Towards Medium-Temperature Hydrogen Fuel Cells with Glassy Proton-Conductive Membranes—Part I: Fundamentals and Single-Anion Matrices

Abstract

The accelerated deployment of hydrogen technologies is widely discussed as a pathway to mitigate climate change and reduce environmental pollution associated with fossil fuel use. In this context, intermediate-temperature proton-exchange membranes that operate in the 120–200 °C window, similar to the one characterizing liquid-acid PAFC systems (much larger in their power range), are sought as a bridge between low-temperature PFSA-based PEMFCs and low-temperature PCFs, thus combining reduced sensitivity to external humidification with solid-electrolyte handling. This Part I review surveys phosphate- and silicate-based glassy proton conductors as single-anion baseline matrices and organizes the literature around a mechanistic screening framework that links processing fingerprints—particularly sol–gel hydrolysis/condensation conditions, aging, drying, and thermal treatment—to pore architecture, hydration state, and the dominant proton-transport regime. Across both families, conductivity is governed by coupled variables: network chemistry (acidic site density and connectivity), water activity (RH), and microstructure-controlled percolation and retention. Reported σ values can arise from fundamentally different regimes, ranging from hopping-dominated transport supported by dense hydrogen-bond networks and proton-bearing groups to carrier-assisted, water-mediated transport in connected porosity, with distinct humidity dependence and stability implications. Accordingly, the review treats σ(T,RH) and activation energy together with hydration/porosity indicators as primary screening metrics, and it records missing durability and device-level information—chemical stability (hydrolysis and leaching/acid migration), mechanical robustness and cycling response, and current/power density where available—as explicit knowledge gaps. While substantial progress has been achieved within single-anion phosphate and silicate glasses, particularly through engineered acidity and microstructural control, most systems remain limited by hydration drift under gradients, thermal/humidity cycling stability, and electrode/electrolyte interfacial constraints when evaluated against intermediate-temperature membrane requirements. These conclusions establish a quantitative baseline and comparison rules for Part II, which will assess mixed-network, composite, and hybrid strategies designed to decouple conductivity from water-retention and durability trade-offs.

1. Introduction

The accelerated deployment of hydrogen technologies is widely discussed as one pathway to mitigate climate change and reduce environmental pollution associated with fossil fuel use [1,2]. Although a large fraction of today’s hydrogen production still carries a substantial carbon footprint [3], rapid progress in low-carbon routes—most notably renewable-powered water electrolysis [4], photolysis-driven processes [5], and selected biotechnological pathways [6]—is improving the feasibility of fuel cells for distributed and mobile power generation. Consequently, multiple fuel-cell concepts are being evaluated for applications that differ in scale and in tolerance to fuel impurities.

Polymer electrolyte membrane fuel cells (PEMFCs) are the industrial benchmark for transport applications because they operate at low temperatures and offer high power densities [7]. However, their stringent requirements for hydrogen purity [8] complicate deployment when hydrogen is sourced from bioprocesses and/or on-site reforming of natural gas or liquid fuels, where residual contaminants such as CO and H₂S may be present [9,10,11,12]. Increasing the operating temperature above ~100 °C can mitigate catalyst poisoning effects and relax fuel-purity constraints, thereby reducing balance-of-plant complexity and overall system cost [11,12]. This motivates the development of electrolytes that enable stable operation in the intermediate-temperature window, typically targeted here as 120–200 °C, with reduced sensitivity to external humidification. Operating at elevated temperatures also offers ancillary advantages, including faster electrode kinetics [13,14] and simplified heat management—particularly relevant for combined heat and power (CHP) as well as co-/tri-generation concepts.

Conventional low-temperature PEMFC stacks rely predominantly on perfluorosulfonic acid (PFSA) membranes, with Nafion^® as the canonical benchmark. In PFSAs, a hydrophobic fluorinated backbone and hydrophilic sulfonic acid side chains phase-separate into hydrated ionic domains that form percolating proton-transport pathways under sufficient water activity. The exact performance of these materials is dependent on both their hydration level and on the details of their chemical–structural balancing of the mechanical and chemical robustness provided by the hydrophobic backbone and conductivity related to the water channels formed by the hydrophilic sulfonated side moieties [15]. Due to this fact, the so-called low-molecular-equivalent-weight PFSA-based analogs [16,17] of Nafion^® can outperform it, reaching 851.76 mW·cm⁻² versus 635.99 mW·cm⁻², producing Nafion^® characteristics in the same operational conditions [18] but with some compromise to the oxidative stability of the polymer. At temperatures exceeding ~100 °C, dehydration disrupts this hydrated network [19], leading to a sharp decrease in proton conductivity and an increased reliance on humidification hardware, which adds complexity and cost. The other drawbacks of PFSA membranes are related to their high cost [20] and growing concerns on the environmental hazards [21], leading to the withdrawal of fluorinated materials. These limitations have intensified interest in alternative electrolytes designed for operation at 120–200 °C with improved water retention and durability. These include Nafion™—zirconium phosphate composites that, according to the report of Yang and co-workers [22], are able to operate up to 150 °C. According to the literature reports, the fuel-cell-applicable polymer–inorganic filler composites may be based not only on the Nafion™ fluoropolymeric matrix [23] but also on PBI one [24] forming PBI-based nanocomposites, reaching 0.12 S·cm⁻¹ at 180 °C [25,26], as well as, sPEEK [27,28,29] =reaching 1.55 × 10⁻² S·cm⁻¹ [30] and a power density of 482.08 mW·cm⁻² [31] (both at 80 °C and under 100% RH), sPS [32], sPPO [33] and silane [34] matrices. A more detailed review of this topic can be found in the paper delivered by Kim et al. [35].

An alternative to organic matrix-based materials worth considering can be found in inorganic amorphous proton conductors. Their main advantages against the former are related to their extreme resilience against oxidative degradation occurring at high operational voltages of the resulting cell (an idle run near an OCV value is the typical case here), as well as to their inherent ability to provide a reasonable (in terms of the system startup) conductivity in a dehumidified state. On the other hand, they lack the flexibility of the organic polymers and thus are prone to cracking and other mechanical failures. It is worth stressing that this discrepancy is at least partially solved these days and flexible inorganic systems are reported in the literature [36,37].

Part I of this two-part review consolidates structure–transport relationships in phosphate- and silicate-based glassy proton conductors. The central aim is to connect composition and processing—especially sol–gel-derived microstructural “fingerprints”—to pore architecture, hydration state, and the dominant proton-transport regime. This framework is used to extract screening criteria and practical design rules, establishing a quantitative baseline for Part II, which will address composite and hybrid strategies intended to overcome the intrinsic trade-offs of single-anion glassy matrices.

From a mechanistic perspective, proton transport in hydrated inorganic networks can proceed via (i) a vehicle mechanism, where proton-bearing species (e.g., H₃O⁺ and acid/water complexes) diffuse through connected porosity, and/or (ii) a Grotthuss-type hopping mechanism along hydrogen-bond networks. In phosphate–silicate glasses, the balance between these regimes is governed by water activity (RH), pore connectivity and size distribution, and the density/strength of acidic sites related to network depolymerization and proton-bearing groups. Accordingly, activation energy, RH dependence, and stability under thermal/humidity cycling are treated here as primary screening metrics rather than auxiliary descriptors.

This review is structured as a two-part series. Part I focuses on phosphate and silicate glassy matrices as the baseline platform, with emphasis on proton-transport mechanisms (vehicle vs. Grotthuss), structure–property relations (network connectivity, NBO/–OH populations, and hydration), and engineering constraints emerging at 120–200 °C.

Part II builds on this baseline and will address approaches that exceed the intrinsic limitations of single-anion matrices, including mixed phosphate–silicate networks, composites and hybrid membranes, and processing-driven microstructural control (e.g., sol–gel porosity and interfacial proton-conduction pathways). In this design, Part I provides screening criteria and quantitative benchmarks that enable coherent evaluation of the strategies discussed in Part II.

2. Definition of the Addressed Problem

2.1. Intermediate Temperature Operation: Why 120–200 °C?

The development of solid proton-conducting electrolytes suitable for operation in the targeted intermediate-temperature window remains a central challenge. While high-temperature fuel-cell technologies are widely explored [38], robust proton conductors operating in the broad 150–400 °C range remain problematic in practice [39]. From a practical point of view, the decreased fragility against CO poisoning of the platinum catalyst is achievable in temperatures around and exceeding 150–200 °C, while a further increase in the operational temperature is counterproductive due to the significant decrease in the protonic conductivity of the glassy material over the course of its dehydration. Therefore, in this review, the primary focus is on glassy materials aimed at enabling stable operation around 120–200 °C, understood as the range related to the final system heat-up of 120–150 °C and its real operational conditions of 150–200 °C, with only marginal attention paid to other inorganic proton conductors such as ceramic [40] or mesoporous [41] ones.

Within the broader family of protonic conductors, perovskite-based ceramics are prominent in intermediate- to high-temperature applications [42,43]. However, their operating temperatures often exceed the “medium” range targeted here, and they can face additional challenges such as CO₂ sensitivity and processing/handling complexity [44,45,46].

2.2. Reference Point: PAFCs as Mature Intermediate-Temperature Technology

The intermediate-temperature operating range is not purely hypothetical: it has long been demonstrated in phosphoric acid fuel cells (PAFCs) [47,48]. In PAFCs, liquid phosphoric acid (with the simplified formula H₃PO₄—see [49,50] for details of the complicated structural chemistry of this compound in elevated temperatures) functions as the electrolyte, enabling operation above the conventional low-temperature PEMFC regime.

PAFCs with immobilized liquid electrolyte in an inert matrix were developed in the late 1960s and remain commercially deployed, with ongoing engineering improvements aimed at performance and durability [51,52]. Despite their maturity and the feasibility of tailoring such systems to different load profiles [53], PAFCs are still widely regarded as a niche technology [54]. Such architectures can, due to the fact of their relative robustness of design and ability to work with low-quality hydrogen (up to ~1 v/v% of carbon monoxide), be easily extended further by incorporating fuel-processing units, enabling the use of practical fuels such as natural gas [55,56], biogas [57,58,59], and methanol [60,61], rather than high-purity hydrogen alone.

Today, industrial-grade systems are offered by a limited number of manufacturers (e.g., Fuji Electric and Doosan), with typical product lines spanning the ~100–440 kW class and targeting stationary generation for small-to-medium industrial and public buildings and residential districts.

2.3. Case Studies and System-Level Value Proposition

To illustrate the system-level value proposition, PAFC-based deployments have been reported in multiple building-scale contexts, including: (i) a residential area installation using a ~200 kW PC25C unit with a condensing boiler for local heating; (ii) a hospital-scale hybrid system combining a ~200 kW fuel cell with internal combustion engines, where the fuel-cell heat stream supported building air conditioning; and (iii) a hotel installation using a ~100 kW class unit operated in grid-interactive mode, where low- and higher-temperature waste-heat streams were utilized for water pre-heating and absorption-based cooling, respectively [62,63]. Across these application categories, fuel-cell generators are commonly associated with improved supply reliability, reduced need for grid reinforcements, compensation for renewable intermittency, and coverage of peak demand.

Beyond building-scale CHP, PAFC-derived architectures have also been proposed in coupled energy systems where solar heat supports endothermic fuel processing (e.g., natural-gas reforming), improving overall energy efficiency relative to reforming schemes that rely on partial fuel combustion for process heat [64]. Related concepts include direct gaseous [56] or liquid [65] hydrocarbon-fueled cell designs that build on similar system thinking.

A distinct application niche that motivates proton glass-based intermediate-temperature fuel-cell concepts is the provision of autonomous power for distributed infrastructure elements, such as natural-gas (NG) distribution stations located away from electrical lines [66,67]. Such sites require reliable power for monitoring, control, and safety systems, while their demand can be relatively modest (typically in the few-hundred-watt range). For small installations, both internal combustion engines and typical (scaled within multi-kilowatt to hundreds-of-kilowatts range [68]) PAFC-based systems can be impractical in terms of installation and maintenance. Even if commercialization of laboratory demonstrators consisting of smaller stacks is considered, liquid-electrolyte PAFCs do not straightforwardly address the niche of compact, small-scale fuel-cell generators discussed above, and those which have been reported [69] are still in the kW, not W, power range. Consequently, compact and reliable electrochemical power sources able to be fed with lower quality hydrogen obtained locally by means of NG reforming are of interest, serving additionally as a convenient local heat supply needed to maintain positive temperatures during gas pressure reduction. In this context, high-temperature solutions such as SOFCs and protonic ceramic conductors may be considered; however, intermediate-temperature units based on polymeric or inorganic membranes could offer a more versatile and cost-effective pathway if suitable electrolytes are available.

A comparison of medium-temperature glassy-electrolyte-based fuel cells to other, similar temperature-range constructions is presented in

Table 1. Benchmark proton-conducting electrolytes for hydrogen fuel cells.

Electrolyte Class	Typical σ (S·cm−1)	Temperature Window	Humidity Sensitivity	Key Limitations/Notes
PFSA (Nafion-class)	10−2–10−1 (humid)	<100 °C	High	Dehydration above 100 °C; cost; humidification needs
PBI (acid-doped HT-PEM)	10−3–10−1	120–200 °C	Low	Acid management; durability trade-offs
PAFC (liquid H3PO4)	10−2–10−1	150–220 °C	Low	Electrolyte management; corrosion; system complexity
Glassy phosphate/silicate	10−6–10−2	120–200 °C	Medium–high	Porosity/hydration control; aging; durability benchmarks scarce

.

2.4. Gap and Positioning vs. Prior Reviews

Among the available synthesis routes for amorphous and porous inorganic electrolytes, sol–gel-derived processes are widely used because of their compositional flexibility and their ability to tailor porosity and hydration state [70]. However, reproducibility can be limited by underreported procedural details (e.g., hydrolysis/condensation conditions, aging, drying, and calcination protocols), which complicates cross-study comparison and translation of reported results.

First, instead of cataloging materials or conductivity values, this review introduces a screening-oriented framework linking processing history → microstructure → hydration state → proton-transport regime → device-relevant constraints. Accordingly, conductivity is treated as a derived parameter requiring joint interpretation with activation energy, σ (T,RH) behavior, hydration modes, and porosity descriptors. This resolves a common comparability issue, where similar σ values may originate from fundamentally different mechanisms (hopping vs. vehicle).

Second, the review formalizes a transport-regime-based comparison logic by explicitly distinguishing hopping-dominated and carrier-assisted proton transport. This distinction, evaluated through σ(T,RH), E_a, and hydration indicators, is rarely used as a primary comparative axis despite its direct implications for stability and durability.

Third, hydration and water retention are reframed as primary design variables rather than experimental conditions. This perspective clarifies regime crossover effects and stability trade-offs characteristic of intermediate-temperature electrolytes.

Fourth, processing fingerprints, particularly in sol–gel systems, are interpreted as structural determinants governing pore architecture, percolation pathways, and hydration stability, extending beyond procedural synthesis descriptions.

Among the many electrolyte chemistries investigated for intermediate-temperature fuel cells, solid-state analogs of PAFC-type phosphate systems have attracted sustained attention. Several reviews, including [39] and the more recent work by Joseph and co-workers [71], summarize phosphate-based solid proton conductors for intermediate-temperature operation. In that framework, materials are often grouped into ammonium polyphosphates, cesium phosphates, pyrophosphates, and further compositions (e.g., systems incorporating La or Zr). These surveys cover crystalline and amorphous materials as well as composites; a dedicated crystallographic perspective is provided by Hatada and co-workers [72]. Notably, despite extensive prior work, key questions remain open—particularly regarding structure–transport relations under controlled hydration, durability under cycling, and the transferability of design principles to multicomponent systems. This gap becomes even more consequential when moving from single-phase materials to composite and hybrid electrolytes. Therefore, before addressing phospho-silicate glass-based composites in Part II, Part I provides an organized baseline review of phosphate- and silicate-based glassy proton conductors and extracts screening criteria that enable coherent cross-comparison.

2.5. Screening Framework and Differences Between Part 1 and Part 2

To address the commonly noted lack of novelty in reviews on glass-based proton electrolytes, Part I formulates an explicit screening framework that goes beyond re-listing conductivity values. The goal is to convert the literature into an actionable map that links processing and network chemistry to the dominant transport regime and to device-relevant constraints, thereby enabling a coherent comparison of strategies discussed in Part II.

A necessary first step is transport-regime identification, because conductivity values are not commensurate if one system is dominated by Grotthuss-type hopping along a hydrogen-bond network while another relies on vehicle-type transport mediated by water- or acid-associated carriers. Activation energies, together with the humidity and temperature dependence of σ, should therefore be interpreted within a unified framework that distinguishes hopping-dominated conduction—typically associated with persistent hydrogen-bond connectivity and a higher density of proton-bearing groups—from carrier-assisted transport, which is more tightly coupled to water activity and the presence of mobile molecular species.

In the 120–200 °C window, hydration management becomes a design variable rather than a boundary condition. Screening must therefore resolve whether the observed conductivity is supported primarily by structurally retained hydroxyl groups or bound water (typically more stable) or by uptake of weakly held water (often higher σ but prone to dehydration under gradients). In practice, σ(RH) at a fixed temperature, interpreted alongside porosity/retention indicators where available, provides a workable proxy for this distinction.

Within the same logic, composition is discussed through network connectivity and acidic site density rather than compositional “rules of thumb”.

Finally, meaningful cross-study comparison requires that microstructure and processing fingerprints be reported with sufficient specificity. For sol–gel-derived glasses, variables such as hydrolysis/condensation conditions, aging time, drying protocol, and calcination temperature directly control pore size distribution, water retention, and percolation pathways. Conductivity data reported without these descriptors are treated here as weakly comparable rather than directly rankable.

Beyond σ and E_a, practical relevance requires engineering metrics and durability considerations, including chemical stability (e.g., hydrolysis resistance and acid migration/leaching), mechanical robustness (brittleness, crack sensitivity, and thermal cycling), and, where available, at least one device-level metric (e.g., current density or power density). When such metrics are absent, they are treated as explicit knowledge gaps rather than ignored.

Table 2. Device-relevant metrics and common reporting gaps for medium-temperature glass electrolytes.

Metric	What It Captures	Minimum Reporting	Typical Gap in the Literature	Use in Part II
σ(T,RH)	Transport performance	σ at ≥2 temperatures + humidity condition	Single-point σ without RH context	Benchmarking improvements from composites/hybrids
Ea	Mechanism proxy	Arrhenius fit range stated	Mixed regimes; inconsistent fit windows	Compare mechanistic shifts via interfaces
Water retention/−OH	Hydration stability at 120–200 °C	TGA/DSC or qualitative bound/free water statement	No hydration quantification	Targeted retention via mixed networks/porosity control
Chemical stability	Hydrolysis, leaching, acid migration	Soak/cycling test + composition check	Short tests or absent	Interfacial stabilization strategies
Mechanical/thermal cycling	Cracking, sealing, CTE mismatch	Flexural/indentation or cycling survival	Rarely reported	Composite reinforcement/lamination routes
Device metric	Practical relevance	Current density or power density at stated T/RH	Often missing	Primary comparative endpoint

summarizes the device-relevant metrics used throughout this series, the minimum reporting needed for comparison, and the typical gaps in the literature, providing a consistent baseline for the mixed-network and composite strategies evaluated in Part II.

3. Polyphosphate-Based Systems

3.1. Polyphosphoric Acids as Reference Systems (Composition and Glass Formation)

A foundational study on the conductivity of binary systems consisting of phosphoric acids and water was reported by Wang et al. [73]. Polyphosphoric acids have also been extensively investigated with respect to their physicochemical properties [49,50] and potential applications. Upon heating, the key reaction is the condensation of orthophosphoric acid molecules, yielding dimers and higher oligomers via elimination of water. An alternative route to polyphosphate formation involves reacting orthophosphoric acid with phosphorous pentoxide [74]. Attempts have been made to determine the composition of condensed phosphoric acid mixtures using analytical techniques such as chromatography. These analyses indicate that orthophosphoric acid (V), pyrophosphoric acid, and linear polyphosphoric acid compounds coexist in mixtures with P₂O₅ contents of 81–85%. For branched polyphosphoric acids, however, chromatographic quantification is less reliable because these species hydrolyze more readily than their linear counterparts. When the content of phosphorus pentoxide exceeds 85%, cyclic phosphoric acids such as trimetaphosphoric acid and tetrametaphosphoric acid are also observed [74]. Importantly, none of the constituent acids crystallize at temperatures down to −60 °C; instead, a stiff glass is formed [75].

From the perspective of intermediate-temperature membranes, these condensed acid systems provide a chemically intuitive reference point: increasing polyphosphate chain length and associated condensation chemistry changes both the density of proton-bearing groups and the balance between bound water and more weakly retained water. These variables later reappear—implicitly or explicitly—in glassy phosphate electrolytes as network depolymerization, acidic site density, and hydration state, and they must be tracked alongside σ(T,RH) to enable meaningful comparison (see Table 2).

3.2. Charge-Carrier Identification and Water-Related Species in Phosphate Glasses

The conductive and viscoelastic properties of these reference systems highlight the potential of phosphate-based proton conductors for fuel-cell applications. Early work on alkali-free proton-conducting glasses, including a report by Abe [76], emphasized that the identity of the mobile charge carriers in these oxide glasses was uncertain. To address this, Abe proposed a DC-driven intercalation experiment in which charge carriers were injected from the glass into a WO₃ layer. Proton mobility was then inferred optically from the color change of tungsten oxide upon proton intercalation. On this basis, proton conductivity was found to correlate with the concentration of -OH residues present in the glassy material.

In another contribution, Abe et al. [76] proposed a more unconventional approach to obtaining enhanced protonic conductivity in Mg(PO₃)₂-based glasses prepared by melt quenching. A proton ion beam with an average energy of 120 kV and a flux of 1018 particles per cm² of the sample’s surface was used to inject charge carriers into the structure. As a result, the conductivity increased to 5 × 10⁻⁴ S·cm⁻¹ for a system whose initial conductivity did not exceed 10⁻¹¹ S·cm⁻¹. In contrast, no comparable enhancement was observed when Ca(PO₃)₂ and SiO₂ melt-quenched materials underwent proton implantation. The source of this discrepancy was identified by spectroscopic investigation. Evidence showed that in the magnesium case, both P-OH moieties and molecular water were present in the material structure, whereas in the other materials, only X-OH (X = P or Si) units (and Si-H units in the silicate-based glass) were detected. These results indicate that molecular water in the magnesium phosphate-based system is responsible for the unique conductivity enhancement observed. The authors further argued that these observations do not originate from accidental humidification, because the presence of water molecules was accompanied by the occurrence of elemental phosphorous in its “red” allotropic form due to ion-beam-induced decomposition following 1/5P₂O + 2H* → H₂O + 2/5P.

Mechanistically, these early studies already highlight the central linkage required for cross-study synthesis: network chemistry and processing control the formation of proton-bearing groups (e.g., P-OH) and the incorporation/retention of molecular water, which in turn determines whether transport is dominated by hopping along a hydrogen-bond network or by water/acid-mediated carriers. Consequently, σ values should be read together with hydration descriptors and spectroscopic evidence of proton-bearing species whenever available (Table 2).

3.3. Modifier Effects in Melt-Derived Phosphate Glasses (Ba/Mg/Zn and Related Systems)

Following this discovery, a family of barium phosphate glasses was investigated with respect to structure and dielectric properties, and proton conductivity was determined by Desphande and Tiple [77]. After partially replacing barium with magnesium, improved proton transport was reported. Further improvement was confirmed by Sumi et al. [78] upon transforming the system into a ternary one by adding zinc oxide. After substituting barium with zinc, which has a lower electropositivity, an increase in conductivity of about two orders of magnitude (from about 10⁻⁵ S·cm⁻¹ to about 10⁻³ S·cm⁻¹) was observed at 250 °C. More detailed electrochemical tests—including the determination of the proton transport number (close to 1.0) and a H₂/O₂ fuel-cell test (giving a power density of 0.4 mW·cm⁻² at a temperature of 250 °C)—confirm improved proton transport; however, the obtained performance remains far below application-relevant expectations. A more detailed study of the conductivity behavior of this ternary system was completed by Anvari et al. [79]. It was found that the activation energy of conduction appears to be unchanged when BaO is replaced with ZnO, up to a total content of 40% in the mixed oxide-based system.

Interpreted through the screening logic of this review, these modifier-driven trends should be discussed in terms of how cation substitution perturbs network connectivity (including depolymerization and non-bridging oxygen populations) and, as a consequence, the density/strength of proton-bearing groups and hydrogen-bond connectivity. The observation that E_a can remain approximately unchanged while σ increases suggests that microstructural/hydration-related factors and carrier availability may be shifting even when the apparent activation barrier is not strongly altered, which reinforces the need to report σ(T,RH) with hydration context rather than as isolated values (Table 2).

A zinc phosphate glass with a similar composition was obtained by Mercier et al. [80] by melting a H₃PO₄-ZnO mixture at a temperature of 900 °C. The system was found to be in temperature-dependent equilibrium with water vapor. Conductivity studies showed that the surface layer of the vitreous sample exhibited increased conductivity as a result of wetting. Thermogravimetric studies further indicated that the presence of this surface layer hindered the rate of water absorption into the bulk of the material. In addition, multinuclear NMR studies showed that at 140 °C, water absorption leads to the formation of proton-transport-active species such as ortho- and pyrophosphoric acids. Therefore, long-term stability of these materials under operating conditions may be influenced by acid migration outside the membrane.

This latter point has direct device relevance and should be treated as more than a qualitative caveat: if operation relies on the formation and redistribution of mobile acid species at the surface, then the apparent σ may reflect a transient or spatially inhomogeneous state, and long-term performance may be limited by acid migration/leaching, compositional drift, and evolving interfacial resistance. As a minimum, such systems should require stability benchmarks under thermal/humidity cycling and post-test compositional checks to establish whether proton transport remains supported by a stable bulk network or by a mobile surface/near-surface phase (Table 2).

3.4. Multicomponent and Composite Phosphate Electrolytes: Routes to Enhanced σ

Another phosphate-based proton conductor was described by Raskovalov et al. [81]. For two-component V₂O₅-P₂O₅ glass, they distinguished reversible physical binding of water at around 100 °C from irreversible chemisorption at higher temperature (150 °C) and linked the latter to the observed proton conductivity. This distinction is important for screening because it separates conductivity supported by weakly held water (often higher σ but more vulnerable to gradients) from conductivity associated with more strongly retained species (typically more stable but not necessarily higher σ).

A related ternary system containing TiO₂ was reported by Melo et al. [82]. They described a proton-conducting inorganic composite comprising a vitreous (V₂O₅-P₂O₅) phase and a ceramic (TiO₂) phase, which exhibited pronounced features in electrochemical impedance spectroscopy. The authors further observed two distinguishable conductivity mechanisms, attributable to contributions from both phases. In this context, it is worth noting that TiO₂—often treated as an inert ceramic filler interacting with a conductive matrix primarily through surface groups—has also been reported to possess a mesoporous internal structure capable of supporting proton transport [83].

More complex glass-forming compositions have also been explored. Soheyli and Shoar [83] combined two transition metals (vanadium and molybdenum) with calcium and lithium, which act as counterions to the phosphate anionic network. Because the proposed system was effectively devoid of mobile protons, conductivity remained extremely low, ranging from 10⁻¹² S·cm⁻¹ at 30 °C to 10⁻⁸ S·cm⁻¹ at 110 °C. This example is instructive as a boundary case: altering composition without preserving proton-bearing group density and hydrogen-bond connectivity can yield structurally valid glasses that are irrelevant as proton electrolytes.

In another study, Suzuki et al. [84] processed the multicomponent vitreous system (36HO_1/2—4NbO_5/2—2BaO—4LaO_3/2—4GeO₂—1BO_3/2—49PO_5/2) into thin layers by high-temperature pressing. This processing yielded a conductivity reaching 10⁻³ S·cm⁻¹ at 300 °C and an order-of-magnitude decrease in ohmic losses under fuel-cell conditions compared to the unprocessed material. From an engineering standpoint, such results emphasize that processing can alter not only σ but also contact resistance and microstructural percolation pathways; therefore, “processing fingerprints” are necessary to interpret whether performance gains reflect intrinsic bulk transport, interfacial improvements, or both.

Abe et al. [85] investigated a BaO–La₂O₃–Al₂O₃–P₂O₅ multicomponent system. In the wet state, the material exhibited conductivities ranging from 10⁻³ S·cm⁻¹ at 250 °C up to 10⁻² S·cm⁻¹ at 200 °C. The authors suggested that the material is applicable in fuel cells operating in this temperature range even without additional humidification. Within the screening framework, such claims require explicit reporting of hydration conditions and stability under cycling, because “wet state” can encompass fundamentally different water activities and retention states that determine both σ and durability.

A different strategy for obtaining a protonically conductive mixed phosphate, WO₃–35NaO_1/2–8NbO_5/2–5LaO_3/2–51PO_5/2, was described by Ishiyama et al. [86]. The sodium-conducting glass was first prepared by conventional melt quenching, and protons were subsequently introduced via electrochemical Na⁺ substitution. This exchange was carried out using liquid tin and thin-film palladium electrodes in contact with a hydrogen-containing atmosphere under a 5V bias for 2 to 30 h under amperometric control. Evidence for the ionic exchange was supported by impedance spectroscopy, which showed a substantial conductivity increase across the investigated temperature range. At 150 °C and 250 °C, conductivity increased from 2 × 10⁻⁷ S·cm⁻¹ and 10⁻⁵ S·cm⁻¹ (sodium form) to 6 × 10⁻⁶ S·cm⁻¹ and 5 × 10⁻⁴ S·cm⁻¹ (protonated form), respectively. Proton concentrations were estimated from the intensity of -OH stretching vibrations observed by FT-IR spectroscopy. This approach highlights an additional design lever: decoupling glass formation from proton introduction can, in principle, tune proton-bearing group populations, but it also introduces process- and stability-related questions (e.g., homogeneity and long-term retention of exchanged species) that require explicit benchmarking.

Building on studies of related systems, Zhang et al. [87] prepared a phosphate-based glass–ceramic composite in which Ca(PO₃)₂ served as the vitreous matrix, while crystalline La(PO₃)₃ was used as the dispersed phase. In humidified air, a conductivity of up to 1.52 × 10⁻⁵ S·cm⁻¹ was registered at 550 °C. Proton transport was attributed to mobile protons acting as the dominant charge carriers, supported by infrared spectroscopy indicating the presence of P-OH moieties in the humidified material.

Amezawa et al. [88] reported related findings for strontium–lanthanum mixed phosphate materials. With a 1%mol content of Sr²⁺ in a La³⁺-based system, conductivities up to 10⁻⁴ S·cm⁻¹ were observed at 700 °C. Because this value was substantially higher than those reported previously by the same authors for similarly composed systems containing a Sr²⁺-doped LaPO₃ phase [89], the authors proposed that the enhanced conductivity reflects the presence of the doped LaP₃O₉ polyphosphate structures.

Another set of composite systems was prepared by Mosa et al. [90] using a sol–gel route, combining a P₂O₅ matrix with a HSO₃⁻-grafted TiO₂ dispersoid. The properties of the resulting materials, described as exhibiting a three-dimensional proton-conducting channel network, depended on preparation factors, including the phosphorus precursor (PCl₃ or H₃PO₄). For materials containing a non-grafted dispersed phase, conductivities of 0.11 S·cm⁻¹ at 140 °C and 80% RH (PCl₃ route) and 0.2 S·cm⁻¹ under the same conditions (H₃PO₄ route) were reported. Functionalization of the filler to introduce HSO₃⁻ acidic groups on the surface further increased conductivity, reaching 0.79 S cm⁻¹ at 140 °C and 80% RH. Structural characterization using SAXS and TEM confirmed the presence of a well-ordered mesoporous structure in these materials. Here, the conductivity gains are plausibly consistent with simultaneously (i) increased acidic site density (via HSO₃⁻ groups) and (ii) stabilized water-mediated pathways in a controlled pore network; however, screening for 120–200 °C membranes still requires durability and aging metrics to establish whether the hydrated conduction network remains stable under thermal/humidity cycling (Table 2).

Non-lanthanum-containing materials were also studied. For example, Souissi et al. [91] investigated a KH₂PO₄—Al₂O₃—P₂O₅ system and concluded that an optimal composition containing 0.1 mol of Al₂O₃ exhibits a relatively low room-temperature conductivity of 7 × 10⁻⁶ S·cm⁻¹. Spectroscopic analysis further indicated that adding the basic component Al₂O₃ depolymerizes the polyphosphoric backbone, promoting formation of Al-O-P groups at the expense of P-O-P linkages. This illustrates that depolymerization alone is not a sufficient predictor of σ unless it is linked to proton-bearing group formation, hydrogen-bond connectivity, and hydration state.

A zirconium phosphate-based system with a surprisingly high proton conductivity, reaching 10⁻² S·cm⁻¹ under ambient conditions, was reported by Abe et al. [92]. The authors also noted that the resulting structures were not hygroscopic and remained stable in atmospheric air, despite containing molecular water immobilized within the structure. A similar conductivity was also reported [93] for a composite material using the same phosphate base with a silicotungstic filler. These observations are potentially important for intermediate-temperature operation because they suggest that high σ need not always require hygroscopicity; however, stability screening must still clarify whether immobilized water and/or acid species remain retained under gradients and cycling and whether interfacial losses remain acceptable.

Phosphate-based glasses can be prepared both by thermal quenching and via low-temperature sol–gel routes. Carta et al. [94] compared glasses of identical chemical composition obtained by these two methods and found that, despite the radically different synthesis routes, the resulting internal structures were nearly identical, as confirmed by PXRD, molecular spectroscopy, and ³¹P NMR. From a screening perspective, this reinforces that process details can control microstructure and hydration, even when the average “bulk structure” appears similar, and thus should be treated as essential descriptors for cross-study comparison (

Table 3. Comparison of phosphate-based systems.

System/Composition	Synthesis Route	Porosity/BET	Hydration/RH	σ(T) and Ea	Stability/Aging Notes	Ref.
Phosphoric acid–water binary system	Liquid mixture	N/A	Water-content controlled by composition (RH not applicable as a primary control variable)	Conductivity characterized vs. composition/temperature; Ea not reported here	Serves as physicochemical baseline; device-level durability not applicable	[73]
Mg(PO3)2 glass after proton implantation	Melt quenching followed by proton ion-beam implantation (120 kV; high fluence)	Not reported	Not an RH study; conductivity linked to implanted species and retained molecular water	σ increased up to ~5 × 10−4 S·cm−1 from ≤10−11 S·cm−1; Ea not reported	Enhancement attributed to coexistence of P-OH and molecular water; durability/cycling not reported	[76]
Ba phosphate glass, partial Ba→Mg substitution (family introduction)	Glass composition tuning	Not reported	not reported	Proton conductivity reported; Ea not reported	Not discussed	[77]
Ba-Mg phosphate glass with ZnO addition	Ternary glass via compositional modification	Not reported	not reported	σ increases ~10−5→~10−3 S·cm−1 at 250 °C; transport number ~1; H2/O2 FC test power density ~0.4 mW·cm−2 at 250 °C; Ea not reported	Demonstrates protonic transport but still far from practical FC performance; durability not reported	[78]
Same ternary Ba–Zn phosphate glass family (detailed σ/Ea behavior)	Glass study of compositional substitution effects	Not reported	not reported	Ea reported as unchanged upon BaO→ZnO substitution up to 40% total mixed oxide content (numeric Ea not quoted here)	Not discussed	[79]
Zn phosphate glass from H3PO4–ZnO melt	Melt at ~900 °C (H3PO4–ZnO)	Not reported	Temperature-dependent equilibrium with water vapor; surface wetting increases σ; water uptake hindered by conductive surface layer	σ values for bulk not quoted here; formation of ortho-/pyrophosphoric acids upon water uptake at 140 °C; Ea not reported	Potential long-term issue suggested: acid migration outside membrane under operating conditions	[93]
V2O5–P2O5 binary glass (water binding regimes)	Glass	Not reported	Water binding: reversible physical around ~100 °C; irreversible chemisorption at ~150 °C	Proton conductivity attributed to chemisorbed water regime	Highlights distinct hydration states; durability not reported	[81]
(V2O5–P2O5) glass + TiO2 ceramic composite	Inorganic glass–ceramic composite	Not reported	Not reported	Two distinguishable conductivity mechanisms attributed to two phases	Not discussed	[82]
TiO2 as mesoporous proton-transport-capable filler	Filler property report	Mesoporous internal structure noted	Not reported	Not reported	Not applicable as a standalone electrolyte in this section	[83]
Multicomponent glass (36HO1/2–4NbO5/2–2BaO–4LaO3/2–4GeO2–1BO3/2–49PO5/2), hot-pressed thin layers	High-temperature pressing into thin layers	Not reported	Not reported	σ reaches ~10−3 S·cm−1 at 300 °C; Ea not reported; ohmic losses reduced by ~order of magnitude vs. unprocessed	Processing-driven improvement; long-term durability not reported	[84]
BaO–La2O3–Al2O3–P2O5 multicomponent glass	Glass	Not reported	“Wet state”; explicit RH not reported; claim of operation without additional humidification	σ ~10−3 S·cm−1 at 250 °C and up to ~10−2 S·cm−1 at 200 °C; Ea not reported	Claimed applicability in FC temperature window; durability/cycling not reported	[85]
WO3–35NaO1/2–8NbO5/2–5LaO3/2–51PO5/2 glass, Na+→H+ electrochemical exchange	Melt-quench Na+ form + electrochemical ion exchange (Sn/Pd electrodes, H2 atmosphere, 5 V bias, 2–30 h)	Not reported	Not reported	At 150 °C: 2 × 10−7→6 × 10−6 S·cm−1 (Na→H); at 250 °C: 10−5→5 × 10−4 S·cm−1; Ea not reported	Demonstrates controllable proton introduction; durability not reported	[86]
Glass–ceramic composite: Ca(PO3)2 glass matrix + La(PO3)3 dispersed phase	Glass–ceramic composite preparation	Not reported	Humidified air	σ up to 1.52 × 10−5 S·cm−1 at 550 °C; Ea not reported	Protonic nature inferred from P-OH after humidification; durability not reported	[87]
Sr-doped La-based phosphate system (polyphosphate phase attribution)	Solid-state ceramic system	Not reported	Not reported	σ up to ~10−4 S·cm−1 at 700 °C; Ea not reported	Higher σ attributed to doped LaP3O9-type polyphosphate presence; durability not reported	[88]
P2O5 matrix + TiO2 dispersoid (HSO3−-grafted), sol–gel composite with 3D channels	Sol–gel; precursor-dependent (PCl3 vs. H3PO4); optional TiO2 surface grafting	Well-ordered mesoporous structure evidenced by SAXS/TEM (BET/pore size not quoted here)	80% RH at 140 °C explicitly	At 140 °C/80% RH: 0.11 S·cm−1 (PCl3 route, non-grafted) and 0.20 S·cm−1 (H3PO4 route, non-grafted); after HSO3− grafting: stable 0.79 S·cm−1; Ea not reported	“Stable” σ at fixed T/RH stated; cycling/leaching/drift not reported	[90]
KH2PO4–Al2O3–P2O5 system (basic additive depolymerizes polyphosphate backbone)	Glass/phosphate network modification study	Not reported	Not reported	Optimal 0.1 mol Al2O3: σ ~7 × 10−6 S·cm−1 at ambient; Ea not reported	Structural change (Al-O-P vs. P-O-P) noted; durability not reported	[91]
Zirconium phosphate-based system (ambient high σ claim; non-hygroscopic)	Not specified here	Not reported	Claimed non-hygroscopic yet contains immobilized molecular water	σ up to ~10−2 S·cm−1 at ambient; Ea not reported	Claimed stability in atmospheric air; detailed durability metrics not given here	[92]
Zirconium phosphate base + silicotungstic filler composite	Composite	Not reported	Not reported	Similar σ magnitude (~10−2 S·cm−1) reported; Ea not reported	Not discussed	[94]

).

3.5. Mechanistic Constraints and Structure–Transport Correlations Across Phosphate Systems

Because substantial discrepancies are reported across studies in the charge-transport capabilities of phosphate-based materials, a closer look at the physicochemical factors governing transport is warranted. While the overall scope of this series targets medium-temperature electrolytes (120–200 °C), selected higher-temperature results are included below where they provide mechanistic context that is useful for interpreting glassy systems and for motivating the screening criteria applied in Part II. Structure–transport correlations in glassy ion conductors are commonly developed starting from simple model systems; accordingly, for phosphate-based materials, the behavior of phosphoric acid itself is an important reference. Wang et al. [95] correlated viscoelastic and dielectric properties of phosphoric acid–water mixtures to estimate the glass transition temperature. They noted that, despite massive production and extensive industrial use of phosphoric acid, its physicochemical properties remain only partially characterized. Even the glass transition temperature (T_g) is not unambiguously established: depending on the experimental method (e.g., NMR, dilatometry, DSC and conductivity studies), reported values span 150 to 250 K. The authors further stated that although proton transport in these systems is clearly associated with the Grotthuss mechanism [96], the corresponding structural relaxations remain insufficiently understood. In their study, both conductivity across the investigated temperature range and the determined T_g values depended on the H₃PO₄:H₂O ratio. Close to T_g, a change in the nature of the temperature dependence was observed: ionic transport in the vitrified glassy state followed Arrhenius-type behavior, whereas the viscoelastic response was described by VTF-type dependencies. This underscores the fact that comparing σ across studies without consistent RH/water-content information conflates fundamentally different transport regimes.

A complementary “edge case” for phosphate-based ion transport is provided by crystalline polyphosphates, which are discussed here only to clarify mechanistic limits (e.g., anisotropy and crystallographic conduction pathways) that are not directly transferable to the unoriented amorphous electrolytes emphasized in this review. Hatada et al. [72] reported properties of lanthanum polyphosphate (LaP₃O₉), describing it as a promising but not fully understood proton-conductive material. Compared with related phases such as La₇P₃O₁₈ [97], LaPO₄ [98], LaP₃O₉ [99] and LaP₅O₁₄ [100] (in all cases doped with Ca, Sr or Ba), LaP₃O₉ was reported to exhibit higher proton conductivity (5 × 10⁻³ S·cm⁻¹ at 700 °C versus 10⁻⁴ S·cm⁻¹ at 600 °C, respectively) and to retain improved transport properties with reduced dependence on the water vapor pressure. Structurally, LaP₃O₉ consists of virtually infinite helices of phosphate polyanions oriented along the c axis, whereas other proton-conductive phosphates are built from isolated monomers and dimers of tetrahedral PO₄ units. On this basis, conduction was expected to be anisotropic and preferentially aligned along the helical chains. This was experimentally supported by single-crystal measurements showing conductivity along the c axis almost an order of magnitude higher than along the ab plane. The uniaxially oriented material was subsequently used to fabricate a polycrystalline fuel-cell membrane, yielding a power density of 10 mW·cm⁻² at around 18 mA·cm⁻² at 550 °C. Because the underlying mechanism and anisotropy are specific to oriented crystalline systems, these results are not directly transferable to the unoriented amorphous electrolytes emphasized in this review.

In amorphous and semi-liquid phosphate-based systems, proton mobility is strongly influenced by the degree of polymerization. Yamaguchi and co-workers [101] studied polyphosphate-based glasses prepared by electrochemical substitution of sodium ions with protons. At 200 °C, the dependence of proton transport on phosphate chain length was found to vary with glass composition. For glasses with the general formula xNaO_1/2–1WO₃–8NbO_5/2–5LaO_3/2–(86 − x)PO_5/2, over the phosphate concentration range x = 28–38, depolymerization enhanced charge-carrier mobility, whereas the composition with the highest concentration of proton-bearing constituents (x = 40) exhibited the opposite trend. Based on molecular spectroscopy investigations, the latter behavior was attributed to an increasing concentration of the protons trapped due to strong O-H bonding within the growing amount of pyrophosphate anions present in the system. Conversely, in highly polymerized systems, conductivity can also be limited because constraints on proton motion along long polyanionic chains increase with chain length. These observations link composition-driven depolymerization to acidic site environments and hydrogen-bond strength, providing a mechanistic axis for interpreting why σ and E_a vary across nominally similar phosphate glasses.

Omata et al. [102] examined proton-transport properties of phosphate glasses at their glass transition temperature across a set of 32 glasses spanning a range of T_g values, focusing on proton conductivity and the effective concentration of charge carriers. They reported that, nearly independently of composition, proton mobility and the related diffusion coefficient at T_g were approximately constant and equal to 2 × 10⁻⁸ cm²·V⁻¹·s⁻¹ and 4 × 10⁻¹⁰ cm²·V⁻¹·s⁻¹, respectively. This was attributed to protons acting as cross-linkers of the polyphosphate lattice due to strong hydrogen bonding. In contrast, the corresponding conductivity values varied, indicating that conductivity remained strongly composition-dependent even at T_g, consistent with differences in carrier concentration and network environments rather than solely mobility.

Sumi and co-workers [78] correlated proton conductivity of phosphate glasses with their structure using NMR, Raman, and AC impedance investigations. For ternary xBaO-(40 − x)ZnO-60P₂O₅ glasses, their properties depended on the ratio between barium and zinc cations present in the system. While pristine zinc phosphate exhibited the lowest conductivity (2 × 10⁻⁸ S·cm⁻¹ at 150 °C), substituting zinc with barium (x = 10 to 40%) increased conductivity. The effect was most pronounced at the smallest x value, yielding an approximately twofold increase in conductivity, whereas changes at x = 40 were smaller. This trend correlated with spectroscopy: the x = 10 sample revealed both the most significant increase in non-bridging oxygen concentration and the corresponding changes in the shape of ¹H and ³¹P NMR traces. The authors suggested that the observed impact of barium cations resulted from structural disordering of the glass lattice related to the breaking of O-P-O moieties, increasing the concentration of mobile protons. The explanation of this phenomenon was attributed to differences in the coordination numbers of Zn²⁺ and Ba²⁺ in their respective M-O environments, reported as four and eight, respectively [103]. Moreover, the lowest concentration of the barium additive was suggested to yield the highest degree of structural disruption and therefore the largest deviations in the transport parameters studied.

Similar conclusions regarding the role of hydrogen bonding were reported by Abe et al. [104]. They found that protons incorporated into phosphate glass matrices were significantly more mobile than sodium and silver ions incorporated into systems of the same basic chemical composition. In contrast to most of the previously reported studies, in this case, charge-carrier transport was determined in a D.C. electrical field. Independently of the applied methodological alteration, the authors reached a similar conclusion regarding weakening of O-H bonds in the glassy matrix due to hydrogen bonding of mobile protons. Moreover, upon studying MO-P₂O₅ structures incorporating alkaline earth metals, the authors devised clear linear dependencies between the position of the OH band originating from O-H stretching and the logarithm of proton mobility as a function of the ionic radius of the cation used. In the former case, increasing cation radius lowered the wavenumber (i.e., 3350 cm⁻¹ for beryllium down to 2800 cm⁻¹ for barium), consistent with stronger hydrogen bonding. In the latter, mobility increased by about four orders of magnitude over the same cation change. The authors further reported that conductivity remained unaffected by incorporation of sodium and silver ions into the glass structure, which they interpreted as evidence that conductivity in these systems is predominantly protonic. Finally, the same author [102] also analyzed relations between electrical conductivity, activation energy, and bonding states characterizing O-H moieties present in the phosphate glass systems. Their work focused on materials based solely on alkaline earth cations or on their combination with aluminum. The results were compared to other similar amorphous systems based on borate and germanate anionic sub-lattices. It was concluded that, similarly to the systems discussed above, the position of the O-H stretching band provides a valuable measure of hydrogen-bond strength and, by extension, proton mobility and conductivity.

Critical synthesis (polyphosphates): High conductivities are typically reported when a connected pore network stabilizes sufficient water activity and when acidic phosphate groups provide a dense H-bond network for proton transfer. Across studies, direct comparison is often hindered by inconsistent RH reporting, differences in pore size distributions caused by aging/drying/thermal treatment, and time-dependent structural relaxation that can reduce water uptake and conductivity. For engineering relevance, datasets should be benchmarked at a defined RH, include activation energies, and report durability under thermal/humidity cycling.

3.6. Phosphate-Based Glasses: Mechanism, Constraints, and Screening Implications

Across phosphate-based glasses, the most informative comparisons are obtained when composition and processing are explicitly connected to network depolymerization (including modifier-driven changes in non-bridging oxygen populations), which in turn governs acidic site density, hydrogen-bond connectivity, and the balance between hopping-dominated versus carrier-assisted transport. In practical terms, σ values in the 120–200 °C window should be interpreted primarily through σ(T,RH) together with E_a and hydration/porosity descriptors, because nominally similar σ can originate from different regimes with very different stability under gradients and cycling. Reports of improved σ in modified or composite phosphate systems are most credible when they simultaneously demonstrate (i) stable retention of proton-bearing groups and/or bound water and (ii) microstructures that sustain percolation pathways without progressive dehydration, pore collapse, or structural relaxation.

From an application standpoint, phosphate systems face a recurring durability constraint: when transport relies on mobile acid species or surface-enriched phases, acid migration/leaching and compositional drift can degrade both bulk conductivity and electrode/electrolyte interfaces over time, increasing contact resistance and reducing device performance. Consequently, long-term benchmarks under thermal/humidity cycling, accompanied by post-test compositional checks, are necessary to distinguish stable network-supported transport from transient, hydration-driven conductivity enhancements.

In phosphate-based glasses, the relevant causality chain can be defined as follows:

(i)

Composition and modifiers (e.g., the type and fraction of modifier oxides, cation substitutions, and the overall “loosening” of the network) together with processing history (which determines what species and structural motifs survive drying and heat treatment) set the degree of network depolymerization. This depolymerization governs how many and what kind of acidic/proton-bearing sites are formed and remain stable in the glass (e.g., the population and local environments of P-OH groups and how strongly protons are bound).

(ii)

The same network state also controls the formation and persistence of a continuous hydrogen-bond network (between proton-bearing groups and, when present, water), as well as water accessibility and retention (i.e., whether water is structurally retained/bound or more weakly held and easily lost under gradients). These factors jointly determine the dominant transport regime: either predominantly Grotthuss-type hopping along a hydrogen-bond network anchored by proton-bearing groups or carrier-assisted (vehicle-type) transport that depends more strongly on mobile water/acid-associated species in connected porosity and near-surface regions.

(iii)

In addition to that, the synthesis methodology plays a decisive role in defining the microstructural features governing proton transport, extending well beyond simple compositional effects. In particular, studies on sol–gel-derived phosphate-based glasses have demonstrated that preparation conditions directly influence the retention of hydroxyl groups and molecular water, which are critical for the formation of hydrogen-bonded networks and local free-volume characteristics. The presence and distribution of these species modify the local proton environment, facilitating proton accommodation and altering the connectivity of potential transport pathways. Consequently, macroscopic proton conductivity should be regarded as a structure-sensitive property controlled by synthesis-dependent organization at the atomic and nanometric scales rather than solely by nominal proton concentration or hydration level. This interpretation is consistent with observations that chemically similar systems prepared via different routes may exhibit markedly different transport characteristics.

(iv)

Finally, only at this point does a meaningful interpretation of σ(T,RH) become possible, because both the magnitude of conductivity and its humidity/temperature dependence follow from the transport regime and the underlying hydration state. Finally, the same chain extends to application-relevant constraints: if conductivity relies on mobile acid/water species or surface-enriched phases, the system becomes more vulnerable to hydration drift, thermal/humidity cycling sensitivity, and acid migration/leaching, which can degrade not only the bulk electrolyte but also the electrode/electrolyte interface (e.g., through increasing contact resistance and chemical incompatibility over time).

Moreover, from a mechanistic perspective, the evolution of proton transport behavior can be rationalized using the conceptual framework established for disordered proton conductors. In structurally rigid networks, proton mobility is typically dominated by localized hopping between fixed coordination sites. However, the retention of hydroxyl groups and weakly bound water, as reported for sol–gel-derived systems, promotes dynamically reconfigurable hydrogen-bond networks that enable structurally assisted proton migration [92,105]. Within the broader model of proton conduction proposed by Kreuer, this behavior corresponds to a transition from structurally constrained proton transfer toward a regime governed by hydrogen-bond fluctuations and rapid reorientational dynamics of the host matrix [106]. The resulting reduction in activation barriers explains the frequently observed increase in conductivity and the apparent “liquid-like” transport characteristics. Importantly, this regime does not imply the formation of a true liquid phase but rather reflects enhanced proton diffusivity arising from dynamic structural reorganization [106].

4. Polysilicate-Based Systems

4.1. Silica Network Chemistry and Silanol Condensation

Orthosilicic acid as the basic compound from the group of silicates is common not only in the earth’s crust but also in living organisms as a component of metabolic changes. Interestingly, it is a structure that is both unstable and difficult to isolate, meaning that solutions to obtain it in a pure and stable form are constantly being sought after [107]. Such a solution could significantly extend the synthesis methodology for various materials, which would significantly influence the production of new silicate materials and other silicon dioxide-based materials. Taking advantage of this aspect, as well as the amorphous nature of these compounds, the practical aspect of using them in electrochemistry is the growth of silicon dioxide on the crystalline plane of a metal or semiconductor at room temperature [107].

Silicon dioxide (SiO₂) is an inorganic compound that occurs naturally, for example as silica-rich minerals and sands, and in multiple structural forms ranging from crystalline phases (e.g., quartz) to amorphous silica [108]. Owing to its high hardness and chemical stability, SiO₂ is widely used in technological applications, including glass and silica-based materials, as well as in electronics (e.g., as an insulating/dielectric material) [109]. In many applications, polysilicate and polysilicic acid-derived networks can be readily obtained via hydrolysis of organosilicon precursors. This route enables the formation of an amorphous solid phase [110] rich in terminal silanol (Si-OH) groups, which makes it susceptible to further condensation reactions.

From the perspective of intermediate-temperature membranes, this chemistry implies an intrinsic limitation: silicate-rich networks typically exhibit lower intrinsic acidity than phosphate-rich analogs, so proton transport often depends strongly on hydration state and the formation of extended hydrogen-bond networks mediated by water in pores and near-surface regions (see Table 2). Consequently, σ(T,RH) in silicate glasses is frequently more sensitive to RH and thermal history than in phosphate-rich systems, and direct comparison requires explicit reporting of hydration conditions and processing descriptors.

4.2. Acid-Site Engineering in Silica Electrolytes

One strategy to enhance proton transport in silica-based structures is functionalization with sulfonic acid groups. Kim and co-workers [111] prepared a membrane via a sol–gel process using alkoxysilane precursors. A thiol-functionalized silane was co-hydrolyzed and incorporated into the inorganic network, followed by oxidation of the -SH groups to generate acidic –SO₃H sites. The maximum conductivity reached 3.71 mS·cm⁻¹; together with the reported low methanol permeability, this material was proposed as a potential membrane for DMFC applications. In screening terms, this approach directly increases acidic site density and can stabilize proton-bearing groups independently of bulk silicate acidity, but it also raises questions of chemical stability and long-term retention of functional groups under 120–200 °C operation.

In another report [112], the authors proposed doping silica gel by combining a poly(vinyl alcohol)-based composite matrix with acids such as HClO₄, H₂SO₄, and H₃PO₄. Depending on the additive and its concentration, conductivities in the range 10⁻⁵–10⁻² S·cm⁻¹ were obtained. The resulting proton conductors were subsequently applied in solid-state electric double-layer capacitors, achieving 31–44 F·g⁻¹ (with activated carbon powder as the high-surface-area electrode material). These systems illustrate a recurring trade-off in silicate-rich conductors: σ gains can be achieved by introducing strong acidic species, but stability then depends on acid retention/migration, polymer–inorganic interactions, and how humidity and thermal treatment modify the conduction pathway.

4.3. Hydration States in Sol–Gel Silica: Porosity, Heat Treatment, and σ(RH,T)

Nogami et al. [113] investigated silicate glasses obtained from tetraethoxysilane by means of the sol–gel process. A water–ethanol mixture was used in the hydrolysis step of the process to limit its kinetics. The addition of formamide to the solution was used to control the porosity of the resulting glass. The viscous sol was dried afterwards for about a week to form a 0.1 mm thick sample of the stiff xerogel. A monotonic decrease in conductivity was observed after heating the sample between 60 and 140 °C. Despite the limited proton mobility in the resulting system, relatively high conductivity was measured. The authors attributed these values to the high porosity of the material produced via the sol–gel route and argued that this porosity enables incorporation of large amounts of water into the internal structure. They classified water in the material into four categories: physically adsorbed, chemically adsorbed, chemically bonded in Si-OH groups present on the surface of the pores, and Si-OH groups belonging to the internal structure of the glass. Heat treatment disrupts the balance of these species and promotes further condensation of the silicate subnetwork, leading to the formation of Si-O-Si bonds and expulsion of water molecules.

In the next stage of sample preparation, gradual heating of the material to 400–800 °C resulted in a decrease in specific surface area from 960 m²·g⁻¹ to 690 m²·g⁻¹. On the other hand, glasses exposed to atmospheric humidity reabsorbed moisture, and therefore the amount of water incorporated into the structure is related to its vapor pressure. Hydration increased rapidly for steam pressures between 0.5 and 0.8 atm, suggesting that in samples equilibrated under the latter conditions, all pores were filled with adsorbed water. This interpretation is consistent with conductivity data: at room temperature, conductivity increased from 2.5 × 10⁻⁴ S·cm⁻¹ to 4 × 10⁻⁴ S·cm⁻¹ as water vapor pressure increased from 0.6 to 0.9 atm. At 100 °C, the difference was smaller, with conductivity increasing by a factor of 1.25 and reaching 5 × 10⁻⁴ S·cm⁻¹ for the most humidified sample.

A more detailed investigation of the role of water molecules in protonic conduction of wet silica gels was reported by Nogami and Abe [114]. In FT-IR spectra recorded between 20 °C and 800 °C, changes were observed in the shape and intensity of the broad band assigned to O–H stretching modes of hydroxyl groups. After deconvolution, a component centered around 3700 cm⁻¹ was assigned to Si-OH moieties, whereas bands around 3200 and 3400 cm⁻¹ were attributed to immobilized water of differing bond strengths. Upon heating, the overall intensity decreased and the dominant component changed: at room temperature, the highest intensity near 3400 cm⁻¹ was attributed to physically adsorbed water, while at 200 °C the dominant band near ~3200 cm⁻¹ corresponded to chemically adsorbed species. Both components largely vanished upon further heating. The water-loss process was found to be at least partially reversible, because the intensities of the 3200 and 3400 cm⁻¹ sub-bands increased with exposure to water vapor. In contrast, the 3700 cm⁻¹ band intensity was not affected by hydration. Further experiments indicated that this heat treatment led to a partially irreversible loss of conductivity. Samples preheated to 600 °C, even after further hydration, exhibited conductivities ranging from 10⁻¹³ to 10⁻¹¹ S·cm⁻¹ with activation energies close to 150 kJ·mol⁻¹.

Together, these studies illustrate the central mechanistic point for silicate xerogels: conductivity can be dominated by water-mediated pathways in connected porosity, and heat treatment can shift the balance from physically/chemically adsorbed water toward irreversible network condensation, which reduces both water retention and σ. Therefore, for sol–gel-derived silicates, the minimum “processing fingerprint” needed for comparison should include, at least qualitatively, hydrolysis/condensation conditions (pH or catalyst identity), aging time, drying protocol, and thermal treatment/calcination schedule, because these variables directly control pore size distributions, percolation, and hydration stability (Table 2).

4.4. Bulk Silicate Glasses Under Controlled Water Content: NBO Effects and Stability Limits

Fanara and Behrens studied a set of proton glassy conductors based on the CaAl₂Si₂O₈—CaMgSi₂O₆ family of compounds [115]. The materials were synthesized using a high-temperature process (1523–1723 K) under high pressure (200 MPa). Under these conditions, a water content (confirmed by Karl Fischer titration) of up to 3% by weight was maintained. A positive correlation between conductivity at 685 K and water content was observed across the full hydration range studied. Combined conductivity and structural analyses further indicated the importance of non-bridging oxides in enabling transport of aqueous charge carriers. However, the measured proton conductivities remained below application-relevant levels, ranging from 4 × 10⁻¹⁰ S·cm⁻¹ to 3 × 10⁻⁶ S·cm⁻¹, while correlated water diffusion coefficients were estimated at 10⁻¹⁷–10⁻¹² m²·s⁻¹. Mechanistically, these results are consistent with a limitation of silicate-rich matrices: even when water is present, transport can remain constrained by network environments and the availability/connectivity of pathways for aqueous carriers, and gains in σ are not guaranteed unless water is both retained and effectively percolated.

Ren and co-workers [116] investigated GaPO₄-SiO₂ gel glasses in terms of their structure and proton mobility using NMR spectroscopy. Although the structural characterization was carefully reported, the lack of conductivity data and electrochemical tests limits the utility of these results for transport screening, illustrating a recurring reporting gap in the silicate literature (Table 2).

In another report concerning protonic conductivity in silicate systems, barium disilicate-based glasses were presented by Behrens and co-workers [117]. Protonic conduction was inferred from impedance measurements and supported by independent NIR investigations indicating the presence of both bonded –OH groups and molecular water. Limited stability was reported for temperatures exceeding 523 K (250 °C), above which conductivity decreased markedly. This behavior was attributed to the conversion of strongly bound –OH groups into water molecules, followed by water diffusion out of the sample, thereby reducing proton-transport capability associated with these moieties. This is an application-relevant failure mode: conductivity supported by hydroxyl/water species can degrade irreversibly if thermal treatment or operating temperature drives conversion and loss of water, emphasizing the need for cycling tests and retention metrics rather than single-point σ.

4.5. Thin-Film Aluminosilicates: Thickness Scaling and Percolation Transport

In thin-film aluminosilicate glassy conductors investigated by Aoki et al. [118], conductivity exhibited complex behavior dependent on both temperature and humidification. According to the authors, this reflects contributions from both Broensted and Lewis protonic carriers. The high-temperature component associated with Broensted protons became independent of humidity above 300 °C and could persist up to 500 °C.

Another thin-film silicate-based proton-conductive system was also investigated by Aoki et al. [110]. A set of M_0.1Si_0.9O_x systems with M = Al, Ga, Hf, Ti, Ta, and La was synthetized and examined. Dry air conductivity was determined for films with thicknesses in the 10–1000 nm range, increasing from about 10⁻⁹ S·cm⁻¹ at 400 K to 10⁻⁵ S·cm⁻¹ at 700 K. This behavior was interpreted theoretically within the framework of percolation theory.

Addressing a related issue, Aoki and co-workers reported a contribution [118] focused on scaling up a previously developed method [119] for manufacturing thin films of aluminosilicate proton conductors. The measured conductivities were relatively low, ranging from 10⁻⁸ to 10⁻⁴ S·cm⁻¹ at 110 and 350 °C, respectively. Conductivity also decreased with increasing film thickness, following an inverse power-law dependence.

Only a limited number of studies have addressed charge-transport mechanisms in waterless and poreless silicate-based systems. Among them, Aoki et al. [110] discussed thickness-related effects in thin films of amorphous M_0.1Si_0.9O_x (M = Al, Ga, Hf, Ti, Ta, and La). Although percolation-type behavior was suggested, conductivities up to 10⁻⁵ S·cm⁻¹ for films thinner than 40 nm are not directly transferable to bulk materials, which are the focus of the present review. Therefore, despite theoretical studies on proton transport in crystalline structures [120], first-principles analyses of proton mobility in amorphous SiO₂ [121], and molecular dynamics simulations of water and proton behavior at extended amorphous silica surfaces [122], the issue remains unresolved experimentally. More detailed experimental characterization is still required. Accordingly, experimental efforts have been devised to provide deeper insight into these phenomena, occurring in crystalline solids [123] where proton injection must be considered, as well as in more complex non-homogeneous materials such as gels [124].

4.6. Transport-Regime Crossover: Anhydrous vs. Hydrated Silica and Pore Size Thresholds

In parallel, it has been shown that preparation routes play an important role in determining proton-transport properties of glassy ionic conductors. Investigations of melt-quenched glasses [125] established a quadratic relationship between conductivity and proton concentration under anhydrous conditions, i.e., in the absence of molecular water. This behavior was attributed to proton mobility being rate-controlled by hydroxyl-group formation. When the same materials were humidified such that [H₂O]/[H⁺] >= 1, conductivity increased drastically by three to four orders of magnitude. Mechanistically, this transition is consistent with a regime change from matrix-coupled hopping limited by hydroxyl formation/dissociation to a more water-mediated, higher-mobility transport pathway, and it exemplifies why σ must be reported with explicit hydration conditions to be interpretable (Table 2).

Significantly different trends were reported by Daiko et al. [126] for sol–gel-derived materials with average pore diameters ranging from 2 to 25 nm. In that work, protonic conductivity was attributed to water molecules within the pore space, which were also claimed to govern related transport descriptors such as electrical modulus and proton diffusion coefficient. An abrupt deviation in properties was reported with a limiting pore size of about 15 nm: above this threshold, pore space was not further filled with water molecules and thus did not contribute to conductivity. For smaller pores, physically bonded water molecules were confirmed and associated with pathways of increased proton mobility. At the smallest pore diameters (~2 nm), the opposite effect was observed, with conductivity decreasing because water becomes predominantly strongly chemically bonded to pore walls, restricting proton hopping. Consequently, an optimal average pore diameter of 4 nm was proposed as most favorable for fast proton conduction. This result captures the silicate trade-off in a quantitative way: increasing porosity and water content can increase σ, but excessively confined pores can immobilize water and suppress hopping, while larger pores may fail to retain water under gradients.

These findings were also supported by Nogami [127] for silica-based glasses obtained from Si alkoxides. In that work, an order-of-magnitude decrease in activation energy was observed upon hydration, decreasing from the ~100 kJ·mol⁻¹ characteristic of “dry” materials to values characteristic of aqueous solutions of strong protonic acids (~10 kJ·mol⁻¹). These two regimes corresponded to values of the product [H⁺][H₂O] ranging from 0.3 to 100, respectively. The transition between these limiting compositions was associated with a change in the conductivity-limiting process from hopping along the solid matrix to liquid-like transport behavior.

Nogami and Abe [127] also analyzed quantitatively the effect of molecular water on the conductivity of porous silicate glasses. For water- and pore-free materials, a quadratic dependence of conductivity on proton concentration was observed. Porous but still water-free glasses behaved differently: activation energy decreased linearly with a logarithmic increase in mobile proton concentration, while conductivity increased in the same order. For glasses in which both mobile protons and molecular water contributed, transport parameters were related to the product of their concentrations, consistent with a mechanism involving hopping between –OH sites bonded to the polysilicate lattice and neighboring water molecules. In this interpretation, activation energy is governed by the energy required for the dissociation of protons from Si-O-H and H-O-H moieties. In these studies, water content was controlled by pre-annealing at 160 to 310 °C for 48 h. Conductivity was measured by D.C. methods using a vibrating reed electrometer from 120 to 310 °C. The resulting conductivities ranged from 2 × 10⁻¹³ S·cm⁻¹ to 4 × 10⁻¹² S·cm⁻¹, while activation energies ranged from 62 kJ·mol⁻¹ for the best conductive sample containing about 1.6 wt.% of water to about 130 kJ·mol⁻¹ determined for the set of materials with the worst conductivity and lowest water content (~0.4%).

4.7. Dielectric Fingerprints and Relaxation Processes in Silica Gels

The impact of preparational details on the dielectric properties of silicate-based gels was described by Gutina et al. [128]. Both conductivity and relaxation behavior were found to depend strongly on preparation temperature, which affects the kinetics of the chemical processes involved in gel formation and thus a broad range of material properties. Some common features were nonetheless identified across matrices and attributed to the complex dynamics of hydrochloric acid molecules remaining from the gel-curing step and their interactions with internal pore surfaces. Based on dielectric spectra, the authors calculated fractal dimensions of pore space and porosity factors describing internal structure. In addition, the overall dielectric response was interpreted as the coexistence of two independent relaxation processes: one resembling behavior typical of glass-type materials and the other resembling polymer-like relaxations. This was argued to be consistent with the chemical constitution and the physical appearance of the samples, described as intermediate between quartz glass and silicone rubber. Mechanistically, such dielectric “fingerprints” provide indirect evidence that a preparation-dependent microstructure and retained species (e.g., acid residues and confined water) can create multiple relaxation/transport pathways, reinforcing that processing descriptors are essential for interpreting σ and aging behavior.

4.8. Silicate-Based Glasses: Transport Regimes, Trade-Offs, and Screening Implications

Silicate-rich glasses differ from phosphate-rich systems primarily through lower intrinsic acidity, which shifts the burden of proton transport toward hydration-mediated pathways and makes σ(T,RH) strongly contingent on water activity and microstructure. In sol–gel-derived silicates, the parameters that define the processing fingerprint—hydrolysis/condensation conditions (often pH/catalyst identity), aging, drying protocol, and thermal treatment—set the pore architecture (BET, pore size distribution, and percolation), which in turn controls water retention and the connectivity of water-mediated conduction pathways. This produces a characteristic trade-off: σ can increase when a connected pore network stabilizes sufficient water activity, but the same water-mediated mechanism can introduce dimensional/mechanical instability and long-term drift under thermal/humidity cycling, particularly when heat treatment drives irreversible condensation and loss or redistribution of water/–OH species.

Mechanistically, the reviewed studies indicate a crossover from matrix-coupled hopping in anhydrous or water-poor silicates (high E_a and σ constrained by hydroxyl formation/dissociation) to more liquid-like, carrier-assisted transport upon hydration (lower E_a and stronger RH dependence), with pore size thresholds controlling whether water is retained, mobile, or immobilized at pore walls. Therefore, intermediate-temperature screening of silicate-based conductors should be based not only on σ magnitude but also on a consistent set of descriptors that ties σ(T,RH) and E_a to a clearly defined hydration state and to evidence that this hydration state and microstructure remain stable under cycling (Table 2). In practice, meaningful comparison requires that sol–gel fingerprints, porosity metrics, and hydration conditions be reported together with σ(T,RH), E_a (fit window stated), and at least qualitative stability/aging observations; otherwise, reported conductivities remain weakly comparable. Nevertheless, for convenience of the reader, available data is gathered in

Table 4. Polysilicate-based glassy proton conductors—comparison checklist (populated with key values from cited studies).

System/Composition	Synthesis Route	Porosity/BET	Hydration/RH	σ(T) and Ea	Stability/Aging Notes	Ref.
Orthosilicic acid/oligomer	Non-aqueous selective synthesis	n/a	n/a	n/a	n/a	[107]
Vitreous silica	n/a	n/a	n/a	n/a	n/a	[108]
Phosphate glasses doped with SiO2	Melt-derived glass study	not reported	not reported	not reported	not reported	[109]
Polysilicic acid-derived networks via hydrolysis of organosilicon precursors	Hydrolysis/condensation of organosilicon precursors	not reported	silanol-rich network implied	not reported	not reported	[110]
Sulfonic acid-functionalized silica membrane (–SO3H sites)	Sol–gel from alkoxysilanes + thiol -> oxidation to -SO3H	not reported	not reported	σmax = 3.71 mS·cm−1; Ea not reported	low methanol permeability; DMFC-oriented	[111]
Acid-doped silica gel in PVA composite matrix (HClO4/H2SO4/H3PO4)	Silica gel + polymer composite + acid doping	not reported	not reported	σ ≈ 10−5–10−2 S·cm−1 (acid/loading-dependent); Ea not reported	EDLC context; durability not emphasized	[112]
Porous silica glass (TEOS sol–gel; formamide porosity control)	Sol–gel (TEOS), controlled hydrolysis; formamide induced porosity; xerogel drying ~1 week	SSA decreases 960 > 690 m2·g−1 after 400–800 °C	water vapor-pressure dependence; example p(H2O) 0.6 -> 0.9 atm	σRT 2.5 × 100−4→4 × 10−4 S·cm−1 for p(H2O) 0.6→0.9 atm; at 100 °C up to 5 × 10−4 S·cm−1; Ea not reported	heating 60–140 °C decreases σ; condensation Si-O-Si and water expulsion	[113]
Water-cooperative proton conduction in silica glasses (FT-IR deconvolution)	Heat treatment + rehydration experiments	not reported	physically vs. chemically adsorbed water distinguished; partial reversibility on rehydration	after preheat to 600 °C: σ ~10−13–10−11 S·cm−1; Ea ~150 kJ·mol−1	partially irreversible loss of σ after strong heat treatment	[114]
Hydrous CaAl2Si2O8–CaMgSi2O6 glasses (oxygen-joined; water-bearing aluminosilicates)	High T (1523–1723 K) under high pressure (200 MPa)	not reported	water up to ~3 wt.%	σ at 685 K: 4 × 10−10–3 × 10−6 S·cm−1; Ea not reported	σ correlates with water content; values low for applications	[115]
Mesoporous GaPO4–SiO2 sol–gel glass (structural/NMR)	Sol–gel mesoporous	not reported	not reported	no σ/no electrochemical tests	limitation: lack of conductivity data	[116]
Hydrous BaSi2O5 glass (barium disilicate-based)	Glass	not reported	bonded -OH + molecular water (NIR)	σ not reported	σ drops above 523 K due to -OH→H2O conversion and out-diffusion	[117]
Aluminosilicate acid thin films (Brønsted + Lewis sites)	Thin film	not reported	humidity-dependent; becomes humidity-independent above ~300 °C	numeric σ not provided	σ persistent up to 500 °C; mixed-carrier mechanism	[118]
Gas-tight nanomembranes of silica-based double oxides	Nanomembrane fabrication	not reported	not reported	not reported	referenced as a fabrication/scaling motif	[119]
First-principles proton conduction in PCFC electrolytes	Computational	n/a	n/a	n/a	background	[120]
Proton mobility in amorphous SiO2	Computational/defect context	n/a	n/a	n/a	background	[121]
Water/amorphous silica interface: H-bond lifetimes and proton transport (MD)	Simulation	n/a	interfacial water	n/a	background	[122]
Proton conduction/injection in solids	Review	n/a	n/a	n/a	background	[123]
Proton diffusion in pores of silicate sol–gel glasses	Sol–gel porous glass (conceptual)	pore diffusion context	pore water	n/a	background	[124]
Quadratic σ vs. proton concentration in melt-quenched glasses (anhydrous)	Melt-quenched; anhydrous condition emphasized	poreless; anhydrous	anhydrous vs. humidified contrast	σ ∝ [H+]2	humidification can raise σ by 3–4 orders	[125]
Pore size effect in sol–gel porous silica glasses	Sol–gel porous silica	avg pore diameter 2–25 nm; threshold ~15 nm; optimum ~4 nm; ~2 nm depletes conduction	pore water; bound vs. free depends on pore size	σ attributed to water in pores; numeric σ/Ea not quoted here	percolation/filled pores; too small pores bind water too strongly	[126]
Nanopore-controlled silica glasses (Nogami; hydration lowers Ea)	Silica from alkoxides; nanopore-controlled; hydration tuned by water content	nanopore-controlled (no BET numbers quoted here)	controlled by product [H+]·[H2O]; hydration triggers regime change	Ea drops ~100→~10 kJ·mol−1 (dry→hydrated); separate dataset: σ 2 × 10−13–4 × 10−12 S·cm−1 (120–310 °C after pre-anneal), Ea 62–130 kJ·mol−1 for 0.4–1.6 wt.% water	explicit hopping→liquid-like transition; pre-anneal governs retained water	[127]
Dielectric properties/relaxation in fast sol–gel glasses	Sol–gel; prep temperature impacts gel kinetics	fractal/porosity factors from dielectric spectra	HCl residues and pore interactions emphasized	conductivity/relaxation depend on prep temperature	two relaxation processes identified; microstructure-sensitive	[128]

below.

5. Conclusions

5.1. Cross-Comparison of Phosphate vs. Silicate Glasses

Taken together, Section 3 and Section 4 show that phosphate- and silicate-based glasses can be discussed within a common causal structure, but they populate it differently. Phosphate-rich matrices generally provide higher intrinsic acidity and a higher density of proton-bearing environments, so improvements in σ more often arise from changes in network connectivity and acidic site chemistry (e.g., modifier-driven depolymerization and hydrogen-bond connectivity) combined with hydration management. In contrast, silicate-rich matrices typically require either explicit acid-site engineering (functional groups and acid doping) or microstructures that stabilize water-mediated pathways; as a result, σ is frequently more RH-sensitive and more tightly coupled to pore architecture and thermal history.

This difference has direct implications for stability and for how a “high σ” should be interpreted. In phosphate glasses, a recurring application risk is acid migration/leaching or formation of mobile surface/near-surface phases that can yield high apparent σ but drive compositional drift and interfacial degradation over time. In silicate glasses, high σ commonly reflects water-mediated transport in pores, which can be beneficial for conductivity but introduces a trade-off with dimensional/mechanical stability and drift under thermal/humidity cycling, especially when irreversible condensation reduces water retention. Therefore, across both families, the practical relevance of a reported σ requires that the transport regime be established (via σ(T,RH) trends and E_a) and that hydration retention and durability be benchmarked rather than assumed.

A unified comparison across Section 3 and Section 4 is only possible when the same minimum dataset is available. At minimum, each candidate should be reported with σ at a defined T and RH (or water content), an Arrhenius analysis with the stated fit range, and quantitative or semi-quantitative hydration/porosity descriptors (e.g., BET/pore size and a bound/free water statement). For intermediate-temperature membranes, this dataset must be extended with stability and engineering metrics—chemical stability (hydrolysis resistance and leaching/acid migration), mechanical robustness and cycling response, and at least one interface/device-relevant metric where available—because both phosphate- and silicate-rich systems can otherwise appear promising by σ alone while failing under realistic operating constraints. This cross-comparison logic is operationalized in Table 2 and Table 3 and motivates a single joint comparison table spanning both families that will be used as the baseline for evaluating composite and hybrid strategies in Part II.

5.2. Further Possibilities

In the 120–200 °C operating window, glassy phosphate and silicate proton conductors can play the role of a starting point towards more complicated mixed systems, such as glassy phosphate–silicate electrolytes, which would occupy a niche between low-temperature PFSA membranes and liquid-acid PAFC systems. Across the reviewed literature, proton conductivity is governed by a coupled set of variables: (i) network chemistry (acidic site density and connectivity), (ii) hydration state and RH, and (iii) pore architecture shaped by processing conditions (including sol–gel aging, drying, and thermal treatment). While phosphate-rich systems can reach higher conductivities than many silicate-rich analogs, the practically relevant performance envelope is often constrained by stability under thermal/humidity cycling, water-retention drift, and interfacial compatibility with electrodes.

Part I of the review presented herein yields a screening-oriented set of design rules intended to make cross-study comparisons commensurable and device-relevant. Specifically, it aimed to: (1) report and benchmark σ at a defined RH and temperature together with activation energies; (2) treat porosity and hydration control as primary design variables, supported by pore metrics (e.g., BET and pore size distributions) and water-uptake/retention data; (3) couple transport metrics with durability, including cycling stability, leaching/acid migration, and mechanical integrity; and (4) evaluate electrode/electrolyte interfaces early in the process, including contact resistance and chemical compatibility. In practice, these criteria also clarify why single-number σ claims are insufficient: the same nominal conductivity can arise from fundamentally different transport regimes and hydration dependencies, which lead to very different stability outcomes under operating gradients.

The literature surveyed in Part I also shows clear progress within single-anion phosphate and silicate matrices. Advances in controlling acidity, hydration, and microstructure—particularly through sol–gel-derived architectures and targeted incorporation of proton-bearing groups—have enabled substantial conductivity gains under humidified conditions and improved mechanistic understanding of water-mediated versus more matrix-coupled transport. At the same time, when assessed against intermediate-temperature membrane requirements as a whole, single-anion networks rarely satisfy the combined constraints simultaneously. Conductivity improvements frequently remain contingent on fragile hydration states and porous architectures that are susceptible to dehydration hysteresis, structural relaxation, and long-term drift; conversely, routes that enhance stability can suppress σ. In addition, chemical durability (hydrolysis resistance and acid migration/leaching), mechanical robustness under cycling, and electrode/electrolyte interfacial losses are reported inconsistently and often remain limiting factors even when σ and E_a appear promising.

Consistent with the cross-comparison developed in Section 3 and Section 4, the dominant limitations also differ in emphasis between the two families. In phosphate-rich glasses, performance gains are often linked to modifier-controlled network connectivity and acidic site environments, but long-term operation can be limited by acid migration/leaching, compositional drift, and associated interfacial degradation. In silicate-rich glasses, lower intrinsic acidity shifts the burden of conductivity toward hydration-mediated pathways in porous architectures or toward explicit acid-site engineering; as a result, σ gains are frequently accompanied by strong RH sensitivity and trade-offs with dimensional/mechanical stability and drift under thermal/humidity cycling, particularly when heat treatment drives irreversible condensation and loss or redistribution of water/–OH species.

Accordingly, the most credible pathway beyond the intrinsic trade-offs of single-anion matrices is to move toward mixed-network, composite, and hybrid designs that deliberately couple acidic site engineering with microstructural control and hydration stabilization while preserving chemical and mechanical durability. Part II therefore builds directly on the screening framework established here to evaluate strategies such as mixed phosphate–silicate networks, inorganic–inorganic and organic–inorganic composites, and interfacial conduction architectures, with an emphasis on whether they deliver concurrent gains in σ(T,RH), stability under cycling, and manufacturability rather than improving a single metric in isolation. Studies in solid-acid-based composite electrolytes, particularly those reported by Tatsumisago and Tadanaga [129,130], have demonstrated that phosphosilicate matrices may actively modify proton transport through interfacial and structural effects rather than acting as passive fillers. Related observations in cesium phosphate and mixed-anion systems, as discussed by Chisholm [131] and Ponomareva [132], further suggest that proton conductivity in hybrid systems is governed by coupled structural and dynamic phenomena, including acid retention, hydrogen-bond network reorganization, and local heterogeneity. Finally, it is worth stressing that unlike conventional solid-acid composites, proton-conducting glasses offer a fundamentally different host environment for proton dynamics, implying that transport–structure couplings may follow distinct mechanistic pathways. These observations point toward a promising research direction involving hybrid glass/solid-acid architectures, where proton conductivity, hydration stability, and interfacial proton mobility should be analyzed within a unified mechanistic framework.