Intrinsic proton relay in poly-phosphamides to bolster proton exchange membrane fabrication and electrocatalytic proton reduction

Abstract

Complex synthetic routes and over-swelling of the perfluoro/sulfonated Nafion-based proton-exchange membrane (PEM) materials at high temperatures and strong acidic pH lead to a continuous quest for stable non-fluoro organic polymers with high proton conductivity. Herein, porous organic polymers of 300–700 nm hydrodynamic diameter containing tripodal polyamine (PPA-1a) and/or ethylenediamine (PPA-2) as the linker and possessing a phosphamide {P(O)–NH} moiety in the repeating unit, as confirmed by the ³¹P and ¹³C (CPMAS) NMR and other spectroscopic characterization studies, are synthesized. A non-phosphamide tripodal polyamine (PPA-1b) is also prepared to establish the pivotal role of the {P(O)–NH} moiety in proton-conductivity and the electrocatalytic hydrogen evolution reaction (HER). Hierarchical mesoporosity with <10 nm average pore diameter and ∼11 m² g⁻¹ surface area of PPA-2 leads to a proton conductivity (σ) of 4.7 × 10⁻² S cm⁻¹ in aqueous solution at pH 4.5 and 358 K, superior to some commercial Nafions. The low activation barrier (E_a) of 0.12 eV indicates facile proton-hopping within the PPA-2 frame following a Grotthuss pathway. Conversely, the absence of phosphamide in PPA-1b and non-porosity results in low proton conduction. The density functional theory (DFT) study predicts that protonation at both “–P=O” and “–NH” sites of the phosphamide is energetically favorable to give stable tautomeric forms, which facilitate the proton-relay within the polymeric frame of PPA-2. The remarkably high proton conduction has led to the fabrication of PEMs using only 1 wt% PPA-2 with the poly(methyl methacrylate) (PMMA) and poly(vinyl alcohol) (PVA) supports, and the optically transparent membranes show structural stability after a successful proton-exchange study with 0.5 M H₂SO₄. Owing to the proton adsorption ability of the {P(O)–NH} moiety, fast proton relay within the framework, and the presence of the redox-active P^V center, PPA-2 behaves as an organo-electrocatalyst for the hydrogen evolution reaction (HER) with a low overpotential of 311 mV at 10 mA cm⁻². The pH dependency in the P^V/IV redox-couple identified in the cyclic voltammetry study indicates a proton-coupled-electron-transfer (PCET) mediated HER. At the same time, the proton adsorption on the {P(O)–NH} sites facilitates the Volmer step of the HER. In this study, phosphamide-based materials are exemplified as Nafion's alternative for PEM design and as metal-free energy materials for the HER.

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Biswarup Chakraborty

Biswarup obtained his PhD in chemistry in 2014 from the Indian Association for the Cultivation of Science (IACS), India. He performed postdoctoral research at Ben-Gurion University of the Negev, Israel. Before joining the Department of Chemistry, Indian Institute of Technology Delhi as an Assistant Professor in 2020, he spent more than a year as a postdoctoral fellow at the Technical University of Berlin, Germany. He currently works as an Associate Professor at IIT Delhi, and his research interests are photo(electro)catalytic water splitting and CO2 and NOx reductions, focusing on the structure–activity correlation and establishing reaction pathways.

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Introduction

The tremendous usage of fossil fuels has led to critical energy and environmental crises, ultimately resulting in the rapid depletion of global energy reserves.¹ Consequently, there is an urgent need to develop cost-effective and highly efficient renewable clean energy resources. Thereby, hydrogen, with its high specific energy density and carbon-neutral combustion waste, has emerged as a promising fuel source.² The use of proton exchange membranes (PEMs) in fuel cells is a sustainable and economical alternative to generate greener energy via converting chemical energy into electrical energy via the oxidation of H₂ at the anode with the reduction of oxygen at the cathode. The applicability of sustainable PEMs is not limited to fuel cells only, but they are also widely used in water electrolyzers for the production of H₂, electrochemical CO₂ reduction into valuable chemicals, wastewater treatment, and other applications.³ Furthermore, the use of proton exchange membranes (PEMs) in water electrolyzers is also a promising green hydrogen technology that utilizes electricity to split water into oxygen and hydrogen. PEMWEs paired with clean power sources can deliver green hydrogen, which is expected to occupy 40% of the green hydrogen market, while the remaining 60% comprises solid oxide and alkaline electrolyzers.⁴ In addition, they further serve as critical components in CO₂ electrolyzers by enabling efficient proton transfer, maintaining ion selectivity, and supporting stable operation under acidic conditions, thereby enhancing the overall efficiency and selectivity of CO₂ electroreduction.⁵ Again, PEMs are also integral to various electrochemical systems, notably in wastewater treatment applications such as microbial fuel cells (MFCs), where PEMs facilitate the selective transport of protons from the anode to the cathode while preventing electron crossover, which is essential for efficient electricity generation and effective wastewater treatment.⁶ However, despite their applicability, the most essential component of PEMFCs is the membrane that behaves as a mediator for the exchange of protons during the operation of proton exchange membrane fuel cells (PEMFCs). Owing to the extraordinary proton conductivity and chemical durability, perfluoro-sulfonated polymer (PSP) based materials such as Hyflon, Flemion, and most frequently Nafion-based membranes are traditionally being used to build PEMFCs.⁷ However, a few concerns, such as high permeability to fuel, loss of mechanical properties at higher temperatures, intricate multistep synthesis steps for the synthesis of the membrane, and its environmental incompatibility, have made PEMFCs economically challenging. The excessive fluorine content in these PEMs makes them environmentally hazardous chemicals.⁸ To circumvent these obstacles, extensive studies have been conducted to replace fluoro-substituted PEMs with low-cost, eco-friendly, and highly proton-conducting membranes.^9–11

In the last decade, metal–organic frameworks (MOFs) have been used as PEM materials.¹² However, MOFs suffer from low conductivity under low humidity conditions, low acid stability over prolonged periods of operation, and degradation of the material to metal oxides at high temperatures, which limit their applications. This has led to the exploration of covalent frameworks (COFs).^13,14 Mimicking the terminal structure of Nafion, researchers have designed polyprotic acids such as sulfonic acid, phosphoric acid, and phosphotungstic acid-rich COFs that enhance the proton conduction many fold in the materials.^15–17 However, the hydrolytic instability of MOFs and synthetic rigidity in COFs and zeolites often limit their applicability under harsh thermal and chemical conditions. Most recent development strategies use modified perfluoro sulphonic acid (PFSA) polymers and acid-functionalized aromatic hydrocarbon-based polymers such as poly (arylene ether) (PAE),¹⁸ poly (ether–ether ketone) (PEEK),¹⁹ poly (phenylene sulfide) (PES),²⁰ poly (arylene ether sulfone) (PAES),²¹ poly (phenylene)s (PP),²¹ and polyimide (PI)²² as potential PEM materials. Sulfonated or phosphonated polybenzimidazoles (PBI), polybenzoxazoles (PBO), and polybenzothiazoles (PBT) have also been investigated for possible use as PEMs.²³ However, these polymers with sulfonic acid side chains undergo desulfonation with the increase in temperature. The high degree of sulfonation in the sulfonated polymers often results in excessive water uptake and over-swelling, which leads to the loss of mechanical characteristics and reduction in dimensional stability, resulting in lower proton conductivity.^24,25 However, non-sulfonated nitrogen-rich polymers like imidazole-containing polymers show promising scope as membrane materials.²⁶ Alcohol group functionalized membranes, such as poly (ethylene oxide) (PEO),²⁷ poly (vinyl alcohol) (PVA),²⁸ poly (acrylamide) (PAAM),²⁹ TPOP,³⁰ and poly (vinyl pyrrolidone) (PNVP),³¹ are also regarded as potential membrane materials. These materials exhibit low proton conductivities (∼10⁻³ S cm⁻¹) without sulfuric acid or phosphoric acid as an addendum. However, the membranes show poor mechanical stability, especially at higher temperatures and with high acid concentrations. Notably, nitrogen-containing heterocyclic rings like imidazole and triazine, due to their intrinsic Brønsted basicity, have recently been favored as building blocks for constructing proton-conducting materials. The microporous framework structures of such materials act especially as proton channels. Chakraborty and co-workers have recently reported a triazole-functionalized polymer framework, and due to Brønsted basicity and porosity, the material boosts proton conductivity.³⁰ However, the fundamental problem with these materials³² is the sustainability of polymer functionalized membranes under harsh acidic or alkaline conditions and stability at high temperatures for a prolonged period. Furthermore, the proton conductivity of most of the reported materials is recorded in the solid state under 20–100% humidity conditions^32–36 instead of embedding them into the membrane and studying the proton conductivity with the PEMs under actual solution-state conditions. These challenges often screened out most of the solid-state proton conductors from practical deployment in PEMFCs. At the same time, proton-conducting porous organic polymers or framework materials are of significant interest as metal-free cathode materials for electrocatalytic water splitting and or battery applications.^37–39 Porous proton exchange membranes (PEMs) offer enhanced hydrated H-bonded network structures for smoother proton hopping and improved gas diffusion compared to non-porous counterparts, leading to more efficient electrochemical performance in PEM fuel cells.⁴⁰ The permanent porosity facilitates effective proton transport while maintaining dimensional stability and minimizing mass transport limitations under varying operating conditions. However, the inherent electron and proton conductivity limitations within these organic frameworks have necessitated new and innovative strategies. Incorporating heteroatoms like phosphorus, sulfur, and nitrogen or modifying the organic backbone can substantially enhance ion conductivity, surface adsorption, and electrochemical activity. P- and N-rich porous organic polymers can be interesting materials not only for fabricating PEMs⁴¹ but also as metal-free organic electrocatalysts for hydrogen production.⁴²

In this context, porous organic polymers (POPs) are a promising alternative possessing structural diversity in the polymeric network with superior chemical and thermal robustness. Among them, polyphosphamide-based POPs represent a unique subclass that integrates the hydrogen-bonding and proton-relaying capabilities of the phosphamide moiety –{P(O)–NH}– into a highly crosslinked, permanently porous architecture. The incorporation of tripodal or ethylenediamine linkers within these networks facilitates hierarchical mesoporosity and enhances structural flexibility, positioning polyphosphamide POPs as a versatile platform that bridges the design principles of MOFs, COFs, and zeolites while introducing novel redox and proton-conductive functionalities for energy and catalytic applications. In this study, two poly-phosphamides (PPA-1a and PPA-2) and one polyamine (PPA-1b) material are prepared, and the presence of the phosphamide {P(O)–NH} repeating unit in PPA-1a and PPA-2 has been ensured by characterizing them through various spectroscopic and microscopic techniques. To demonstrate impressive proton conduction by these P- and N-rich porous organic materials, the materials are fabricated on a carbon cloth (CC) surface, and impedance analyses are conducted at various pH levels and temperatures. The comparison of solution-state proton conductivity (σ) and the activation energy (E_a) is performed to emphasize the essential role of the {P(O)–NH} repeating unit in facile proton conduction. Density functional theory (DFT) studies have also been conducted to show that the facile protonation on the P–O and –NH groups is energetically more feasible. Furthermore, to showcase the practical implications of poly-phosphamides as a PEM material, transparent membranes are fabricated, and proton conduction across the membrane has been demonstrated. Due to the electron-rich nature and the presence of the redox-active {P(O)–NH} moiety in the polymers, they are further used as the metal-free electrocatalysts for the HER study. The comparative study with poly-phosphamides (PPA-1a and PPA-2) versus polyamine (PPA-1b) infers the crucial role of phosphamide {P(O)–NH} in faster proton adsorption and migration through the straight chain skeleton of the polymer to facilitate the proton relay. The redox-active P^V centers of poly-phosphamides participate in proton-coupled electron transfer for the HER. Therefore, poly-phosphamides are demonstrated here as viable alternatives to Nafion-based PEM materials and as an organic material for renewable energy applications. This study also showcases an approach to bypass the need for precious and environmentally hazardous metal ions and sulfonate functionalized additives for PEMs.

Experimental section

Synthesis of PPA-1a and PPA-2

PPA-1a and PPA-2 were prepared through the reported procedure with partial modification.^43–45 In a 50 mL two-neck round bottom flask attached with a condenser, tris(2-aminoethyl)amine or ethylene diamine (0.6 mmol) was dissolved in 15 mL dry toluene, followed by the addition of 2.0 mmol (279.1 µL) triethyl amine. The solution and overhead space were degassed with nitrogen, and the temperature of the R. B. flask was maintained at 0 °C. Phenylphosphinic dichloride (0.9 mmol, 128 µL) was added to the reaction mixture, and 1 h of stirring at 0 °C, and the mixture was heated to 60 °C for 24 h. After cooling down, the product was obtained after washing with toluene, ethanol, chloroform, and diethyl ether for PPA-2. PPA-1a was further washed with water. Finally, the semi-solid obtained product was left at 70 °C overnight in a hot air oven. Yield: 0.34 g.

Synthesis of PPA-1b

PPA-1b was prepared following a procedure reported in the literature.³⁰ Tris(2-aminoethyl)amine (TREN) (1 mmol) was dissolved in 10 mL DMSO, and formaldehyde (1 mmol in 10 mL DMSO) was slowly added to the TREN solution. The whole reaction mixture was subjected to refluxing conditions in an oil bath at 160 °C for 12 h. After cooling down to RT, the pale white precipitate was isolated by centrifugation, followed by washing thoroughly with water and ethanol to be finally dried at 70 °C in the oven. Yield: 0.42 g.

Proton conductivity study

Electrochemical impedance spectroscopic (EIS) studies were used to measure the solution-state proton conductivity of the polymeric materials. EIS studies were conducted in a three-electrode set-up, using Ag/AgCl as the reference electrode, Pt wire as the counter electrode, the polymeric materials drop-casted on carbon cloth (CC: 1 cm² geometric surface area) as the working electrode, and 0.5 M Na₂SO₄ as the electrolyte solution. To the electrolyte, 0.5 M H₂SO₄ solution was added to achieve a different pH of the solution for the proton conductivity studies. For the preparation of the working electrodes for proton conductivity measurements, a slurry ink was made using 3–12 mg of the PPA-1a, PPA-1b, and PPA-2a in a Nafion : ethanol mixture (40 : 60), followed by drop casting on the carbon cloth (CC). A scan of the amplitude of the sinusoidal wave in the fixed frequency range of 100 kHz to 1 mHz was applied to obtain the EIS spectra. All the EIS data were fitted into the Nyquist plot and the corresponding equivalent RC circuit model. The diameter of the semicircle in the Nyquist plots represents the value of the charge-transfer resistance (R_ct). A Gamry 1010E-29165 potentiostat, commanded through the Gamry Framework software package, was used to obtain charge transfer resistance from Nyquist plots, which has been used to measure the electrical conductivity (σ) using eqn (1).

σ = L/R_ctA (1)

where L and A are the sample's thickness and contact area, respectively, and R_ct is the resistance. The thickness and area of the samples were 0.032 cm and 1 cm², respectively.

The activation energy was estimated using the Arrhenius relation eqn (2).

σ_dc = A exp(−E_a/k_BT) (2)

where σ_dc is the proton conductivity, E_a is the activation energy, k_B is the Boltzmann constant, and T is the absolute temperature.

Preparation of the PPA-2/PVA film

For the PPA-2/PVA film preparation, 99 mg of polyvinyl alcohol (PVA) polymer was dissolved in the appropriate amount of water with continued stirring at 60 °C. PPA-2 solution (1 mg of PPA-2 dissolved in 0.3 mL of water) was added to the aqueous PVA solution. The transparent solution of PPA-2 and PVA was poured into a circular mould and dried at 60 °C for a few hours. The thin film of PPA-2/PVA was then microscopically characterized.

Preparation of the PPA-2/PMMA film

For the PPA-2/PMMA film preparation, 99 mg of polymethyl methacrylate (PMMA) polymer was dissolved in a desired volume of dimethyl sulfoxide (DMSO) while stirring constantly at 130 °C. Afterward, PPA-2 solution (1 mg of PPA-2 in 0.3 mL of H₂O) was added to the PMMA-DMSO solution, which resulted in instant aggregation. The aggregated mixture was dissolved by adding a few mL of DMF, followed by heating. Then, the transparent solution of PPA-2 and PMMA was poured into a circular mould and dried overnight at 80 °C. The thin film of PPA-2/PMMA was then microscopically characterized.

Proton exchange (permeation) membrane test

A two-compartment setup was built by putting the PPA-2/PVA or PPA-2/PMMA film in between two compartments (A and B). In compartment A, an acidic (0.5 M H₂SO₄) solution of pH 0.9 was placed, while compartment B was filled with distilled water of pH 7.9. A pH meter was placed in compartment B, and the pH of the solution was noted with time.

Electrochemical hydrogen evolution reaction (HER)

A similar three-electrode cell setup (mentioned above for the EIS study) was used to study the electrocatalytic HER with PPA-1a, PPA-1b, and PPA-2a deposited on CC electrodes. For the HER study, linear sweep voltammetry and cyclic voltammetry (CV) studies were performed in 0.5 M H₂SO₄ electrolyte and using a Hg/Hg₂SO₄ reference electrode. Mass loading (3–12 mg) of PPA-1a, PPA-1b, and PPA-2a was varied for optimum activity. LSV and CV measurements were done at the scan rates of 1 mV s⁻¹ and 5 mV s⁻¹. The LSV data provided are calibrated to the reversible hydrogen electrode (RHE) scale using the relation provided below:

E(RHE) = E(Hg/Hg₂SO₄) + 0.098 V + (0.059 × pH) V (3)

The LSV plots were corrected with 85% iR compensation using the solvent resistance (R_s) obtained from the Nyquist plot (impedance study). For the double layer (C_dl) measurements, CV cycles were recorded within −0.8 to 0.7 V (vs. RHE).^46,47

Results and discussion

Synthesis and characterization of polyphosphamides

Polyphosphamides (PPAs) were prepared via stoichiometric condensation of di-/tri-amines with phenylphosphinic dichloride (PPDC) (Scheme 1) and para-formaldehyde. PPA-1a was synthesized by following the previously reported methodology.⁴⁴ PPA-1b was synthesized through condensation of N′,N′-bis(2-aminoethyl)ethane-1,2-diamine with formaldehyde in DMSO at 160 °C.³⁰ PPA-2 was prepared via condensation of a stoichiometric amount of ethylene diamine and phenylphosphinic dichloride (PPDC) in the presence of triethylamine in dry tetrahydrofuran (THF) under refluxing conditions at 125 °C. The ¹³C CP-MAS solid-state NMR spectrum of PPA-1b showed three sharp peaks at 37 ppm, 46 ppm, and 58 ppm, which correspond to the adjacent and distant methylene carbon to the tripodal nitrogen and the linker methylene group (Fig. S1a†). The FTIR spectrum gives all characteristic bands for secondary amines and aliphatic groups (Fig. S1b†). The PXRD pattern of PPA-1b displayed a broad band at a 2θ value of 15–25°, indicating the amorphous nature of the material (Fig. S1c†). Furthermore, the morphology and atomic percentage of PPA-1b were evaluated by field emission scanning electron microscopy (FESEM) imaging, which suggests aggregated particles of irregular shape in the surface morphology (Fig. S2†). The elemental mapping from Energy Dispersive X-ray (EDX) implies the homogeneous distribution of carbon (66.2%) and nitrogen (33.8%) elements in PPA-1b (Fig. S3 and S4†), which is also complemented by an elemental distribution of carbon (56.9%), hydrogen (12.07%) and nitrogen (31.03%) by CHN analysis. The ³¹P CP-MAS solid-state NMR of PPA-2 revealed a prominent peak at 14.1 ppm with two pairs of satellite peaks at 76.2 ppm and −48.4 ppm, 137.3 ppm, and −108.9 ppm, implying the presence of a phosphamide unit in the polymer (Fig. 1a). In addition, the ¹³C CP-MAS solid-state NMR spectrum of PPA-2 depicted carbon signals at 38.3 ppm and 129.4 ppm, corroborating the ethylene moiety and phenyl ring, respectively (Fig. 1a). A weak carbon signal appeared at 44.9 ppm, perhaps due to the self-condensation of ethylene diamine or the branching of polymers.⁴⁸ The size distribution of the aggregation of the as-synthesized polymers was studied thoroughly by dynamic light scattering (DLS) experiments. The dynamic light scattering (DLS) study of PPA-1a, PPA-1b, and PPA-2 dispersed in DMSO has revealed a hydrodynamic diameter of 459 nm, 614 nm, and 257 nm, respectively, which emphasizes self-aggregation of the polymeric networks to deliver a macromolecular structure of the materials (Fig. S5a–c†). However, the narrow hydrodynamic size distribution in DLS for PPA-2 highlighted less branching or self-aggregation present in the solution phase. The FTIR analysis suggests a characteristic band at 1148 cm^—1, which represents the incorporation of a phosphamide functionality in PPA-2 similar to that observed for PPA-1a and other polyphosphamides reported recently (Fig. S6a†).^44,45,49 Furthermore, a Raman band at 995 cm⁻¹ corresponds to the phosphamide group present in PPA-2 (Fig. 1b). Similarly, the broad diffraction observed in the powder X-ray diffraction (PXRD) of PPA-2 highlighted the polycrystalline nature (Fig. S6b†). The size distribution of the polymer in the solid state was analyzed by small-angle X-ray scattering (SAXS) studies. The pair distance distribution function (PDDF) revealed an aggregate size of 70–80 Å corresponding to the partial crystalline nature of PPA-2, whereas the Guinier plot revealed the aggregation of particles (Fig. S7a–c†). The incorporation of the thermodynamically robust –{P(O)NH}– moiety into the polymeric phosphamide materials synergistically enhanced the thermal stability, which resists structural degradation at elevated temperatures up to 250–300 °C. TGA results for PPA-2 often reveal a multi-step degradation process, where organic substituents decompose at lower temperatures, followed by the breakdown of phosphamide bonds. In DSC, phosphamide-functionalized materials exhibit a change in heat flow at a similar temperature compared to the decomposition temperature in TGA due to increased intermolecular interactions and restricted segmental mobility.⁵⁰ Thermogravimetric analysis (TGA) demonstrated the thermal stability up to 270–280 °C for PPA-1a and PPA-2, with approximately 80% weight loss observed above this temperature due to the chemical instability of the framework (Fig. 1c, S8a and b).† However, PPA-1b showed structural degradation even at 130 °C. The differential scanning calorimetry (DSC) study further validated the TGA data. The differential scanning calorimetry (DSC) study of PPA-2 and PPA-1a demonstrated a sharp shift in the heat flow after 310 °C, whereas, for PPA-1b, the sharp shift in the heat flow was noted at around 140 °C as expected owing to the absence of phosphamide functionalization (Fig. 1c, S8c and d†).

Scheme 1 Schematic representation of the synthesis of polyphosphamides PPA-1a and PPA-2 using tris(2-aminoethyl)amine (TREN), ethylene diamine (EDA), and phenylphosphinic dichloride (PPD).

Fig. 1 (a) ³¹P and ¹³C CPMAS NMR spectra of PPA-2, (b) Raman spectrum of PPA-2, (c) thermal analysis using TGA and DSC for PPA-2, (d) deconvoluted core-level P 2p and N 1s XPS spectrum obtained for PPA-2, (e) HRTEM images of a selected area (inset: SAED pattern of PPA-2), (f) FESEM image of a selected area of PPA-2 and the corresponding EDX elemental mapping for (g) C, N, O, and P, and (h) N₂ sorption isotherm measured with p-PPA at 77 K (inset: pore distribution curve using the BJH model).

The elemental valence state and composition of PPA-2 were further investigated by X-ray photoelectron spectroscopy (XPS). The survey scan from XPS revealed the presence of C, N, O, and P in the PPA-2 with an approximate ratio of 6.9 : 1.1 : 1.1 : 0.9 (Fig. S9a†), which corroborates with the EDX mapping with an elemental ratio of 7.1 : 1.2 : 1.1 : 1.2. The X-ray photoelectron spectroscopy (XPS) study of PPA-2 found that the binding energy at 284.7 eV in the XPS spectrum corroborates to core level C 1s. Furthermore, the binding energy at 283.1 eV is attributed to the aromatic carbon of the phenyl ring (Fig. S9b†).⁵¹ The XPS spectrum of core-level P 2p with binding energy at 131.8 eV indicates the presence of a single peak at binding energy for the P(V) state in the phosphamide polymer corresponding to the spin–orbit components 2p_3/2 and 2p_1/2 (Fig. 1d).⁵² The core level XPS spectrum of N 1s displays binding energy at 399.6 eV, confirming the existence of P–N–C linkage (Fig. 1d). The O 1s XPS spectrum of oxygen is deconvoluted to 530.6 eV and 528.9 eV, corresponding to the O²⁻ of the phosphamide moiety alongside a minor peak at 528.9 eV, perhaps owing to the tautomeric form of the phosphamides or the adsorbed moisture (Fig. S9c†).⁵³ High-resolution transmission electron microscopy (HRTEM) imaging of PPA-2 revealed needle-shaped particles throughout the material, and the SAED pattern displayed some extent of crystallinity throughout the specimen, complementing the results obtained from the PXRD diffractogram of PPA-2 as well as the particle size obtained from the DLS study (Fig. 1e). Field emission scanning electron microscopy (FESEM) imaging unveiled a plate-shaped regularity with some extent of porous surface morphology due to the spontaneous aggregation of particles (Fig. 1f and S10†). Elemental mapping derived from FESEM-energy dispersive X-ray (EDX) analysis revealed a nearly uniform distribution of C, N, O, and P, and the EDX spectrum provided an atomic ratio of C : N : O : P that closely matched the calculated ratio for the repeating unit of the polymer (Fig. 1g and S11†). The CHN analysis from the combustion study confirmed the chemical composition of PPA-2 as C_6.8H₁₀N_1.2P_1.1O_1.1, translating to C 54.27%, H 5.79%, and N 9.23%. Notably, the elemental analysis revealed a phosphorus content of 10.3% in PPA-2.

The specific surface area (SA_BET) for all three polymers was determined by N₂ adsorption and desorption at 77 K within the range of partial pressure of 0 to 1. The Brunauer–Emmett–Teller (BET) study was used to interpret the data. Furthermore, the hysteresis loop extends from the lower to the higher pressure region, reflecting a hierarchical pore architecture.⁵⁴ From the multi-point BET study of PPA-2, a surface area (SA_BET) of 11.202 m² g⁻¹ with a total pore volume of 1.7255 × 10⁻¹ cc g⁻¹ corresponding to slit-type pores for PPA-2 was observed, whereas no specific ordering of the pores was observed for PPA-1a and PPA-1b (Fig. 1h). The pore size distribution (PSD) using the BJH method for PPA-2 demonstrates that most pores are of diameters <10 nm with a frequent distribution of pores at 1.19 nm, 2.01 nm, 2.55 nm, and 8.13 nm with an average diameter of 4.21 nm. For PPA-1a and PPA-1b, no hysteresis loop with a specific ordering of pores was observed (Fig. S12a and b†).⁵⁵ For PPA-1a and PPA-1b, a specific surface area (SA_BET) of 0.048 m² g⁻¹ and 0.142 m² g⁻¹ with a total pore volume of 3.456 × 10⁻⁴ cc g⁻¹ and 5.357 × 10⁻³ cc g⁻¹, respectively, was recorded. Furthermore, the pore size distribution (PSD) plot from the BJH method exhibits mostly flat distribution patterns for PPA-1a and PPA-1b (Fig. S13a and b†). Among PPA-2, PPA-1a, and PPA-1b, only PPA-2 possessed a hierarchically mesoporous architecture in its crystalline surface, implying the existence of a pore channel within the material.^10,56,57

Proton conductivity study

Spontaneous proton conduction through porous membranes is an essential requirement for the proton-exchange membrane (PEM) fuel cells or electrolyzers.⁵⁸ Traditionally, PEMs are made of sulfonated fluoro polymers such as Nafion. In the last decade, organic polymers functionalized with sulphonate, phosphate, carboxylate, or amine groups have been shown to be potential PEM materials.^21,59 However, the studies are limited to four-probe solid-state impedance analysis under 50–100% humidity conditions.³² Solution-state proton conduction data were not reported, which are practically more relevant for the PEM study. Recently, Chakraborty and co-workers reported a triazole-based organic polymer that shows an impressive solution-state proton conductivity (0.025 S cm⁻¹).³⁰ Solution-state proton conductivity is more realistic in comparison to the solid-state measurement since the actual deployment of membranes for proton conductivity involves two different solutions separated by a membrane.⁶⁰ Therefore, the proton conductivity for all three materials, PPA-1a, PPA-1b, and PPA-2, was recorded by casting them on the electrode surface and at variable pH, and within the temperature range of 25–85 °C, electrochemical impedance spectroscopy (EIS) was performed and the subsequent Nyquist plots were obtained. During the EIS study in the three-electrode configuration using the material deposited on carbon cloth, as the pH of the electrolyte was varied from 8.9 to 4.5, proton conductivity was substantially enhanced for PPA-1a, PPA-1b, and PPA-2. At room temperature and pH 4.21, the proton conductivity of PPA-1a, PPA-1b, and PPA-2 was found to be 7.33 × 10⁻³ S cm⁻¹, 4.3 × 10⁻⁴ S cm⁻¹ and 8.53 × 10⁻³ S cm⁻¹, respectively (Fig. 2a, S14a–c and Table S1†). There was nearly a tenfold enhancement of proton conductivity of PPA-1a and PPA-2 compared to PPA-1b, which highlights the crucial role of phosphamide units present in PPA-1a in proton conduction. However, the proton conductivity at room temperature within the pH range of 8.9–4.2 of PPA-2 was slightly higher than that of PPA-1a, which can be attributed to the crystallinity and porous channels present in PPA-2 in facilitating proton conductivity. Thermal activation can facilitate a faster shuttling of protons within the polymer frame. As the cell temperature increased from 303 K to 358 K, keeping the pH fixed at 4.5 showed the shrinking of the semi-circle of the Nyquist plot, which corresponded to the notable drop of the charge transfer resistance (R_ct) value (Fig. 2b). A sharp and steep linear growth of the proton conductivity (σ) to a low value of 0.047 S cm⁻¹ (at 358 K and pH 4.5) for PPA-2 was observed with the increase in cell temperature. In contrast, for PPA-1b, an almost negligible increase in the proton conductivity was observed, while a moderate enhancement of proton conductivity with temperature was observed for PPA-1a (Fig. 2c and Table S2†). Notably, PPA-1a and PPA-1b exhibited the σ value of 0.013 S cm⁻¹ and 0.00197 S cm⁻¹, respectively, at pH 4.5 and at 358 K (Table S3, Fig. S15a and b†). The superior cell performance of PPA-2, with a proton conductivity of 0.047 S cm⁻², is better than the materials reported in the literature. As the temperature of the electrolyte solution increases, the mobility of protons in water-filled pores within the polyphosphate material (PPA-2) facilitates a proton relay through hydrogen-bonded water molecules, aided by an excess of H⁺ ions. Recent studies have revealed that certain porous organic polymers rely on the presence of Brønsted components within their molecular structures to facilitate proton conduction.⁶¹ Additionally, amine functionality in metal–organic frameworks has been shown to promote proton conductivity. However, all of these materials were reported using pelletized samples under highly humid conditions and solid-state impedance analysis. However, under harsh alkaline or acidic conditions, the robustness and recyclability were not studied, which is a prime factor for practical deployment in PEM fuel cells. The metal-based materials are also not reported under actual solutions separated by membranes, whereas PPA-2, as a metal-free material, was incorporated into PVA and PMMA-based membranes for applications.

Fig. 2 (a) Plot of proton conductivity with respect to pH, (b) Nyquist plots (from EIS) at different temperatures for PPA-2. (c) Change in proton conductivity with temperature and (d) Arrhenius plot for the activation energy of PPA-2, PPA-1a & PPA-1b. (e) Energy profile diagram for the protonation of a phosphamide species with a hydronium ion, (f) equilibrium pathway of different tautomeric forms of the phosphamide and hydronium ion obtained during the proton hopping, and (g) NBO analysis of the different tautomeric forms using natural population analysis.

Activation energy (E_a) is a critical parameter for the understanding of proton conduction or shuttling mechanisms.⁶² Proton shuttling can occur via two distinct pathways, namely the Grotthuss mechanism and the vehicle mechanism. In the Grotthuss mechanism, protons migrate through successive hydrogen bond dissociation and formation with adjacent water molecules, effectively “hopping” between two oxygen atoms of different water molecules. This pathway is characterized by a lower activation barrier (E_a < 0.4 eV) due to the relatively moderate energy required to break a hydrogen bond (2–3 kcal mol⁻¹ or 0.09–0.13 eV). Conversely, the vehicle mechanism involves proton migration through diffusion, requiring a higher activation energy (E_a > 0.4 eV).^63,64 For the temperature-dependent proton conduction study, the activation energy was estimated to be 0.12 eV and 0.14 eV for PPA-2 and PPA-1a, respectively, whereas a high activation energy of 0.59 eV for PPA-1b (Fig. 2e). The E_a data, therefore, strongly suggest that proton conduction in PPA-2 and PPA-1a primarily adheres to the Grotthuss mechanism facilitated by the phosphamide unit. Apart from the presence of the phosphamide functionality, the superior proton conductivity of PPA-2 over PPA-1a is also facilitated due to the dual structural effect: (i) the presence of some extent of crystallinity within the material, and (ii) mesoporosity in the material. In contrast, the absence of both a phosphamide unit and a porous channel in PPA-1b leads to a predominance of the vehicle mechanism, demanding higher energy for the proton conduction.⁶⁵ Overall, the remarkable pH responsiveness of PPA-2 emphasizes the significance of the phosphamide unit and crystalline porous architecture in enhancing proton conductivity. Furthermore, these findings shed light on proton relay within the poly-phosphamide framework of PPA-2, which has profound implications for the development of materials for PEMs.

Density functional theory (DFT) studies were performed on all the plausible tautomeric forms of phosphamide repeating units to understand the proton hopping within the POP. The Gibbs free energy change associated with the protonation of the “P=O” and “–NH” functionalities of the POP repeating unit by a hydronium ion was negative to form the oxonium (ΔG: −62.3 kcal mol⁻¹) and ammonium form (ΔG: −16.1 kcal mol⁻¹) (Fig. 2e). Based on the free energy change, protonation of the –NH of the {–P(O)–NH–} is more favorable than the oxygen of the P=O unit.⁶⁶ The charge calculation of P, N, and O from a natural bond orbital (NBO) analysis implies that as the oxide form converts to imine form, the charges on P and N are continuously reduced, whereas the charges on O reduced to iminium form then raised up significantly implying a high degree of delocalization of charges along the N–P–N bond compared to the P=O bond. After protonation of the oxygen to give rise to the oxonium form, both the P=O and P–N bonds depict a partial double bond character, which enhances P=O while making the P–N bond shorter (Fig. 2f). In the ammonium form, both P=O and P–N bond distances shrink more compared to the oxonium form. This could be due to the higher stability of the ammonium form over the oxonium, as evident from the free energy value and NBO analysis (Fig. 2g). A natural population analysis (NPA) was carried out to understand the role of the phosphamide moiety in the proton relay. From the NPA analysis, it is clearly evident that due to protonation on either of the Brønsted base sites, i.e., nitrogen center or oxygen center, the electron density is partly reduced. It could also be explained by the lowering of energies for ammonium and oxonium forms by 1.06 eV and 0.85 eV, owing to the delocalization of charge along the O–P–N moiety (Fig. 2g). However, the emergence of long-range intermolecular hydrogen bonding networks between the protonated phosphamide and hydronium ion allows the flow of protons throughout the polymeric network. However, the positive free energy change from the oxonium/ammonium to iminium and imine form is attributed to the loss of thermodynamic stability owing to the formation of the P=N bond over the P=O bond. Therefore, proton hopping within the {–P(O)–NH–} unit to form ammonium or oxonium is most energetically favorable while hopping from one {–P(O)–NH–} unit on POP to another one {–P(O)–NH–} unit, which is facilitated by the water network or channel.

Transparent proton-exchange membrane fabrication

Owing to the facile proton conductivity of PPA-2, transparent and semi-permeable membranes were fabricated to demonstrate facile proton exchange across the membrane. Polyvinyl alcohol (PVA) and polymethyl methacrylate (PMMA) were used as the support material to fabricate the membranes (Fig. 3a, S16 and S17†).^67,68 FE-SEM images of the PPA-2/PVA membrane revealed a smooth surface morphology (Fig. 3b, S18 and 19†), whereas the EDX-elemental mapping showed a uniform distribution of C, N, O, and P over a large area of the membrane (Fig. 3b inset and S20†). The PPA-2/PMMA membrane also showed a similarly smooth surface and homogeneous distribution of PPA-2 in the PMMA matrix (Fig. S21–S23†). The mechanical properties, such as tensile strength, elongation at break, and Young's modulus, were measured for the membranes. From the stress–strain plot, in the dry state for both PVA and PMMA-based membranes, a tensile strength of 10.9 MPa and 10.2 MPa was recorded, respectively, whereas the Young's moduli showed a value of 0.5 GPa and 4.6 GPa, alongside 83% and 5% greater lengthening in the elongation break being observed (Fig. S24 and Table S4†). The mechanical tests of films revealed a low stiffness for PPA-2/PVA films, whereas PPA-2/PMMA films demonstrated brittle behaviour. Water swelling tests such as water uptake (WU) and swelling ratio (SR) play a crucial impact on the performance of PEMFCs because absorption of water by membrane facilitates proton conductivity i.e. enabling of proton transport via the Grotthuss mechanism, in which the O–H bond migrates by recurrent forming and breaking of bonds instead of drifted by carrier species via the vehicle mechanism. However, an excess quantity of water uptake results in swelling of the membrane and diminishing of the stability, which leads to the rupture of the membranes. The PPA-2 functionalized PVA and PMMA demonstrated a significantly higher water uptake and lower swelling ratio, implying superior dimensional stability of polyphosphamide-based membranes. In real-life implementation of PEMFCs, the incomplete reduction of oxygen generates free radicals such as hydroperoxyl (˙OOH) and hydroxyl (˙OH) that make PEMs more susceptible to chemical degradation.⁶⁹ Thus, the oxidative stability of PPA functionalized PVA and PMMA membranes was investigated using Fenton's reagent (3% aqueous H₂O₂, 2 ppm of FeSO₄ solution) test to demonstrate peroxide and oxidative resistance of the membranes. The durability testing of membranes using Fenton's solution was performed over 24 h at 80 °C with an interval of 8 h (Fig. S25a–h†).⁶⁹ The comparison of FESEM images of membranes after 0 h, 8 h, 16 h, and 24 h revealed a partial change in the surface morphology of the membranes after 24 h. The pristine membranes demonstrated a flat, smooth morphology without any cracks, whereas cross-sectional imaging after 12 h revealed the appearance of minor pinholes. For PPA-2/PVA membranes, these cracks become prominent after 24 h, whereas for PPA-2/PMMA membranes, the film remained mostly intact after exposure to Fenton's solution for a prolonged time.

Fig. 3 (a) 1 wt% PPA-2 incorporated PVA membrane (PPA-2/PVA) under day light (inset: membrane under 365 nm UV light), (b) FESEM image of a selected area of PPA-2/PVA (inset: corresponding FESEM-EDX elemental mapping for C, N, O, and P), (c) schematic diagram demonstrating proton permeability through the PPA-2/PVA membrane in a two compartment set up consisting of pH 0.9 solution in compartment A and pH 7.5 in compartment B. Plot of pH in compartment B with time during the proton permeability experiment using the (d) PPA-2/PVA membrane and (e) PPA-2/PMMA membrane.

To demonstrate the proton exchange across the membranes, the proton (H⁺) permeability was investigated in a customized two-compartment H-type diffusion cell (with a minimum osmotic pressure effect) using PPA-2/PVA, PPA-2/PMMA films as well as control PVA/H₂0, PMMA/DMF, and Nafion 117 membranes (Fig. S26†).^70,71 The membranes were placed as a separator between two solutions containing a pH 0.9 aqueous solution (compartment A: 0.5 M H₂SO₄ solution) and freshwater of pH 7.5 (compartment B) (Fig. 3c). Within 1 h, the pH of compartment B increased to pH 2.9, while the pH of compartment A was 1.1, only highlighting the exchange of protons across the PPA-2/PVA membrane (Fig. 3d and e). Under similar experimental conditions, when the PPA-2/PMMA membrane was used, the pH of compartment B dropped to 3.9 after 1 h, highlighting a comparatively poor proton exchange of the PPA-2/PMMA membrane (Table S5†), which could be due to the hydrophobicity of PMMA. The relatively more hydrophilic nature of the PVA structure preferred faster proton exchange. Under similar conditions, using Nafion 117 film as the separator, the pH value of compartment A dropped to 2.3. Notably, under the same experimental set-ups, the PVA and PMMA membranes without the PPA-2 polymer showed almost negligible proton transfer across the membrane. The pH of compartment B dropped from 7.5 to 6.9 over a period of 1 h. Hence, the PPA-2 polymer is an integral part of the PPA-2/PVA and PPA-2/PMMA membranes for the diffusion of a proton (Fig. 3d and e). The results of the proton permeability experiment imply that Nafion 117 exhibited comparable proton permeability with PPA-2/PVA. Notably, in the commercially obtained Nafion 117 membrane, a significantly high Nafion loading was used, whereas 1 wt% PPA-2 was loaded the into PPA-2/PVA film and the proton permeability of both PPA-2/PVA and Nafion is almost comparable. It is therefore worth mentioning that the increase in the loading of PPA-2 from 1 wt% to a higher percentage will exceed the proton permeability of commercially available Nafion membranes. Notably, the smooth flow of protons with the phosphamide incorporated membrane establishes the crucial role of long-range cooperative interaction through the formation of hydrogen-bonded networks via a synchronized orientation of the polyphosphamide units within the membranes.⁷² However, the PPA-2/PVA membrane was further characterized after the proton permeability experiments, and almost comparable surface morphology and elemental composition were found, implying a recyclability of the membrane (Fig. S26–S28†).

Electrocatalytic hydrogen evolution reaction (HER) study

In the past few years, redox-active organic polymers or framework materials have been used as electrode materials to perform water-splitting reactions. In this context, polyphosphamide materials studied herein were fabricated on carbon cloth (CC) with an active surface area of 1 cm², which was then used as the working electrode in a three-electrode cell configuration to study the hydrogen evolution reaction (HER) in a 0.5 M H₂SO₄ electrolyte. During the cyclic voltammetry (CV) and linear sweep voltammetry (LSV) studies, polarization curves were recorded within 0.7 to −0.7 V (vs. RHE). The loading of PPAs was systematically varied from 3 to 12 mg per cm⁻² to optimize the mass loading of the catalyst. From the LSV curve recorded with 9 mg PPA-2 loaded CC, the geometric current density (j_geom) of 207 mA cm⁻² was obtained at a potential of −0.6 V (vs. RHE), giving an overpotential of 318 mV at 10 mA cm⁻² for the HER (Fig. 4a and S29–S33†). Under similar conditions, for the PPA-1a/CC electrode, the highest current density of 98 mA cm⁻² was achieved at a potential of −0.6 V (vs. RHE), and the LSV curves provided an overpotential of 342 mV at 10 mA cm⁻² for the HER (Fig. 4b, S34–S38 and Table S6†). Conversely, PPA-1b/CC showed a significantly lower cathodic current response and required an overpotential of 526 mV to reach the j value of 10 mA cm⁻², indicating the inferior HER performance of PPA-1b. The poor performance of PPA-1b in comparison to PPA-1a and PPA-2 highlighted the importance of the phosphamide moiety in the HER (Fig. S39–S43†). This enhances proton mobility and ensures continuous HER progression. In general, the HER on any heterogeneous surface consists of two fundamental steps, i.e., Volmer and Heyrovsky steps.⁷³ Poly-phosphamides possess a redox active {P(O)–NH} moiety, and they also acted as efficient proton adsorption sites for facile proton conduction and reduction. The adsorption of a proton into the {P(O)–NH} moiety and conduction through a polymeric network play a synergistic role in the hydrogen evolution reaction (HER), accelerating the adsorption of protons to catalytically active centers. In the next step, an adsorbed proton (H⁺) combines with electrons to generate (H*), and subsequently, a Tafel–Heyrovsky step may follow to produce molecular hydrogen (H₂).⁷⁴ Efficient H⁺ adsorption on the polyphosphamide accelerates the first steps and reduces the activation energy of the Volmer step.⁷⁵ Recently, Chakraborty et al. reported a poly-phosphamide that showed notable proton reduction and CO₂ reduction.⁷⁶ The HER performance and proton hopping of PPA-2 are remarkable compared to other metal-free polymers and other metal-based materials reported in the literature (Table S7†).^77–79 The sharp increase in geometric current density and a notable reduction of overpotential strongly suggest that the functionalization of phosphamide units bolsters the electrocatalytic HER. Notably, for PPA-1a, CV cycles performed at a scan rate of 5 mA cm⁻² for a potential range of 0.7 to −0.7 V (vs. RHE) depicted prominent redox features within −0.35 V to 0.6 V in forward (cathodic) and backward (anodic) scans (Fig. 4c). Very similar redox features were observed for PPA-2 within the potential window of −0.3 V to 0.6 V (Fig. 4c), which can be attributed to the reduction of the P^V center of {P(O)–NH}.⁸⁰ A large peak-to-peak separation of the cathodic and anodic redox peaks indicated that the electron transfers are associated with some partial chemical alteration of the {P(O)–NH} unit in the structure. The redox peak, followed by a crossover point noted in the higher cathodic region, before the catalytic current also supports that electron transfer is associated with a chemical modification in poly-phosphamide. For PPA-2, an almost 118 mV anodic shift of the cathodic redox response was observed, which implies that the reduction of the phosphamide units became more favorable to PPA-2 compared to PPA-1a (Fig. 4c). Furthermore, the anodic shift of the cathodic reduction peak with increasing proton concentration indicated that proton-assisted reduction occurs at the {P(O)–NH} center in PPA-2. As the pH of the electrolyte solution varied from 0.1 M to 1.0 M H₂SO₄, a significant linear decrease in the cathodic redox response was observed, further demonstrating the role of the acidic medium for proton conductivity and HER studies (Fig. 4d).⁸¹ The DFT studies described above also validate that protonation at the oxygen of P(O) is energetically favourable. Therefore, the CV studies demonstrate that the O/NH centers are the proton adsorption sites, while the {P^V} centers of the phosphamide units of PPA-2 and PPA-1b are the proton reduction sites during the HER.

Fig. 4 (a) LSV polarization curves for different mass loadings of PPA-2 (in 0.5 M H₂SO₄, scan rate of 1 mV s⁻¹), (b) comparison of LSV polarization curves for bare CC, PPA-2@9, PPA-1a@9, and PPA-1b@12 (in 0.5 M H₂SO₄, at a scan rate of 1 mV s⁻¹), (c) CV cycle recorded with PPA-2 (red curve) and PPA-1a (blue curve): (scan rate: 5 mV s⁻¹) (inset: selected region of the CV cycle), (d) CV cycles recorded at different H₂SO₄ conc. (0.1 M, 0.25 M, 0.5 M, and 1.0 M H₂SO₄) in the electrolyte solution (inset: plot of reductive peak potential of the P^V/IV redox feature of phosphamide versus conc. of H₂SO₄ in the electrolyte). (e) Tafel slopes, (f) C_dl values, and (g) Nyquist plots for PPA-2@9, PPA-1a@9, and PPA-1b@12. (h) CA study performed for 10 h with PPA-2/CC and PPA-1a/CC at −0.311 V and −0.351 V (vs. RHE), respectively.

The electrochemical HER performance of PPA-2 and PPA-1a over PPA-1b can also be correlated to the electrokinetics. The Tafel slope value for PPA-2 was 138 mV dec⁻¹ which is lower than the Tafel slope values of 170 mV dec⁻¹ and 186 mV dec⁻¹ for PPA-1a and PPA-1b, respectively (Fig. 4e). The electrochemically active surface area (ECSA) gives a direct measure of the number of catalytically active sites participating in the HER, and the ECSA can be experimentally determined by recording double layer capacitance (C_dl) from the CV cycles in non-faradaic regions. For PPA-2, the C_dl value was 1.19 mF cm⁻², which is ∼1.5 fold higher compared to the C_dl value of 0.87 mF cm⁻² for PPA-1a and ∼3 fold higher than PPA-1b (0.43 mF cm⁻²), respectively (Fig. 4f).^42,82 The presence of phosphamide subunits in PPA-2 and PPA-1a makes the surface of the polymeric matrix more polar, which facilitates the formation of an electrical double layer at the surface.⁸³ The C_dl data correlate that PPA-2 possesses the highest ECSA among the three polymers tested herein. The EIS studies and Nyquist plots over a broad frequency range were conducted to correlate the HER activity trend with the charge transfer resistance (R_ct). An equivalent of Randles circuit fitting (R∥C) of the Nyquist plot depicted a minimum R_ct value of 3.9 Ω and 5.8 Ω for PPA-2 and PPA-1a, respectively, while the R_ct value for PPA-1b was 10.4 Ω, indicating a better charge transfer across the electrode–electrolyte interface for the PPA-2/CC electrode (Fig. 4g). The notable HER activity of PPA-2 can be attributed to its low Tafel slope and charge-transfer resistance, indicating a facile charge transfer. PPA-2's phosphamide units within the porous channel and/or its framework contribute significantly to its superior HER performance. The long-term HER performance and stability of PPA-2/CC were explored by the chronoamperometric (CA)-HER study, followed by post-catalytic characterization. The PPA-2/CC cathode during the CA-HER at −0.311 V vs. RHE provided a stable current of 10 mA cm⁻² for 10 h (Fig. 4h). During the CA-study at −0.351 V vs. RHE, PPA-1a also demonstrated a stable current at 10 mA cm⁻² over a period of 10 h (Fig. 4h). The post-catalytic spectroscopic and microscopic characterization studies of PPA-2/CC and PPA-1a/CC indicated negligible alteration of elemental composition and structural integrity (Fig. S44–S51†).

Conclusions

Although Nafions are the best materials for PEM design, the fluorine-related environmental concern and multi-step synthetic routes limit their use in designing cost-effective PEMFCs. The key objective of this study is to find a safe and suitable alternative to perfluoro/sulfonated polymers with a high proton conduction rate and long-term stability in highly acidic pH and at high temperature. In this context, two porous polyphosphamides (PPA-1a and PPA-2) and one non-phosphamide polymer (PPA-1b) were prepared through a single-step synthetic protocol adopted by choosing a flexible organo-linker. Despite the presence of tripodal polyamine in PPA-1a and PPA-1b, both of them appeared as non-porous polymers. Conversely, the presence of a diamine organo-linker and phosphamide repeating unit in PPA-2 resulted in a higher degree of porosity with a BET surface area (SA_BET) of 5.173 m² g⁻¹. The ³¹P and ¹³C CPMAS and FESEM-EDX mapping and studies ensured the presence of the {P(O)–NH} repeating unit and less degree of branching in PPA-2. Notably, the introduction of porosity and the presence of the phosphamide unit in PPA-2 facilitate the proton conductivity many-fold higher than that of PPA-1a, PPA-1b, and some commercial Nafions. The facile proton conduction in PPA-2 was due to the low activation energy (0.12 eV), implying a fast proton hopping following the Grotthuss mechanism. The DFT study emphasized the crucial role of the {P(O)–NH} moiety in energetically favourable adsorption of protons in the “–P=O” and “–NH”, giving rise to stable tautomeric forms. The proton conductivity of PPA-2 remained unaltered when transparent PEMs were fabricated using 1 wt% PPA-2 and PMMA and PVA as base materials with an equilibrium time of only 1 h. Microscopic analyses of the PPA-2/PVA film after the proton-permeation study depicted the stability of PPA-2 in strong acidic pH. Owing to the high porosity, superior proton adsorption and conduction ability, and presence of redox-active P^V, PPA-2 was utilized as a metal-free cathode material for the HER study. The low overpotential, Tafel slope, charge-transfer resistance, and high double-layer capacitance of PPA-2 compared to PPA-1a, PPA-1b, and some reported organic polymers establish the new role of poly-phosphamide as an organo-electrocatalyst for sustainable hydrogen production. In this study, poly-phosphamides are shown as potent non-fluoro/sulfonated PEM materials for high-temperature PEMFC applications. The presence of alkyl and aromatic groups in the poly-phosphamides will allow post-synthetic modifications, which will fine-tune the electronic structure and finally can improve the proton-conductivity and other electrocatalytic performances. Therefore, poly-phosphamides are shown here as a potential alternative to Nafions and a reliable organic material for energy conversion or storage applications. This study opens up the possibility to design more superior structural variants of poly-phosphamides, which can be the ideal material to fabricate PEMs to test multiple electrochemical membrane applications as an alternative to Nafions or other perfluorosulfonated polymers.

Data availability

All the relevant experimental data, comparison tables, and Cartesian coordinates for theoretical calculations supporting this article have been included in the ESI.†

Conflicts of interest

The authors declare no competing financial interests.

Acknowledgements

AM (PMRF ID: 1403193 & ORCID id: 0009-0002-6272-5915) thanks Prime Minister Research Fellowship, India. LM (ORCID id: 0000-0002-8175-0131) thanks, IIT Delhi, for his institute fellowship. BC and SB sincerely acknowledge the Science and Engineering Board (SERB) (now ANRF), India, for the startup research grants, SRG/2021/000079 and SRG/2022/000866, respectively. The authors also acknowledge CRF-IIT Delhi for providing its research facilities for the characterization of the materials.

References

1. C. Gunathilake, I. Soliman, D. Panthi, P. Tandler, O. Fatani, N. A. Ghulamullah, D. Marasinghe, M. Farhath, T. Madhujith, K. Conrad, Y. Du and M. Jaroniec, Chem. Soc. Rev., 2024, 53, 10900–10969.
2. Z. Rao and S. Wang, Renew. Sustain. Energy Rev., 2011, 15, 4554–4571.
3. M. E. Scofield, H. Liu and S. S. Wong, Chem. Soc. Rev., 2015, 44, 5836–5860.
4. C. R. Wang, J. M. Stansberry, R. Mukundan, H. M. J. Chang, D. Kulkarni, A. M. Park, A. B. Plymill, N. M. Firas, C. P. Liu, J. T. Lang, J. K. Lee, N. E. Tolouei, Y. Morimoto, C. H. Wang, G. Zhu, J. Brouwer, P. Atanassov, C. B. Capuano, C. Mittelsteadt, X. Peng and I. V. Zenyuk, Chem. Rev., 2025, 125, 1257–1302.
5. W. Fang, W. Guo, R. Lu, Y. Yan, X. Liu, D. Wu, F. M. Li, Y. Zhou, C. He, C. Xia, H. Niu, S. Wang, Y. Liu, Y. Mao, C. Zhang, B. You, Y. Pang, L. Duan, X. Yang, F. Song, T. Zhai, G. Wang, X. Guo, B. Tan, T. Yao, Z. Wang and B. Y. Xia, Nature, 2024, 626, 86–91.
6. S. Subhadarshini, J. S. Sravan, O. Sarkar, S. Venkata Mohan, T. K. Roy and T. Jana, ACS Appl. Energy Mater., 2023, 6, 1422–1438.
7. G. Zhao, H. Zhao, X. Zhuang, L. Shi, B. Cheng, X. Xu and Y. Yin, J. Mater. Chem. A, 2021, 9, 3729–3766.
8. M. Segale, T. Seadira, R. Sigwadi, T. Mokrani and G. Summers, Mater. Adv., 2024, 5, 7979–8006.
9. K. D. Kreuer, S. J. Paddison, E. Spohr and M. Schuster, Chem. Rev., 2004, 104, 4637–4678.
10. P. Bhanja, A. Palui, S. Chatterjee, Y. V. Kaneti, J. Na, Y. Sugahara, A. Bhaumik and Y. Yamauchi, ACS Sustain. Chem. Eng., 2020, 8, 2423–2432.
11. S. Lv, X. Li, M. Lu, T. Lu, W. Lv, W. Liu, X. Dong, Z. Liu and B. Yang, J. Mater. Chem. A, 2024, 12, 9345–9370.
12. D. W. Lim and H. Kitagawa, Chem. Soc. Rev., 2021, 50, 6349–6368.
13. R. Sahoo, S. Mondal, S. C. Pal, D. Mukherjee and M. C. Das, Adv. Energy Mater., 2021, 11, 2102300.
14. X. Meng, H. N. Wang, S. Y. Song and H. J. Zhang, Chem. Soc. Rev., 2017, 46, 464–480.
15. H. Ma, B. Liu, B. Li, L. Zhang, Y. G. Li, H. Q. Tan, H. Y. Zang and G. Zhu, J. Am. Chem. Soc., 2016, 138, 5897–5903.
16. G. Jiang, W. Zou, Z. Ou, L. Zhang, W. Zhang, X. Wang, H. Song, Z. Cui, Z. Liang and L. Du, Angew. Chem., Int. Ed., 2022, 61, e202208086.
17. W. Zou, G. Jiang, W. Zhang, L. Zhang, Z. Cui, H. Song, Z. Liang and L. Du, Adv. Funct. Mater., 2023, 33, 2213642.
18. S. Feng, G. Wang, H. Zhang and J. Pang, J. Mater. Chem. A, 2015, 3, 12698–12708.
19. D. Marani, A. D'Epifanio, E. Traversa, M. Miyayama and S. Licoccia, Chem. Mater., 2010, 22, 1126–1133.
20. N. Yokota, H. Ono, J. Miyake, E. Nishino, K. Asazawa, M. Watanabe and K. Miyatake, ACS Appl. Mater. Interfaces, 2014, 6, 17044–17052.
21. F. Liu and K. Miyatake, J. Mater. Chem. A, 2022, 10, 7660–7667.
22. S. Rogalsky, J. F. Bardeau, S. Makhno, O. Tarasyuk, N. Babkina, T. Cherniavska, M. Filonenko and K. Fatyeyeva, Mater. Today Chem., 2021, 20, 100453.
23. K. Jiao, J. Xuan, Q. Du, Z. Bao, B. Xie, B. Wang, Y. Zhao, L. Fan, H. Wang, Z. Hou, S. Huo, N. P. Brandon, Y. Yin and M. D. Guiver, Nature, 2021, 595, 361–369.
24. P. Khomein, W. Ketelaars, T. Lap and G. Liu, Renew. Sustain. Energy Rev., 2021, 137, 110471.
25. L. Zhang, M. Liu, D. Zhu, M. Tang, T. Zhu, C. Gao, F. Huang and L. Xue, Nat. Commun., 2024, 15, 3409.
26. S. Bureekaew, S. Horike, M. Higuchi, M. Mizuno, T. Kawamura, D. Tanaka, N. Yanai and S. Kitagawa, Nat. Mater., 2009, 8, 831–836.
27. T. R. Farhat and P. T. Hammond, Adv. Funct. Mater., 2005, 15, 945–954.
28. Y. Li, L. Liang, C. Liu, Y. Li, W. Xing and J. Sun, Adv. Mater., 2018, 30, 1707146.
29. B. Yin, Y. Wu, C. Liu, P. Wang, L. Wang and G. Sun, J. Mater. Chem. A, 2021, 9, 3605–3615.
30. M. Chaudhary, S. Kumar, L. Mallick and B. Chakraborty, ACS Appl. Polym. Mater., 2023, 5, 8235–8240.
31. J. Zhang, S. Lu, H. Zhu, K. Chen, Y. Xiang, J. Liu, M. Forsyth and S. P. Jiang, RSC Adv., 2016, 6, 86575–86585.
32. S. C. Pal, D. Mukherjee, Y. Oruganti, B. G. Lee, D. W. Lim, B. Pramanik, A. K. Manna and M. C. Das, J. Am. Chem. Soc., 2024, 146, 14546–14557.
33. B. Joarder, J. B. Lin, Z. Romero and G. K. H. Shimizu, J. Am. Chem. Soc., 2017, 139, 7176–7179.
34. S. Pili, S. P. Argent, C. G. Morris, P. Rought, V. García Sakai, I. P. Silverwood, T. L. Easun, M. Li, M. R. Warren, C. A. Murray, C. C. Tang, S. Yang and M. Schröder, J. Am. Chem. Soc., 2016, 138, 6352–6355.
35. T. Yamada, M. Sadakiyo and H. Kitagawa, J. Am. Chem. Soc., 2009, 131, 3144–3145.
36. G. Xu, K. Otsubo, T. Yamada, S. Sakaida and H. Kitagawa, J. Am. Chem. Soc., 2013, 135, 7438–7441.
37. B. C. Patra, S. Khilari, R. N. Manna, S. Mondal, D. Pradhan, A. Pradhan and A. Bhaumik, ACS Catal., 2017, 7, 6120–6127.
38. C. Hu, X. Yang, P. Liu, Z. Song, Y. Lv, L. Miao, M. Liu and L. Gan, J. Mater. Chem. A, 2024, 12, 11867–11874.
39. J. Bitenc, K. Pirnat, O. Lužanin and R. Dominko, Chem. Mater., 2024, 36, 1025–1040.
40. D. Joseph, J. Büsselmann, C. Harms, D. Henkensmeier, M. J. Larsen, A. Dyck, J. H. Jang, H. J. Kim and S. W. Nam, J. Membr. Sci., 2016, 520, 723–730.
41. A. Ghorai, S. Roy, S. Das, H. Komber, M. M. Ghangrekar, B. Voit and S. Banerjee, ACS Appl. Polym. Mater., 2021, 3, 4127–4138.
42. H. Coskun, A. Aljabour, P. de Luna, H. Sun, N. Nishiumi, T. Yoshida, G. Koller, M. G. Ramsey, T. Greunz, D. Stifter, M. Strobel, S. Hild, A. W. Hassel, N. S. Sariciftci, E. H. Sargent and P. Stadler, Adv. Mater., 2020, 32, 1902177.
43. N. Kumari, A. Mahata and B. Chakraborty, Inorg. Chem., 2023, 62, 7728–7737.
44. N. Kumari, A. Mahata and B. Chakraborty, Inorg. Chem., 2023, 62, 19622–19630.
45. A. Mahata, N. Kumari and B. Chakraborty, ChemCatChem, 2024, 16, e202400729.
46. M. K. Adak, L. Mallick, K. Samanta and B. Chakraborty, J. Phys. Chem. C, 2023, 127, 154–168.
47. A. Kundu, B. Kumar and B. Chakraborty, J. Phys. Chem. C, 2022, 126, 16172–16186.
48. C. N. Brodie, A. E. Owen, J. S. Kolb, M. Bühl and A. Kumar, Angew. Chem., Int. Ed., 2023, 62, e202306655.
49. N. Kumari, A. Jagadeesh, P. Galav, A. Kundu and B. Chakraborty, J. Org. Chem., 2024, 89, 15851–15863.
50. X. Cao, Y. Li, X. Li, J. Zheng, J. Gao, Y. Gao, X. Wu, Y. Zhao and Y. Yang, ACS Appl. Mater. Interfaces, 2013, 5, 11494–11497.
51. K. Kanamura, S. Shiraishi, H. Takezawa and Z. i. Takehara, Chem. Mater., 1997, 9, 1797–1804.
52. P. E. R. Blanchard, A. P. Grosvenor, R. G. Cavell and A. Mar, J. Mater. Chem., 2009, 19, 6015–6022.
53. B. G. Janesko, H. C. Fisher, M. J. Bridle and J. L. Montchamp, J. Org. Chem., 2015, 80, 10025–10032.
54. C. T. Hung, L. Duan, T. Zhao, L. Liu, Y. Xia, Y. Liu, P. Qiu, R. Wang, Z. Zhao, W. Li and D. Zhao, J. Am. Chem. Soc., 2022, 144, 6091–6099.
55. S. A. Jenekhe and X. L. Chen, Science, 1999, 283, 372–375.
56. Y. P. Zhu, Y. Liu, Y. P. Liu, T. Z. Ren, G. H. Du, T. Chen and Z. Y. Yuan, J. Mater. Chem. A, 2015, 3, 11725–11729.
57. L. Jiménez García, A. Kaltbeitzel, W. Pisula, J. S. Gutmann, M. Klapper and K. Müllen, Angew. Chem., Int. Ed., 2009, 48, 9951–9953.
58. L. Duclos, R. Chattot, L. Dubau, P. X. Thivel, G. Mandil, V. Laforest, M. Bolloli, R. Vincent and L. Svecova, Green Chem., 2020, 22, 1919–1933.
59. X. Q. Xu, L. H. Cao, Y. Yang, F. Zhao, X. T. Bai and S. Q. Zang, ACS Appl. Mater. Interfaces, 2021, 13, 56566–56574.
60. S. J. Peighambardoust, S. Rowshanzamir and M. Amjadi, Int. J. Hydrogen Energy, 2010, 35, 9349–9384.
61. C. Hong, D. W. Kang, J. Song, K. Lee, H. Lee, J. Kim, H. Lee and J. Kim, J. Mater. Chem. A, 2017, 5, 17492–17498.
62. H. Lin, F. Zhou, C. P. Liu and V. Ozoliņš, J. Mater. Chem. A, 2014, 2, 12280–12288.
63. D. W. Kang, K. A. Lee, M. Kang, J. M. Kim, M. Moon, J. H. Choe, H. Kim, D. W. Kim, J. Y. Kim and C. S. Hong, J. Mater. Chem. A, 2020, 8, 1147–1153.
64. S. S. Park, A. J. Rieth, C. H. Hendon and M. Dincă, J. Am. Chem. Soc., 2018, 140, 2016–2019.
65. G. A. Ludueña, T. D. Kühne and D. Sebastiani, Chem. Mater., 2011, 23, 1424–1429.
66. L. Wang, N. Deng, G. Wang, J. Ju, B. Cheng and W. Kang, ACS Appl. Mater. Interfaces, 2019, 11, 39979–39990.
67. R. Rudra, V. Kumar and P. P. Kundu, RSC Adv., 2015, 5, 83436–83447.
68. X. Yan, X. Xiao, C. Au, S. Mathur, L. Huang, Y. Wang, Z. Zhang, Z. Zhu, M. J. Kipper, J. Tang and J. Chen, J. Mater. Chem. A, 2021, 9, 21659–21684.
69. S. Roy and S. Banerjee, ACS Appl. Polym. Mater., 2024, 6, 3456–3469.
70. J. Peron, A. Mani, X. Zhao, D. Edwards, M. Adachi, T. Soboleva, Z. Shi, Z. Xie, T. Navessin and S. Holdcroft, J. Membr. Sci., 2010, 356, 44–51.
71. X. Che, W. Tang, J. Dong, D. Aili and J. Yang, Sci. China Mater., 2022, 65, 683–694.
72. L. Martinez Crespo, G. F. S. Whitehead, I. J. Vitórica Yrezábal and S. J. Webb, Chem. Sci., 2024, 15, 17120–17127.
73. X. M. Li, L. Z. Dong, S. L. Li, G. Xu, J. Liu, F. M. Zhang, L. S. Lu and Y. Q. Lan, ACS Energy Lett., 2017, 2, 2313–2318.
74. Q. Tang and D. e. Jiang, ACS Catal., 2016, 6, 4953–4961.
75. Y. Ye, L. Zhang, Q. Peng, G. S. Wang, Y. Shen, Z. Li, L. Wang, X. Ma, Q. H. Chen, Z. Zhang and S. Xiang, J. Am. Chem. Soc., 2015, 137, 913–918.
76. M. Chaudhary, M. K. Adak, D. Gupta, N. Kumari, A. Kundu, H. K. Basak, T. Karmakar and B. Chakraborty, Inorg. Chem., 2024, 63, 494–507.
77. H. Xu, S. Tao and D. Jiang, Nat. Mater., 2016, 15, 722–726.
78. P. Ramaswamy, N. E. Wong, B. S. Gelfand and G. K. H. Shimizu, J. Am. Chem. Soc., 2015, 137, 7640–7643.
79. K. Feng, L. Liu, B. Tang, N. Li and P. Wu, ACS Appl. Mater. Interfaces, 2016, 8, 11516–11525.
80. J. S. Elias, C. Costentin and D. G. Nocera, J. Am. Chem. Soc., 2018, 140, 13711–13718.
81. A. Goyal and M. T. M. Koper, Angew. Chem., Int. Ed., 2021, 60, 13452–13462.
82. B. C. Patra, S. Khilari, L. Satyanarayana, D. Pradhan and A. Bhaumik, Chem. Commun., 2016, 52, 7592–7595.
83. M. Azimzadeh Sani, N. G. Pavlopoulos, S. Pezzotti, A. Serva, P. Cignoni, J. Linnemann, M. Salanne, M. P. Gaigeot and K. Tschulik, Angew. Chem., Int. Ed., 2022, 61, e202112679.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d5ta01738f

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