cm², density of quartz crystal (ρ_q) as 2.648 g cm⁻³, and shear modulus of quartz (μ_q) as 2.947 × 10¹¹ g cm⁻¹ s⁻¹.^2,77 The total mass of water sorbed by Nafion films (m_{water sorbed}) at each % RH was obtained by analyzing the corresponding frequency change (Δf = f_{wet Nafion film on crystal} − f_{dry Nafion film on crystal}). The water mass sorption by each ionomer film was corrected further by subtracting the water absorption by bare QCM crystals at the same % RH, following the standardized protocol reported earlier.^11,77–79 This corrected mass of water sorbed by a Nafion film at certain % RH (m_{water sorbed, corrected}) and dry mass of Nafion film (m_{dry Nafion film}) were used to obtain hydration number (λ_w), moles of water per mole of sulfonic acid, at certain % RH using eqn (2):

Here, M_H2O is the molecular mass of water, and IEC is the ion-exchange capacity of Nafion.

Results and discussion

Au surface modification

To investigate the effect of interfacial engineering, we created a self-assembled layer of Cysteamine (Cys) on Au prior to spin-coating 15–100 nm thick Nafion films on top. A standard protocol was followed to immobilize Cys on Au via Au–thiol (–SH group of Cys) interactions, as detailed in the sample preparation section and illustrated in Fig. 1. Quartz crystal microbalance (QCM) measurements revealed that ∼0.515 ± 0.033 μg of Cys attached to the active area of the Au-coated QCM crystals (1.27 cm²) (Fig. S1), forming a ∼1 ± 0.27 nm (10 ± 2.7 Å) thick chemisorbed layer of Cys on Au (from ellipsometry). This experimentally determined thickness of the Cys layer closely matched what was predicted by molecular dynamics (MD) simulations (∼6.6 Å), and suggested the likelihood of the formation of a monolayer of Cys on Au, agreeing with the literature.^80,81

Cys adsorption was further confirmed via elemental analysis using XPS (Fig. 2a–c) and chemical mapping using SEM-EDX (Fig. 2d–g). XPS analysis showed the emergence of N 1s and S 2p peaks after treatment with Cys, a molecule having both N (amine, –NH₂) and S (thiol, –SH) (Fig. 2a–c). In contrast, no N 1s and S 2p peak was observed for bare Au samples (Fig. 2a–c). Specifically, the S 2p peak at 162.1 eV (Fig. 2b) was attributed to the formation of S–Au bonds,⁸² while a smaller secondary peak at 168 eV likely corresponded to oxidized sulfur species (sulfonates or sulfinates), commonly reported byproducts of Au-SH reactions.^82–84 Notably, our Cys immobilization protocol (6 h reaction under dark conditions) was optimized to minimize the formation of oxidized sulfur compounds, as indicated by the relatively low intensity of the 168 eV peak. The N 1s peak at 399.08 eV (Fig. 2c) further confirmed the presence of –NH₂ groups from immobilized Cys.⁸² The corresponding elemental analysis showed an increase in both nitrogen and sulfur content upon Cys modification: N: 0% (bare Au) to 6.22% (Cys–Au); S: 0% (bare Au), 4.61% (Cys–Au) (Table S1).

Fig. 2 (a–c) XPS (a) survey spectra, (b) S 2p and (c) N 1s spectra of bare Au (blue) and Cys on Au (Cys–Au, red) substrates. (d–g) SEM-EDX images showing the elemental mapping of sulfur: on Au (d) and Cys–Au (e); nitrogen: on Au (f) and (g) Cys–Au.

The water contact angle measurements (Fig. 3a) indicated that Cys-modification increased the hydrophilicity of the Au surface (bare Au: 63°; Cys–Au: 46°). Birefringence of this Cys layer on Au was positive (Δn = +0.0018), indicating surface normal orientation of Cys linkers on Au. A similar surface-perpendicular orientation of Cys on Au was reported by others.^84,85 These results aligned with the anticipated orientation of Cys molecules: thiol (–SH) groups of Cys bind to the Au surface, leaving terminal –NH₂ groups exposed at the air interface. Subsequent deposition of Nafion films on both unmodified and Cys-modified Au resulted in similarly hydrophobic film-air interfaces, with minimal differences in water contact angles (Nafion on bare Au: 98°, Nafion on Cys–Au: 92°) (Fig. 3a). This suggested that the Cys modification impacted Nafion chains residing next to the substrate interface more than residing near the air interface. Having said that, the difference in hydrophilicity between Nafion-substrate and Nafion-air interfaces saw an increase upon Cys modification.

Fig. 3 (a) Water contact angle of bare Au (yellow), Cys-modified Au (green), 60 nm thick Nafion thin film spin-coated on Au (blue), and 60 nm thick Nafion thin film on Cys-modified Au (red). (b) Hydration number (λ_w = moles of H₂O/moles of SO₃H) as a function of % relative humidity (% RH) for ∼60 nm thick Nafion films on unmodified (blue) and Cys-modified Au (red). Measurements were done in triplicate, and error bars were calculated based on standard deviations.

Based on AFM (Fig. S2), Cys modification reduced the roughness of the Au surface by one-third (1.56 nm (bare Au); 0.53 nm (Cys–Au)). This indicated that Cys modification created a smoother/flatter interface for Nafion chains to sit on the substrate. The increased substrate hydrophilicity and a flatter, smoother Cys interface may favor parallel Nafion backbone orientation⁸⁶ immediately next to Cys, which may dictate the –SO₃H groups of Nafion side chains to point toward Cys and interact with its –NH₂ groups. This experiment-based prediction agreed with findings from MD simulations (Fig. 4).

Fig. 4 MD simulations. (a and b) Probability distribution of cos θ (P(cos θ)) for the first layer of Nafion on (a) Au (considering 4 Å due to van der Waal forces and first Nafion layer ∼5.4 Å thick) and (b) Cys-modified Au (considering the thickness of Cys layer as ∼6.6 Å). (c and d) P(cos θ) beyond the first layer of Nafion on (c) Au and Cys–Au (d). Representative snapshots are provided of the Nafion molecule orientation with respect to the Au surface for the different maxima in the distributions.

Fig. 4a and b show the probability distribution of cos θ (P(cos θ)) for the first layer of Nafion on bare Au and Cys–Au, respectively, where θ is the angle between the Nafion bisector vector (V_NB) and surface normal vector (V_s) (Fig. S3). For instance, when the two segments of the Nafion backbone lie parallel to the substrate, the Nafion bisector lies in the plane of the substrate and forms a 90° angle with the surface normal, resulting in cos θ = 0. In contrast, the bisector aligns at 180° or 0° with the surface normal vector when the Nafion backbone fragments create a V––shaped conformation oriented downward or upward, corresponding to cos θ values of −1 and +1, respectively. As can be seen in Fig. 4a, Nafion chains in the first layer on bare Au exhibit a broad distribution of orientation, with cos θ ranging between −0.25 and +0.75. In contrast, Fig. 4b showed a sharp, high-intensity peak with a maximum at cos θ = 0.4 with a shoulder near cos θ = 0 for the first layer of Nafion chains on Cys-modified Au. Also, the edge sulfonyl groups tend to point towards the Cys surface. The near-zero cos θ value indicates that Nafion backbones predominantly lie surface-parallel on top of Cys-terminated Au. These computational findings align with experimental observations (flat, smooth Cys surface from AFM roughness), suggesting that Nafion backbones adopt a more parallel orientation with higher order next to the substrate when deposited on Cys–Au compared to bare Au. It is important to highlight that the simulated Nafion film thickness (∼3.4 nm) is notably smaller than the experimental range (∼15–100 nm), and thus the results primarily reflect interfacial effects rather than bulk-like regions.

Previous works have shown that surface-parallel backbone orientation of ionomer chains could lead to chain pinning to substrate (confinement) and reduced proton conductivity. But interestingly, in this work, the highly ordered surface-parallel backbone orientations of Nafion chains on Cys–Au made a positive impact, enhancing in-plane proton conductivity (σ_IP) (unlike Nafion films on unmodified substrates). We measured σ_IP of annealed Nafion films (∼15–100 nm thick) on unmodified Au and Cys-modified Au IDEs. Diameters of semicircles in Nyquist plots for σ_IP decreased as the Nafion film was deposited on Cys–Au (Fig. 5a–c). Since the decrease in semicircle diameter indicates a decrease in film resistance and an increase in film proton conductivity, the Nyquist plots indicated that Cys modification of the substrate interface can effectively improve in-plane proton conductivity (σ_IP) of Nafion in thin films. This improvement was consistently observed across the entire film thickness range we studied (Fig. 5d), with the deposition of a Cys layer beneath the Nafion film yielding up to a 4-fold increase in σ_IP. For instance, a ∼65 nm thick Nafion film showed σ_IP of 27.6 mS cm⁻¹ on Cys–Au, while it showed 7.5 mS cm⁻¹ on bare Au IDEs at 85% RH (Fig. 5d). Even under low humidity conditions (except 20% RH), a 2–3-fold enhancement in σ_IP was observed when Nafion films were deposited on Cys–Au (Fig. 5e). The improved in-plane proton conductivity could be a combined outcome of reduced interfacial confinement effect (discussed later with more evidence) as well as Nafion side chains creating highly ordered ion-conduction pathways in the in-plane direction (i.e., parallel to substrate). The thickness range investigated here is comparable to that of ionomer binder layers used in fuel cell electrodes, demonstrating the potential of interfacial engineering strategies to alleviate ion transport limitations in catalyst layers.

Fig. 5 In-plane Nyquist plots of (a) ∼15 nm, (b) ∼60 nm, and (c) ∼100 nm thick Nafion films on unmodified and Cys-modified Au (Cys–Au) electrodes at ∼85% RH. (d) In-plane proton conductivity (σ_IP) of Nafion films on unmodified and Cys–Au electrodes at ∼85% RH as a function of film thickness. (e) σ_IP of ∼60 nm thick Nafion films on Au and Cys–Au as a function of % RH. The equivalent circuit model used to obtain σ_IP values is shown in Fig. S5b in the SI. Measurements were done in triplicate (n = 3), and the results in Fig. 5d and e are presented as mean ± standard deviation (error bars).

To understand the contribution of Cys, we measured the conductivity of Cys-modified Au IDE (without a Nafion layer on top). The resistance of the Cys layer was very high: 8 × 10⁷ ohms, which was about 5 orders of magnitude higher than that of Nafion films. This confirmed that Cys's intrinsic conductivity was very low and had a negligible contribution to the improvement of the Nafion film conductivity.

The hydration behavior, studied using QCM, showed that Nafion films on Cys–Au absorbed slightly more water than those on bare Au (Fig. 3b). However, such a small difference in water uptake (λ_w ∼ 0.8 (Nafion on Au) vs. 1.2 (Nafion on Cys–Au)) should not account for the 4-fold improvement in proton conductivity (Fig. 5), suggesting a complex interplay of hydration with other factors, such as chain orientation within Nafion films, in modulating proton conductivity. In fact, based on experimental evidence and MD simulations, it can be inferred that the Cys layer was mainly responsible for internal change in Nafion chain distribution and organization next to the substrate (Fig. 4a and b) and within the film (Fig. 4c and d), causing improvement in σ_IP of Nafion thin films. The more surface-parallel orientation of the Nafion backbone induced by Cys–Au immediately next to it was discussed earlier. MD simulation also allowed us to predict the water (H₂O), hydronium (H₃O⁺) ion, and Nafion chain distribution next to the substrate.

Fig. 6a and b show the normalized density distribution and probability of having hydronium ions (H₃O⁺) and water (H₂O) (with representative snapshots in Fig. S4c and S4d, respectively), while Fig. 6c represents the probability of having Nafion chains as a function of z-distance from the Au surface (with representative snapshots in Fig. S4a and S4b). For Nafion films on bare Au, the density of H₃O⁺ ions was consistently low across 45 Å thickness starting from the Au interface (Fig. 6a, blue). While Nafion film on bare Au did not exhibit any H₃O⁺-rich region next to the substrate, a sharp increase in H₃O⁺-ion density was observed at ∼8 Å distance from Au when Nafion was deposited on Cys–Au (Fig. 6a, red). A similar spike in water (H₂O) density was observed at 9.3 Å distance from Au when Nafion was deposited on Cys–Au (Fig. 6b, red). This corresponded to a distance of ∼2.7 Å from the Cys layer. Also, the peak intensity of the first water layer was almost double for the Cys-modified Au surface as compared to the neat Au surface (Fig. 6b). This observation indicated that water molecules preferentially reside in close proximity to the Cys–Nafion interface, rather than within the Cys layer.

Fig. 6 Normalized density distribution of (a) hydronium (H₃O⁺) ion and (b) water (H₂O) for Nafion on Au and Cys–Au. (c) Probability distribution of Nafion chains with the z-distance from the substrate.

In addition to forming this water-rich layer next to Cys, the starting depth of the Nafion chain population shifted up next to Cys when the Au surface was modified with Cys (Fig. 6c). The shifted molecular distribution of Nafion chains (i.e. next to Cys, not next to Au) can be rationalized via several findings: (i) as per MD simulations, the Cys layer is ∼7 Å thick. (ii) As per QCM, area-specific density of Cys was 3.19 × 10¹⁵ molecules per cm², which was of the same order of magnitude of what is theoretically expected for a monolayer of Cys, assuming each Au atom is attached to one Cys molecule (1.58 × 10¹⁵ molecules of Cys per cm²). Thus, the dense Cys layer next to Au may have inhibited Nafion chain intercalation within Cys. The negligible density of H₃O⁺ ion and water within the Cys layer, on the other hand, can be rationalized by the hydrophobic nature of the aliphatic linker in the mid-region of Cys molecules. While Cys's hydrophobic linker may have inhibited the water localization within the Cys layer, the terminal hydrophilic –NH₂ groups of Cys may have stimulated Nafion (with –SO₃H), water, and H₃O⁺ ions sit on top of it, creating a water- and H₃O⁺ ion-rich layer next to Cys. This aligns with prior neutron reflectometry-based observations⁸⁶ where hydrophilic and smooth surfaces create a more water-rich layer next to the substrate when hydrated. Having said that, the average hydration number (moles of water per moles of –SO₃H) for the entire film did not see much change due to Cys modification (Fig. 3b). This likely suggested that the water-rich layer next to the substrate may have sacrificed hydration in the upper part of the Nafion film (i.e., adjacent to the air interface). This also suggested that the Cys-induced improvement in in-plane proton conductivity may predominantly be a localized effect. The orientational distribution, revealed from MD simulations, also supported that it is immediate next to Cys where chains are more parallel to substrate (beneficial for improving in-plane proton conductivity), but as we moved away from that position to up into the film (toward air interface), the Nafion chain orientation became more random with broad distribution of cos θ values (Fig. 4d), just like what we observed at similar position in Nafion films on bare Au (Fig. 4c).

The improvement in in-plane proton conductivity (σ_IP) was also in complete agreement with the in-plane (H₃O⁺) ion- and water (H₂O) self-diffusion coefficient (D_{self, H3O⁺, IP} and D_{self, H2O, IP}) in Nafion thin films. We observed about 4-fold increase in σ_IP (SI, Fig. 5d) as well as D_{self, H2O, IP} (4.3 × 10⁻⁹ m² s⁻¹ (Nafion on Cys–Au) vs. 1.4 × 10⁻⁹ m² s⁻¹ (Nafion on Au)) (Fig. 7a) and D_{self, H3O⁺, IP} (5.57 × 10⁻¹³ m² s⁻¹ (Nafion on Cys–Au) vs. 1.9 × 10⁻¹³ m² s⁻¹ (Nafion on Au)) (Fig. 7b). As is known, self-diffusion coefficient of a species is correlated to its mobility by the equation: D_self = μ × K_B × T, where μ is the mobility of the species, K_B is Boltzmann constant, and T is temperature. Thus, an increase in the water self-diffusion coefficient indicates an increase in water mobility.

Fig. 7 Mean square displacement of water and hydronium (H₃O⁺) ions in Nafion films on Au and Cys–Au: (a) water, in-plane; (b) H₃O⁺ ions, in-plane; (c) water, out-of-plane; (d) H₃O⁺ ions, out-of-plane. Corresponding self-diffusion coefficients (D_self) of species in a certain direction are shown in insets.

Especially, the likelihood of increasing near-substrate proton conductivity (discussed in earlier section), likely points towards alleviated confining effects due to interfacial interactions. Although the –NH₂ groups on top of Cys–Au can still induce some H-bonding and electrostatic interaction-induced confinement, the Cys-modified surface may still be considered as a source of soft confinement for both water molecules and Nafion chains, whereas bare Au as a source of hard confinement. This is because Au creates a rigid, non-penetrable boundary (hard confinement) that restricts molecular movement only on one side (i.e. on top of Au surface). In contrast, Cys, being a little away from Au, provides water and Nafion chains multi-directional freedom to move around, even though they are anchored to Cys to some extent and hanging in there (soft confinement). It is like a fruit hanging from a tree vs. one lying on the ground. The hanging fruit can swing freely in all directions (i.e., retains its dynamic freedom), while the fruit on the ground has its motion completely blocked on one side. Thus, despite some interfacial interactions, water may have a higher degree of rotational freedom, and Nafion chains may have higher segmental motion on top of Cys compared to on top of bare Au. Literature has shown this hard vs. soft confinement effect in terms of glass transition temperature, segmental dynamics, and structural relaxation.^87,88 This translational/rotational freedom (or mobility) favors proton transport.

Additionally, the storage modulus of Nafion films (Fig. 8) decreased when it was made on Cys-modified Au. This indicated that Cys modification allowed plasticization of Nafion film overall, which agreed with predicted faster water-ionomer chain dynamics owing to reduced confinement. This is in contrast with the typical response of Nafion films, which antiplasticized on affine substrates, but not in bulk membranes (free of interfacial effects).^5,6 This again presented a testament that if the interfacial effects are less, film stiffening will be less. Our prior observation in multiple cases also supported this connection, as we saw that when near-substrate –SO₃H accumulation decreased, substrate pinning and storage modulus reduced.^10,60 Overall, a simultaneous increase in in-plane proton conductivity and in-plane water self-diffusion coefficient indicates that Cys immobilization next to the substrate likely alleviated the confining effect exerted by the substrate on water and protons in Nafion film, leading to an increase in in-plane proton conductivity.

Fig. 8 Storage modulus of ∼60 nm thick Nafion films on unmodified and Cys-modified Au at different % RH. Error bars are from the standard deviations of triplicate measurements.

While we achieved an improvement in in-plane proton conductivity (σ_IP) in the presence of Cys, we did not see any significant increase in out-of-plane proton conductivity (σ_OP) (Fig. S7). A similar observation was made for out-of-plane water and H₃O⁺ ion self-diffusion coefficients (Fig. 7c and d). This could likely be explained by inspecting back the orientation of Nafion chains next to the substrate (simulation, Fig. 4a and b). MD simulations showed that Cys promotes a surface-parallel orientation of Nafion backbones near the substrate. This alignment arranges the –SO₃H groups parallel to the surface (Fig. 4b), creating lamellar, in-plane hydrophilic domains that facilitate lateral (in-plane) proton transport. However, such excessive surface-parallel ordering may have reduced vertical connectivity of ionic pathways and water channels, thereby limiting proton transport across ionomer films (σ_OP). Consistently, both out-of-plane proton conductivity and water/H₃O⁺ diffusion coefficients showed no significant improvement.

What did we learn?

In summary, these observations pointed toward the fact that localized chain order can impact proton conductivity in a certain direction in ionomer films. When Nafion chains align more parallel to the substrate, in-plane proton conductivity improves. However, excessive surface-parallel orientation can compromise out-of-plane conductivity. Similar conclusions have been drawn by others,^36,89,90 where oriented surface-parallel lamellae of ionic block copolymers on electrodes enhanced in-plane ionic conductivity, but suppressed through-plane conductivity. In contrast, our earlier work demonstrated that engineering the electrode interface with spiderweb-like amino silanes disrupted surface-parallel lamellae, while inducing some chain order in the out-of-plane direction, resulting in a significant increase in out-of-plane proton conductivity.¹⁰ In another relevant work, an ionomer having macrocyclic cavity-forming units and self-assembly capabilities created ellipsoidal features, where surface-parallel stacking of ionomer chains improved in-plane proton conductivity, while stacking of cavity-forming units perpendicular to the substrate enhanced out-of-plane proton conductivity.¹ These findings collectively demonstrate that improvement in proton conductivity in a certain direction has a lot to do with how the chains are packed and how the ion-conducting pathways are oriented. Therefore, strategic ionomer design, control over ionomer packing, and interfacial engineering can play a critical role in gaining control over ion-conduction pathways in electrochemical systems.

Our work draws attention to another intriguing observation. Birefringence (Δn) measurements,^10,25 surface enhanced FTIR,³² Raman⁴⁰ spectroscopy, and simulations (both from previous studies⁴¹ and this work showing chain distribution (Fig. 6c) and orientation (Fig. 4a and b)), indicate that on unmodified Au electrodes, Nafion chains or chain fragments tend to align parallel to the substrate, interacting directly with Au via Nafion's SO₃⁻ groups. In contrast, on Cys-modified Au, Nafion chains interact primarily with the terminal –NH₂ groups of Cys while remaining slightly lifted away from the substrate. Although Nafion chains experience interactions in both cases, the better in-plane proton conductivity observed on Cys-modified Au could be rationalized by a milder confinement effect exerted by Cys on the Nafion chains. Nafion chains anchored to Cys while hanging a bit up from the substrate experience less confinement and greater mobility compared to chains that are directly anchored to and lying on bare Au. Au, being a hard, non-penetrable wall-like substrate, totally blocks molecular mobility in one direction (i.e., below Au surface); while Cys-anchored, hanging ionomer chains can still swing and sway and retain multi-directional mobility. Similar concepts are discussed in polymer physics as the impact of hard vs. soft confinement.^87,91 Preventing direct contact of ionomer chains with substrate and/or modulating ionomer-water localization across the ionomer film may thus allow greater chain freedom and rotational relaxation of water molecules, both critical for proton transport.

Conclusions

In this work, we demonstrated that modifying Au electrodes with cysteamine (Cys) significantly enhances the in-plane proton conductivity of Nafion thin films by alleviating the confinement imposed by direct ionomer–metal interfacial contact and interactions. The Cys modification created a thin, uniform monolayer that facilitated a more ordered, surface-parallel Nafion backbone orientation near the interface while allowing the chains to remain slightly elevated from the substrate, thereby reducing direct chain pinning to Au and likely enabling greater segmental mobility. The reduced interfacial confinement, critical for proton transport, was also supported by the reduced storage modulus of the Nafion films on Cys–Au as compared to on bare Au. MD simulations showed that this interfacial engineering increased the population of water and hydronium ions near the substrate and enhanced their self-diffusion coefficients, aligning with the experimentally observed 4-fold increase in in-plane conductivity (σ_IP). However, out-of-plane conductivity (σ_OP) remained almost unchanged, as Nafion chains predominantly oriented parallel to the substrate, supporting in-plane proton transport rather than across the planes. Overall, this study highlights that interfacial design using tailored linkers, like Cys, can effectively mitigate confinement effects in ionomer thin films and improve proton transport in targeted directions. Our future work will focus on systematic exploration of linker chemistry and architecture to simultaneously optimize in-plane and out-of-plane ion conduction pathways, thereby advancing electrochemical device performance.

Conflicts of interest

The authors declare no conflict of interest.

Data availability

All data supporting the findings of this study are provided in the article and its supplementary information (SI). Supplementary information: QCM, AFM, out-of-plane EIS measurements, and simulation data. See DOI: https://doi.org/10.1039/d5lp00245a.

Acknowledgements

The research was primarily supported by the National Science Foundation (NSF) CAREER Award (NSF-DMR #1750040) and NSF-DMR# 2520078. The authors also acknowledge partial support through the U.S. Department of Energy (DOE), Office of Science, Basic Energy Sciences (BES) Early CAREER Award (#DE-SC0020336), and NSF-CBET#2310185. K. A. C. thanks the Fulbright Scholars Program and the University of Nebraska Collaboration Initiative Grant for their support. R. K. and G. A. acknowledge support from the NSF (#1845795). Part of this research (XPS, birefringence, AFM) was performed in the Nebraska Nanoscale Facility: National Nanotechnology Coordinated Infrastructure and the Nebraska Center for Materials and Nanoscience, which are supported by the NSF under Award ECCS: 2025298, and the Nebraska Research Initiative. The authors also acknowledge the Nano-Engineering Research Core Facility (NERCF) at UNL, which is partially funded by the Nebraska Research Initiative, for SEM-EDX and CR-AFM studies. K. A. C. and O. A. O acknowledge Sourav Sutradhar for assistance in AFM imaging experiments. The simulations were conducted on LSU HPC facilities and the LONI HPC facilities.

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Footnote

† These authors contributed equally.

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