A focused review on structures and ionic conduction mechanisms in inorganic solid-state proton and hydride anion conductors

Abstract

Solid-state proton and hydride anion conductors are an important family of materials as electrolytes for solid state electrochemical cells such as fuel cells, batteries, sensors, and gas separation membranes. Searching for new proton and hydride-anion conductors has been an active research area for many decades. The focus of this article is on reviewing the types and mechanisms of each proton/hydride-anion conductor developed and their pros and cons. This review starts off with the most studied and most promising perovskite structured oxides as proton conductors, followed by other types of perovskite-related structures such as the Ruddlesden–Popper phase, pyrochlores and rare earth orthoniobates/orthotantalates. This review then moves to solid polyanionic compounds as proton conductors, including sulfates, nitrates, and phosphates, which is followed by hydrates and nanocomposites. This review finally discusses the types and conduction mechanisms of new hydride-anion conductors that emerged recently.

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1. Introduction

Many solids exhibiting fast proton (H⁺) and hydride ion (H⁻) conduction have been reported for decades. It is generally understandable that H⁺ and H⁻ move fast in solids due to their size and charge.¹ Like any ionic conductors, H⁺/H⁻ conductors can be used as an electrolyte in electrochemical cells for H₂ conversion and storage; the latter is a vital clean technology to combat climate crisis and build a carbon pollution-free economy. In general, H⁺/H⁻ conductors are found in both organic and inorganic materials. Organic materials, such as Nafion (perfluorosulfonic acid polymer), show high proton conductivity at low temperatures (<90 °C) under hydrated conditions, while inorganic materials, on the other hand, show H⁺/H⁻ conduction in a higher and wider temperature range. In this review, we will focus primarily on the inorganic H⁺/H⁻ conductors. Despite many existing reviews on H⁺ conducting materials in the literature, they are mostly focused on material properties and applications in energy conversion and storage.^2–5 A review encompassing both H⁺/H⁻ conductors with a focus on structure and ionic conduction mechanisms is still lacking. In this article, we aim to fill this gap by reviewing recent progress of inorganic solid-state fast H⁺/H⁻ conductors in the perspectives of synthesis, structure and phase transformations, conductivity, and conduction mechanisms.

The inorganic proton conductors reviewed in this article include solid oxides, solid polyanionic salts, and hydrates. The solid-oxide family includes perovskites, Ruddlesden–Popper intergrowths, and ortho-niobates/ortho-tantalates; the solid polyanionic salt family includes inorganic acids, sulphates, and phosphates/pyrophosphates; the hydrate family includes hygroscopic mesoporous materials. Compared to the well-studied proton conductors, hydride anion (H⁻) conductors are under studied. Although H⁻ has a larger size than H⁺, it is still small enough for fast ion transport and is a strong reducing agent that could find unique applications in energy storage and conversion devices. For these reasons, there has been growing interest in H⁻ conductors in recent years. This class of materials includes hydrides, nitride-hydrides, perovskite oxyhydrides, and hydroflurides. In the following, we review the key aspects of these H⁺/H⁻ conductors.

2. Solid oxide proton conductors

2.1 Perovskite oxides

One of the most studied proton conductors is perovskite oxides. The typical chemical formula of the perovskite oxides is ABO₃, where A and B denote two different cations, see Fig. 1. A typical ABO₃ structure consists of a large 12-coordinated A cation (A = Ca, Sr, Ba, La, and Gd), and a small 6-coordinated B cation (B = Nb, W, Ce, Zr, Ti, Mn, Fe, Co, and Ga). The ideal ABO₃ structure is primitive cubic, but many ABO₃ exhibit a lower symmetry (e.g., hexagonal and orthorhombic) with slightly distorted unit cells because of the ionic size effect as described by the Goldschmidt's tolerance factor. Based on the types of A and B cations, ABO₃ can be an ionic conductor, semiconductor, and electronic conductor, finding applications in electrochemical devices such as solid oxide fuel cells and catalytic membrane reactors.^6–12

Fig. 1 The ideal structure of ABO₃.¹³

For a perovskite oxide to be a proton conductor, a proton in the form of on the O²⁻-sub-lattice can be generated by a hydration process under humid conditions or a hydrogenation process under reducing conditions as illustrated by the following Kröger–Vink defect reactions:

Clearly, the concentration of OH⁻ is favored by higher partial pressure of H₂O (pH₂O), oxygen vacancy () concentration, and lattice oxygen () concentration. Since the reactions are exothermic, e.g., the enthalpy (ΔH°) of the reaction (1) for BaCe_0.9Y_0.1O_3−δ and BaZr_0.8Y_0.2O_3−δ are −163.3 and −79.5 kJ mol⁻¹, respectively,⁴ the proton concentration () decreases with temperature.^3,14,15

Two proton conducting mechanisms have been proposed to describe the H⁺ transport in the perovskite-type oxides:^15,16 (1) the Grotthuss mechanism, Fig. 2(a), where H⁺ migrates via reorientation of H⁺ and the formation and cleavage of H-bonds with the adjacent lattice O, and (2) the vehicle mechanism, Fig. 2(b), in which H⁺ is first bonded with O²⁻ to form OH⁻ (or () by Kröger–Vink notations) and then diffuses through . There has been increasing evidence in recent years showing that both Grotthuss and vehicular mechanisms are involved in the H⁺ transport in perovskite oxides.

Fig. 2 Schematics of two H⁺ conducting mechanisms in a BaZrO₃-based perovskite oxide: (a) the Grotthuss mechanism and (b) the vehicle mechanism.¹⁷

Due to the presence of , perovskite proton conductors can also be good oxide-ion conductors, depending on the conditions. For example, H⁺ conductivity dominates under high pH₂O, low pO₂, and low temperature, whereas O²⁻ conductivity is favored by dry and high temperatures.¹⁸ The magnitude of proton conductivity also depends on the material composition. For example, Ba-containing perovskites such as BaCeO₃ and BaZrO₃ are among the best in all proton conductors studied so far.

The first reported perovskite H⁺ conductor was doped SrCeO₃.^19,20 Since then, doped BaZrO₃ and BaCeO₃ have become the most studied proton conductors due to their higher conductivity.^10,21–26 In general, doped BaCeO₃ exhibit a higher conductivity (e.g. ∼2 × 10⁻² S cm⁻¹ at 600 °C) than doped BaZrO₃,^27,28 but their chemical stability in H₂O and CO₂ is generally poorer. To achieve a good balance among chemical stability, ionic conductivity, and mechanical property, “solid solution” mixing between BaCeO₃ and BaZrO₃ is a widely adopted strategy. However, to further increase the degree of hydration and thus proton conductivity, the B site needs to be partially doped by rare earth cations such as Y³⁺, Yb³⁺, Gd³⁺, and Sm³⁺. For example, Y- and Zr-doped BaCeO₃ (BaCe_1−x−yZr_xY_yO_3−δ, BCZY) has been one of the most studied systems for H⁺ conductors with high conductivity and chemical stability under CO₂, H₂O, or H₂S-containing atmosphere.^29,30 Similarly, Y-, Zr- and Yb-triple-doped BaCeZrO₃ with compositions of BaZr_0.1Ce_0.7Y_0.2−xYb_xO_3−δ (x = 0 to 0.2, BZCYYb) has also been reported with good ionic conductivity and chemical stability.²³Fig. 3 shows the comparison of ionic conductivity in 3%H₂O moistened O₂ among BZCYYb, YSZ (Y₂O₃-stabilized ZrO₂), GDC (Gd₂O₃-doped CeO₂) and BZCY; it is evident that BZCYYb shows the highest conductivity below 750 °C.

Fig. 3 Ionic conductivities of BZCYYb, BZCY, GDC, and YSZ as measured at 400° to 750 °C in wet oxygen (with ∼3 vol % H₂O).²³

The fuel cell performance with BZCYYb as an electrolyte was further demonstrated with a high faradaic efficiency (90–98%) and >97% overall electric-to-H₂ energy conversion efficiency (based on the lower heating value of H₂) at a current density of 1 A cm⁻².²² With a co-fed CO₂ and H₂O, the same cell was shown with a round-trip efficiency (electricity-to-H₂-to-electricity) >75% and stable operation at a degradation rate of <30 mV per 1000 h.

A similar composition, BaZr_0.4Ce_0.4Y_0.1Yb_0.1O₃ (BZCYYb4411), has been tested as a H⁺-conducting electrolyte in a solid oxide cell setting with a mixed conductor PrBa_0.5Sr_0.5Co_1.5Fe_0.5O_5+δ (PBSCF) as the air electrode, and a composite of Ni–BZCYYb4411 as the fuel electrode, see Fig. 4(a) and (b).³¹ The low-temperature solid oxide fuel cell (SOFC) performance, see Fig. 4(c), is outstanding and the long-term steam electrolysis performance under 500 °C and constant 1.3 V, see Fig. 4(d), is also high and stable cells, further demonstrating the potential of BaCeO₃–BaZrO₃-based materials as a good H⁺ conductor.

Fig. 4 (a) Cross-sectional view of 15 μm thick, dense electrolyte BZCYYb4411 and a 15 μm thick, porous PBSC air electrode supported on a porous fuel electrode. (b) Microstructure of the PBSCF air electrode. (c) V–I and P–I curves under fuel cell modes. (d) Long-term steam electrolysis performance at 550 °C and a fixed 1.3 V.³¹

Despite the broad interest in perovskite H⁺ conductors, practical applications of these conductors in large and thin film formats have been severely limited due to the poor mechanical strength, high reactivity with adjacent components, and difficult processing conditions. Therefore, how to implement the BaCeO₃–BaZrO₃-based H⁺-conductors into solid oxide cell (SOC) stacks remain an engineering challenge.

2.2 Perovskite proton conductors mixed with molten carbonate

To overcome the chemical stability issues and manufacturing difficulties associated with the BZY H⁺-conductor, a mixed proton and carbonate ion conductor (MPCC) comprising of a pre-sintered proton conductor BaZr_0.8Y_0.2O_3−d (BZY) skeleton infiltrated with molten carbonate (MC) phase was reported.³² The H⁺-conductivity of MPCC reached 0.33 S cm⁻¹ at 600 °C in 3% H₂O–air and and 0.38 S cm⁻¹ in 3%H₂O–H₂, which is roughly two orders of magnitude higher than pure BZY, see Fig. 5(a). With the aid of MC, not only is the H⁺-conductivity of BZY improved, but also a fully dense microstructure is achieved by the MC phase.

Fig. 5 Plots of effective ionic conductivity of BZY-(Li_0.62K_0.38)₂CO₃ (MC) (a) vs. temperature in 3%H₂O–air and 3%H₂O–H₂.³² Migration pathways for H⁺-migration along the Y direction in MPCC, (b) side view and (c) top view. The migration pathways are marked with arrows.³³

To explain the conductivity enhancement at high temperature (e.g. T = 600 °C), a new mechanism based on defect reactions involved in the BZY and MC phases was proposed:

At the BZY/MC interface, simultaneous H⁺ conduction takes place in the bulk BZY and along the interface of BZY/MC, where extensive H⁺ exchange and transfer can occur. To support the above hypothesis, a density functional theory (DFT) was conducted on the system to map out the pathways and energetics of H⁺ migration in BYZ/MC.³³ The calculations explicitly show that the H⁺ migration prefers a curved pathway along the ZrO₂-terminated (110) surface in BaZrO₃, see Fig. 5(b) and (c). However, the MC phase provides an alternate, more conductive pathway for H⁺ to migrate along the surface of BaZrO₃, thus facilitating the H⁺ exchange between BaZrO₃ and MC phase across the interface. The study modeled H⁺ migration in CO₃²⁻, Li₂CO₃ crystals, and (Li₂CO₃)₈ clusters and show that the H⁺ transfer in (Li₂CO₃)₈ clusters, a surrogate of MC, rather than in solid carbonates, is fast, confirming the feasible and fast H⁺ transport in the MC phase.³⁴ Overall, the experimental and theoretical results suggest the potential application of dual-phase BYZ/MC MPCC as an electrolyte for intermediate-temperature solid oxide cells.

2.3 Ruddlesden–Popper oxides

The Ruddlesden–Popper (RP) oxides refer to the family with a general formula of A_n+1B_nO_3n+1, where n ≥ 1; A is a rare-earth or an alkaline-earth element, and B is a transition metal.³⁵ The formula can also be otherwise viewed as (AO)(ABO₃)_n, where n is the number of connected layers of corner-sharing BO₆ octahedra. The structure is built of consecutive perovskite (ABO₃)_n blocks alternating with rock-salt AO layers along the c crystallographic axis, see Fig. 6.

Fig. 6 Idealized representation of the crystal structures of Ruddlesden–Popper (RP) phases.³⁶

Ln₂NiO_4+δ (Ln = La, Nd, and Pr) is an n = 1 RP compound. They have recently attracted attention due to their low thermal expansion coefficients (TECs), high oxygen bulk diffusion and surface exchange coefficients, which makes them a good candidate for SOCs. While being oxide-ion conductive, Ln₂NiO_4+δ is also conductive to protons and electrons, thus expanding the active sites for electrochemical reactions.³⁷Fig. 7 shows the possible crystallographic sites for H₂O insertion in Ln₂MO_4+δ (Ln = rare earth element) to produce protons via the following defect reaction:³⁶

Here, represents H attached to the interstitial O and represents H attached to the lattice O. This mechanism is different from perovskite proton conductors mentioned above, where H₂O reacts with to produce . Here in the structure of Ln₂NiO_4+δ, the interstitial O () is a prevalent point defect. Therefore, it is envisioned that H⁺ hops between and to enable H⁺ conduction. While the above mechanism still needs experimental verification, many studies agree that the H⁺ conducting mechanism is different between RP and traditional perovskite oxides.

Fig. 7 The crystal structure of Ln₂MO_4+δ with the arrow indicating the possible water insertion sites.³⁶

The applications of RP materials in H⁺-conducting solid oxide electrolysis cells (SOECs) have been tried in laboratory studies. The results are sometimes contradicting. For example, an H⁺-SOEC based on BaZr_0.2Ce_0.6Y_0.2O_3−δ (BZCY) electrolyte and Pr₂NiO_4+δ (PNO) anode was shown with good chemical compatibility and catalytic activity towards water splitting.³⁷ In another study on Ln₂NiO_4+δ (Ln = La, Pr, Nd) in steam and CO₂ environments, however, notable secondary hydroxides and carbonates were detected on the surface after being exposed to high steam pressure (40 bar) at 550 °C.³⁸

Despite numerous early works, the H⁺ transport mechanisms and properties in RPs are not well understood and attained as perovskite counterparts.³⁹ More comprehensive and in-depth studies are, therefore, needed for understanding H⁺ conducting mechanisms and measuring H⁺ conductivity in RP oxides in the future.

2.4 Pyrochlores

Pyrochlore oxides with the composition of A₂B₂O₇ have also been found to be proton conductive.^40–42 The acceptor doped A₂B₂O₇ have been shown with improved proton conductivity over the undoped A₂B₂O₇.⁴³ The doped La_1.95Ca_0.05Zr₂O₇ exhibits a total conductivity of 6.8 × 10⁻² at 600 °C, see Fig. 8(a), in which partial H⁺ conduction is dominating.⁴³ The oxygen vacancies created by the acceptor doping are the vehicles for proton conduction as illustrated by reaction (1). MD simulations support the necessity of acceptor doping in A₂B₂O₇ on either A- or B-site for long-range H⁺ conduction.⁴⁴Fig. 8(b) shows the structure of La₂Zr₂O₇, in which Zr ions (invisible in Fig. 8(b), embedded inside the octahedra) are located at the centers of the distorted octahedra and O ions (small, white balls) at their corners. The La ions (large, dark-shaded balls) together with O′ (medium-size, white balls) constitute an interpenetrating chain-like network. The suggested proton pathway is illustrated in Fig. 8(c) by following (ii) → (ii) → (v) → (v) → (ii) → (ii) → …. Notably, only sites (ii) and (v) must be visited to form a continuous trajectory. The (ii) to (ii) separation is 1.54 Å, whereas (v) to (v) and (ii) to (v) are 0.76 Å and 1.22 Å, respectively. It is also found that negatively charged dopants are associated with H⁺, and consequently, the proton diffusivity is hindered in their vicinity. However, the binding energy of the associate is quite small in the case of Sr and Ca, making them a reasonable conductor.⁴⁴

Fig. 8 (a) Arrhenius plot of conductivity of (La_1.95Ca_0.05)Zr₂O₇ in wet H₂, (■) σ_H, (•) σ_O and (○) σ_total represent partial proton, oxide-ion conductivity, and total conductivity, respectively;⁴³ (b) La₂Zr₂O₇ structure; (c) suggested proton conduction pathway.⁴⁴

Several applications using pyrochlores as electrolytes in fuel cells, hydrogen sensors, hydrogen separation membranes, and electrochemical ammonia synthesis have been demonstrated. However, the major issue for pyrochlores is the low ionic conductivity, high sintering temperature (>1600 °C), making their practical use in SOCs less attractive.

2.5 Rare-earth ortho-niobates and ortho-tantalates

The rare-earth ortho-niobates and ortho-tantalates refer to AMO₄ (A = rare earth; M = Nb, Ta). With 1% Ca acceptor doping, RE_1−xA_xMO₄ exhibits prominent proton conductivity in wet atmospheres below 800 °C. The highest proton conductivity was observed in 1% Ca-doped LaNbO₄, which was on the order of 10⁻³ S cm⁻¹ at 800 °C, see Fig. 9(a).⁴⁵ However, this class of material experiences monoclinic fergusonite-type structures to tetragonal scheelite-type structures as temperature increases. The proton conductivity suffers a sudden drop as the structure is transformed from a tetragonal to a monoclinic structure, see Fig. 9(a).

Fig. 9 (a) Protonic conductivities of 1% Ca-doped RENbO₄;⁴⁵ (b) proton migration paths with low potential barriers in the tetragonal phase of LaNbO₄. Proton migration pathways are labelled as 1–4 in the figure.⁴⁷

In addition, the proton mobility in rare-earth ortho-niobates and ortho-tantalates is found to decrease with decreasing ion size, which is understandable because it leads to decreased lattice parameter, the polarizability of the lattice and proton-hosting oxygen sublattice. Typically, the smaller, stiffer lattices give rise to a less dynamic oxygen lattice, leading to its lowered momentum to match proton transfer. The proton mobility is also thought to be favored by the symmetry of the lattice.⁴⁶

Using DFT method, the in-plane H⁺-migration pathways in tetragonal LaNbO₄ are identified, see Fig. 9(b), in which H⁺ diffuses by an inter-tetrahedral mechanism within a cation sublattice.⁴⁷ Briefly, pathways 1 and 2 show the proton rapidly shift position between two adjacent a sites and b sites, respectively, with a comparable potential barrier (0.04 eV and 0.06 eV, respectively) for so-called “oscillatory proton transfer”. Pathway 3 is the additional migration path connecting a and b sites for long-distance proton migration within the plane with the third lowest potential barrier of 0.41 eV. It is also found that the proton rotates around the oxide ion before and after it jumps between the two tetrahedrons, i.e. rotation and hopping, which has also been found computationally in other proton conducting oxides, such as BaCeO₃ and BaZrO₃. For the out-of-plane proton migration, pathway 4 describes proton migration between two layers with a higher potential barrier of 0.74 eV. So far, only a few literature are available on the proton conductivity in rare-earth ortho-niobates and ortho-tantalates, leaving great space for future exploration of this class of materials.

3. Solid polyanionic salts as proton conductors

3.1 Sulfates

Solid sulphates such as Li₂SO₄, Li₂Mg(SO₄)₂, LiAgSO₄, and Rb₂SO₄ were first found to be proton conducting in the early 1970s.⁴⁸ Among these, Li₂SO₄ exhibits proton conductivity in its α-phase stable between 577 °C (first-order transition temperature) and 860 °C (melting temperature). The proton conducting α-phase has a face centered cubic structure with a space group of Fm3m, see Fig. 10(a). It was later discovered that the first-order transition temperature of the proton-conducting phase could be reduced to ∼450 °C by adding a small amount of Al₂O₃ into Li₂SO₄.^50,51

Fig. 10 (a) Crystal structure of the face-centered cubic α-Li₂SO₄, where SO₄ tetrahedron and its rotation pushing out of Li⁺ and H⁺ are illustrated.⁴⁹ (b) Arrhenius plot of proton conductivity of Li₂SO₄-45 mol%Al₂O₃ (reproduced from ref. 50).

The mechanism of proton conduction follows the paddle wheel model,⁴⁹ in which steps of proton-attachment to SO₄²⁻, i.e. (H-SO₄)⁻, rotation of SO₄²⁻, and pushing-out of H⁺/Li⁺ are considered, see Fig. 10(a). The proof of proton conductivity was provided by hydrogen concentration cells. The cell with Li₂SO₄–Al₂O₃ as the electrolyte showed zero EMF (electromotive force) in O₂ atmospheres, but proportional to the logarithm of partial pressure of H₂.⁵⁰ A high proton conductivity in the order of 10⁻² to 10⁻¹ S cm⁻¹ has been observed in α-Li₂SO₄, see Fig. 10(b).⁵⁰

3.2 Nitrates

Like solid sulphates, solid nitrates (MNO₃, M = Li, Na, K, Rb, Cs, etc.) were also found to exhibit high proton conductivity. It was also found that proton conductivity increases with crystallinity of nitrates and cationic radius as a result of the soft M–O bonds and open crystal structures.⁵² To stabilize and increase the conductivity of MNO₃, Al₂O₃ is often added. The conductivities of RbNO₃–Al₂O₃ and CsNO₃–Al₂O₃ measured by hydrogen concentration cells are comparable to Li₂SO₄–Al₂O₃, in the order of 10⁻² to 10⁻¹ S cm⁻¹ above 400 °C, see Fig. 11(a).⁵³ For face-centered cubic RbNO₃, for example, protons migrate through the interfacial coordinative mechanism between RbNO₃ and Al₂O₃, see Fig. 11(b).

Fig. 11 (a) Arrhenius plot of proton conductivity for CsNO₃- and RbNO₃–Al₂O₃ composites (reproduced from ref. 52); (b) a schematic showing bonding between the M⁺ and NO₃⁻ may be broken, which provides a coordination site for the proton. Thus, both Rb⁺ and H⁺ may stay at the interface between two phases in the RbNO₃–Al₂O₃ system; here a, b, c, and d represent four non-equivalent M-sites around NO₃⁻ (reproduced from ref. 53).

The major drawbacks of solid sulphates and nitrates, however, are their weak mechanical strength and low chemical and electrochemical stability, as they could decompose at higher temperatures if not properly stabilized. For instance, nitrates may decompose at a temperature >350 °C,⁵⁴ while sulphates may undergo reduction to form H₂S and H₂O in H₂-containing atmospheres.⁵⁵ These drawbacks have limited the practical applications of solid sulphates and nitrates.

3.3 Pyrophosphates

Compared to sulphates and nitrates, phosphates or complex heteropolyacids, such as pyrophosphates, have a better chemical stability. An early study by Zhu et al. suggests that K₃PO₄–Al₂O₃ with a cubic α-phase similar to sulphates and nitrates has a proton conductivity of 10⁻³ to 10⁻² S cm⁻¹ at a temperature above 550 °C.⁵⁰ A series of studies have also shown that Sr or Ca doped rare earth phosphates (LnPO₄, Ln = La, Ce, Pr, Nd, and Sm) exhibit high proton conductivity under wet reducing conditions.^56–60

The protons’ conduction mechanisms in doped phosphates are generally understood as follows:⁵⁸

where represents Sr dopant on the La-lattice; represents (P₂O₇)⁴⁻ residing on the (PO₄)³⁻ lattice; represents proton species residing on the (PO₄)³⁻ lattice. Assisted by ¹H and ³¹P magic angle spinning-nuclear magnetic resonance (MAS-NMR) and Fourier transform (FT)-Raman spectroscopy measurements, the formation of pyrophosphate ions at adjacent orthophosphate ions (PO₄)³⁻ induced by the doping of Sr²⁺ into La³⁺ has been confirmed, which becomes the active site for water incorporation and the formation of (HPO₄)²⁻ at (PO₄)³⁻ sites, resulting in proton conduction.

The conductivities of 1 mol% Sr-doped LaPO₄ and CePO₄ reached 10^−5.2–10^−3.5 S cm⁻¹ and 10^−5.2–10^−3.4 S cm⁻¹, at 500–925 °C, respectively, under H₂O/H₂ and D₂O/D₂ atmospheres (pH₂O = 4.2 kPa and pH₂ = 3 kPa); the results are shown in Fig. 12.⁵⁶ The lowered conductivity in D₂O/D₂ indirectly confirmed the nature of proton conduction in these materials.

Fig. 12 Conductivities of undoped and 1 mol % Sr-doped (a) LaPO₄ and (b) CePO₄ under wet reducing conditions; pH₂O or pD₂O = 4.2 kPa, and pH₂ or pD₂ = 3 kPa.⁵⁶

In addition to phosphates, solid acid pyrophosphates TP₂O₇ (T = Sn, Ti, Si, Ge, Ce, and Zr) have also been found to exhibit proton conductivity in the order of 10⁻³ to 10⁻² S cm⁻¹ within 100 to 400 °C under anhydrous conditions.^61–64 The pyrophosphates typically have a face centered rock salt cubic structure with TO₆ and P₂O₇ units sharing corners, see Fig. 13(a). Moreover, Nagao et al. reported that partial substitution of In³⁺ for Sn⁴⁺ in SnP₂O₇ can improve the proton conductivity from 5.56 × 10⁻² to 1.95 × 10⁻¹ S cm⁻¹ at 250 °C, see Fig. 13(b).⁶⁵ By studying the H/D isotope effect, the proton conduction in TP₂O₇ was proposed to follow the hopping mechanism, i.e., migration via dissociation of O–H bonds. Moreover, it was also found that the presence of P₂O₇ ions has a deficiency in the lattice of TP₂O₇, leading to a lower proton conductivity than undoped pyrophosphates, due to the decrease in proton mobility.

Fig. 13 (a) Crystal structure of TiP₂O₇,⁶⁶ TiO₆ octahedra and PO₄ tetrahedra are shown in blue and green in color, respectively, and (b) Arrhenius plot of the conductivity of Sn_1−xIn_xP₂O₇ in dry air (pH₂O = ∼0.0075 atm).⁶⁵

The improved stability of phosphates and pyrophosphates, compared to solid sulphates and nitrates, makes them more practical. Recently, the use of pyrophosphates composite electrolytes has drawn attention due to their stability in a wider voltage and temperature range. For example, Hibino et al. reported that the Sn_0.95Al_0.05H_0.05P₂O₇ (SAPO)-polytetrafluoroethylene (PTFE) composite shows a high proton conductivity of 0.02 S cm⁻¹ at 200 °C with a highly condensed H₃PO₄ electrode ionomer in a supercapacitor configuration. Such an electrolyte also exhibits a wide voltage range of ±2 V. The resultant performance of such a supercapacitor reaches an energy density of 32 W h kg⁻¹ at 3 A g⁻¹ and stable 7000 cycles from room temperature to 150 °C.⁶⁷ Given their relative low cost and ease of preparation, they remain an attractive class of proton conductors for future consideration.

3.4 Solid acids

A group of solid acids, MHAO₄ and M₃H(AO₄)₂ with monovalent alkali metal cations M = Rb, Cs, NH₄ and tetrahedral oxyanions AO₄ (A = S, P, Se) exhibits super proton conductivity above a phase transition temperature between 50 and 250 °C even in a dry atmosphere. This class of material was first reported by Baronov et al.,⁶⁸ in which high proton conductivity (10⁻³–10⁻² S cm⁻¹ at 140–230 °C) was observed in CsHSO₄ and Rb₃H(SeO₄)₂, see Fig. 14(a). The high proton conductivity was attributed to the unique proton disorder in H-bond networks. CsH₂PO₄ was found later to have a higher transition temperature (230 °C) and decomposition temperature (375 °C in high humidity) than CsHSO₄ (140 °C and 200–230 °C, respectively) and exhibit a higher proton conductivity of 2.2 × 10⁻² S cm⁻¹ at 200 °C and 4.0 × 10⁻² S cm⁻¹ at 240 °C, respectively.⁶⁹ The proton conduction mechanism in these acids follows successive intra-bond jumps, reorientations of the O–H group and breaking of an old H-bond and forming of a new one, see Fig. 14(b).⁷⁰

Fig. 14 (a) Arrhenius plot of proton conductivity of CsH₂PO₄, Rb₃HSeO₄, and CsHSO₄ (reproduced from ref. 55) and (b) illustration of two proton diffusion movements in CsH₂PO₄, confirmed via¹H pulsed field gradient NMR, ¹H NMR spin-lattice relaxation time (T_l), and quasi-elastic neutron scattering measurements.⁷⁰

The potential of phosphate-based solid acids for application in fuel cells has been shown by Haile et al.⁵⁵ With a thick CsHSO₄ membrane (thickness: ∼1.5 mm) as the electrolyte, the work showed a promising OCV of 1.11 V, and output of 44 mA cm⁻² at 150 °C. With a 25 μm thick CsH₂PO₄ as an electrolyte film, the fuel cell produced a peak power density of 415 mW cm⁻² at 240 °C under humidified (pH₂O = 0.3 atm) H₂ and pure O₂.⁶⁹ However, a further application of solid acids in fuel cells is limited by the strict operating conditions (e.g. high humidity and narrow temperature range). Similar to pyrophosphates, a combination of solid acids with a metal–organic coordination framework shows promising conductivity and stability, especially under low humidity. For example, Ponomareva et al. reported that a composite of CsH₅(PO₄)₂ and a metal–organic coordination framework, Cr-MIL-101 matrix, exhibits a conductivity of ∼10⁻² S cm⁻¹ at 130 °C at RH = 10–15%.⁷¹ The high proton conductivity under low RH suggests that the hybrid solid acids composite could find potential applications in fuel cells and H₂ pumps.

4. Hydrates as proton conductors

The materials that become proton conductive when they are in a fully hydrated state are classified as proton-conducting hydrates. Since this class of materials relies upon the degree of adsorbed water to conduct protons, their operating temperature is typically low.

4.1 Mesoporous nanocomposites

Some mesoporous inorganic materials with nanosized pores in 2 to 50 nm become proton conductive due to physically adsorbing water molecules. Yamamda et al. reported an inorganic TiO₂–P₂O₅ mesoporous nanocomposite with a proton conductivity of 10⁻² S cm⁻¹ at 160 °C, under fully saturated humidification conditions, see Fig. 15.⁷² The high proton conductivity is attributed to the strong P-POH chemisorption enabled by P₂O₅ under high humidity conditions.⁷³

Fig. 15 (a) Proton conductivity of (○) TiO₂–P₂O₅ mesoporous nanocomposites and (□) mesoporous silica measured from the RT to 160 °C under 100% relative humidity (RH) conditions. (b) Arrhenius plots of the conductivity of the TiO₂–P₂O₅ mesoporous nanocomposites. A solid line is the least-squares fitting. The activation energy (E_a) of the proton transport was estimated from the slope. Reprinted with permission from J. Am. Chem. Soc., 2005, 127(38), 13092–13093. Copyright (2007) American Chemical Society.⁷²

Similarly, Marschall et al. reported that highly ordered mesoporous silicas with sulfonic acid groups (SO₃H) when combined with Nafion to form a functional group (Si-MCM-41, Mobil Composition of Matter No. 41) exhibited very high proton conductivity up to 0.2 S cm⁻¹ under fully saturated humidification conditions.⁷⁴ The high proton conductivity in the presence of water is derived from the Grotthus mechanism, see Fig. 16.⁷⁵ In this mechanism, the proton transport is realized by protons hopping from one water molecule to another. To a smaller extent, the diffusion of H₃O⁺ also enhances proton transport. At higher temperatures, hopping and diffusion become faster since the rotation and vibration of anchoring propyl chains increase, allowing the SO₃H groups at the end of the chains to encounter each other more freely, thus facilitating the direct proton transport. A recent study by Vinothkannan et al. shows substantially high proton conductivity (45.1 mS cm⁻¹) under RH = 50% and 120 °C for the optimized Nafion/SO₃H-1 wt% unzipped graphite nanofiber (UGNF) membrane.⁷⁶

Fig. 16 (a) Proton conductivity of 40% SO₃H-MCM-41 (mw) at 0% (◆), 50% (○), and 100% (●) RH, Reprinted with permission from Chem. Mater. 19, 26 (2007): 6401–6407. Copyright (2007) American Chemical Society.⁷⁴ and (b) simplified schematic representation of proton transport via the surface mechanism, Grotthuss mechanism, and vehicle mechanism.⁷⁵

In contrast to Nafion fouls, for which the proton conductivity decreases drastically above 100 °C due to the water loss, water seems to be kept inside the channels of mesoporous nanocomposites due to the geometry of the pores which fixes the sulfonic acid groups, thus maintaining proton conductivity even at higher than 100 °C. It was also found that the proton conductivity increased with the extent of substitution of metasilicate with MCM. The highest proton conductivity of 0.2 S cm⁻¹ was achieved with a 40% SO₃H-MCM-41 (mw) sample due to the high pore filling with SO₃H groups and the absence of parasitic pores.

4.2 Heteropolyacid hydrates

Early studies have found that some heteropolyacids in a hydrated state, such as H₄SiW₁₂O₄₀·28H₂O, H₃PW₁₂O₄₀·29H₂O, and H₃PMo₁₂O₄₀·29H₂O, exhibit high proton conductivity.^77–79 The proton conductivities are in the order of 10⁻² and 10⁻¹ S cm⁻¹ at room temperature. Large and globular anions (e.g. [Mo₁₂PO₄₀]³⁻) with Keggin structures are present in these hydrates, with crystal water showing violet thermal motion which can be easily dehydrated and rehydrated. These crystal waters contribute to the high proton conductivity, but only within a narrow window. Temperature and humidity are key factors in maintaining stable hydrates, where 80 °C and 70–80% RH are typically required conditions. In addition, pelletization or other processing methods can lead to the loss of crystallization in these hydrates. Thus, some porous oxides such as SiO₂ and Al₂O₃ have been often used as a matrix to reinforce the stability and robustness of heteropolyacid hydrates, which makes their application in solid electrolytes, sensors, and electrochromic displays possible.^80,81

4.3 Layered hydrates

Another group of hydrates studied as a proton conductor has a layered structure. Examples include HUO₂PO₄·4H₂O,⁸¹ α-Zr(HPO₄)₂·nH₂O,⁸² γ-Zr(PO₄)(H₂PO₄)·2H₂O,⁸³ and γ-Zr sulfoaryl phosphonates.⁸⁴ They are acid groups that exhibit high proton conductivity of 10⁻⁴ to 10⁻² S cm⁻¹ over a wider temperature range, but thermally more stable than heteropolyacids hydrates. In these compounds, a phase transition occurred above ∼50 °C,⁸⁵ resulting in a significant increase in proton conductivity, due to the formation of an interlamellar H-bonded network with abundant H₃O⁺ and high concentration of H⁺.

The conduction mechanism is similar to that in mesoporous nanocomposites, with a hybrid Grotthuss-vehicle mechanism involving an intermolecular transfer (hopping) and an intramolecular transfer of protons. The former could be facilitated by the high concentration of H₃O⁺ in the structure, and the latter is most likely facilitated by the high concentration of H-bond vacancies.

Some oxides, such as Sb₂O₅,⁸⁶ SnO₂,⁸⁷ V₂O₅, ZrO₂⁸⁸ and VO₂,⁸⁹ are also found to be proton conductive in their hydrated states. Proton conductivities of 10⁻⁴ to 10⁻² S m⁻¹ are found below 300 °C. Typically, maintaining saturated water vapor pressure is required to achieve sustainable proton conductivity. The proton conduction in such hydrates typically follows a pathway involving first, rotation around the O atom, and then migration between O atoms, see Fig. 17.

Fig. 17 Migration pathways for H⁺ in H-VO₂.⁸⁹

Despite the high proton conductivity of all these hydrates, their practical application is still limited due to the difficulty of retaining water and poor thermal stability during operation.

5. Hydrides as hydride-anion conductors

In chemistry, a hydride is formally an anion of hydrogen, H⁻. Therefore, hydrides often refer to compounds and ions in which hydrogen is covalently attached to a less electronegative element. In such cases, the H center has a nucleophilic character, which contrasts with the protic character of acids.

5.1 Hydrides

Early studies revealed that some alkaline earth metal hydrides, such as CaH₂ and SrH₂, exhibit hydride anion (H⁻) conductivity up to 10 S cm⁻¹ above 780 °C,⁹⁰ and application of such materials as electrolytes for fuel cells was also reported.⁹¹ More recently, a study on BaH₂/BaD₂ by Verbraeken et al. suggested that the high proton conductivity was attributed to H⁻ and achieved in a Ni₂In-type structure transformed from a cotunnite structure at ∼470 °C, where a sharp increase in conductivity was observed.⁹² The phase transition from cotunnite to high-symmetry Ni₂In type structures was also promoted by high pressure. A large concentration of hydride vacancies was also observed in the structure with a composition of roughly BaH/D_1.84, which contributes to the proton conductivity. In addition, small Ba-to-H or deuteride bond length, high isotropic temperature factors for D1 and D2 (7.6 and 11.6 Å), small mass and high mobility of D ions, see Fig. 18(a), all give rise to the superionic conductivity in BaD₂. The conductivity of BaH₂ is in the range of 0.04–0.2 S cm⁻¹ between 470 and 630 °C, see Fig. 18(b).

Fig. 18 (a) Crystal structure of BaD₂, and (b) Arrhenius plot of hydride-anion conductivity of BaH₂.⁹²

Despite the high hydride anion conductivity, the narrow operation window largely restricts the practical application of BaH₂. More recently, Ubukata et al. reported a solid solution of Ba_2−δH_3−2δX (X = Cl, Br, or I) with a layered anion-ordered structure, which exhibited excellent hydride conductivity at low temperatures (e.g. 1 mS cm⁻¹ at 200 °C).⁹³ The layered anion ordering and Schottky defects in such materials likely suppress the structural transition commonly occurring during cooling, thus retaining a highly symmetric hexagonal lattice at lower temperatures for H⁻ conduction. Furthermore, AELiH₃ (AE: alkaline earth metals) perovskites are also found to exhibit high H⁻-conduction at low temperatures (e.g. 5.0 × 10⁻⁶ S cm⁻¹ at room temperature for Sr_0.925Na_0.075LiH_2.925),⁹⁴ implying that the employment of aliovalent cations on the A site in perovskite-type hydrides leads to the formation of hydrogen vacancies, thus enabling fast H⁻-conduction even at very low temperatures. These new discoveries widen the potential of hydrides for applications over a wider temperature range.

5.2 Perovskite and fluorite-type oxyhydrides

Perovskite and fluorite-type oxyhydrides containing H⁻ in the lattice, such as La_2−xSr_xLiH_1+xO_3−x,⁹⁵ Ba₂ScHO₃,⁹⁶ LaSrCoO₃H_0.7,⁹⁷etc., have been shown to conduct H⁻ in a similar manner to the above BaH₂ hydrides. Kobayashi et al. reported that rare-earth lithium oxyhydrides of La_2−x−ySr_x+yLiH_1−x+yO_3−y conduct H⁻ over oxygen sublattices. One of the promising compositions, La₂LiHO₃, has been successfully synthesized into a K₂NiF₄-type structure with an orthorhombic structure at 923 K and 2 GPa for 30 min. In this structure, both H⁻ and O²⁻ are ordered on the axial anion sites, see Fig. 19(a). The conductivity of La_2−xSr_xLiH_1−xO₃ series (0 ≤ x ≤ 0.2, y = 0) is shown in Fig. 19(b), suggesting that Sr-induced vacancies increase H⁻ conductivity. With La₂LiHO₃ as the electrolyte, a cell consisting of Ti/o-La₂LiHO₃/TiH₂ has been constructed and tested; the results are shown in Fig. 19(c).⁹⁵ The cell exhibited an initial open circuit voltage of 0.28 V at 300 °C, consistent with the theoretical value of a Ti/TiH₂ redox couple. Under a constant discharge current of 0.5 mA, a steep drop-off in the cell voltage from 0.28 to 0.06 V was observed for the first reaction step, corresponding to the increase in the H⁻ content at the anode, followed by the cell voltage decreasing gradually to 0.0 V as H⁻ was continued added into Ti, implying the reaction between H and Ti was completed.

Fig. 19 (a) The crystal structure of o-La₂LiH₃O₃, (b) Arrhenius plot of the ionic conductivity of La_2−xSr_xLiH_1−xO₃ (y = 0, 0 ≤ x ≤ 0.2), and (c) discharge curve of cell Ti/o-La₂LiHO₃/TiH₂. The inset shows the cell configuration and the proposed electrochemical reactions.⁹⁵

The synchrotron XRD indicates that the La₂LiHO₃ electrolyte is stable when in contact with the Ti and TiH₂ electrodes during the reaction, while phase changes were detected for both the cathode and anode. In the cathode, the initial δ-TiH₂ phase was found to transform into α-Ti upon releasing hydrogen; the corresponding lattice shrinkage was confirmed by the shift to a higher 2θ in the XRD pattern.

A more recent work shows that Ba–Li oxyhydride, Ba_1.75LiH_2.7O_0.9, synthesized under ambient pressure with a K₂NiF₄-type structure exhibited a high H⁻ conductivity (∼10⁻² S cm⁻¹ at 300 °C) after an order–disorder transition.⁹⁸ Overall, H⁻ conductivity in such oxide-based framework structures is found to be almost independent of temperature but closely related to the anion ordering and vacancies.

Other than the H⁻-conducting oxyhydrides, mixed-electronic and ionic conductivity was also found in some oxyhydrides such as BaTiO_3−xH_x.⁹⁹ The high electronic conduction (e.g. 10⁻⁴ S cm⁻¹ at 25 °C in BaTiO_2.4H_0.6) might find these materials as electrodes in an electrochemical cell.

Fluorite-type oxyhydrides are another group of hydrides with high H⁻ conductivity comparable to perovskite oxyhydrides. LaH_3−2xO_x (x = 0.54) demonstrates the highest H⁻-conductivity of 2.6 × 10⁻² S cm⁻¹ at T = 342 °C, see Fig. 20.¹⁰⁰ The crystal structure of LaD₂O_0.5, an isotopic analogue of LaH₂O_0.5, is tetragonal, as determined by neutron powder diffraction (NPD). The light mass and large polarizability of H⁻ and the framework comprising of densely packed H⁻ in LaH_3−2xO_x are crucial factors that enable significant temperature dependence of the H⁻ conductivity.

Fig. 20 (a) Crystal structure of LaD₂O_0.5 determined by NPD. Solid and dashed lines represent the unit cells of the FCC and tetragonal structures, respectively. Green, pink, yellow, and grey colors denote the fractions of La³⁺, D⁻, O²⁻, and vacancies at each crystallographic site, respectively, and (b) Arrhenius plots of LaH_3−2xO_x and H⁻ conductivity measured by AC impedance spectroscopy.¹⁰⁰

The successful demonstration of an all-solid-state electrochemical cell using oxyhydrides as an electrolyte presents the possibility of developing solid electrochemical devices based on H⁻ conductors. Although recent studies have shown the possibility of ambient pressure synthesis of oxyhydrides, the requirement for high pressure in synthesizing the desired oxyhydrides and the lack of stability study are major obstacles for practical applications of oxyhydrides.

5.3 Nitrohydrides and hydrofluorides

Like oxyhydrides, H⁻ diffusion was also observed in nitrohydrides such as Ba₂NH,^101,102 Sr₂NH,¹⁰³ and Sr₂LiH₂N.¹⁰⁴ Studies have discovered that these nitrohydrides exhibit layered structures at high concentrations (up to 15%) of H-vacancies, which favors H⁻ diffusion. The H⁻ migration proceeds in the planes occupied by H⁻, in uncorrelated jumps between regular lattice sites, see Fig. 21(a). The high and pure mobility of H⁻ gives rise to the high conductivity of 1 S cm⁻¹ above certain temperatures (e.g. 525 K for Ba₂NH), see Fig. 21(b). However, there is a lack of systematic study on the origin of such high conductivity, leaving nitrohydrides group wide open for further investigations.

Fig. 21 (a) Layered type structure of Ba₂NH with distorted Ba-octahedra centered alternately by H⁻ and N³⁻, with corresponding distance, (b) Arrhenius plot of H⁻ conductivity in Ba₂NH, and (c) diffusion of hydride anions in Ba₂NH via uncorrelated jumps between regular hydrogen sites.¹⁰¹

Hydrofluorides, for example, CaF_2−xH_x,¹⁰⁵ and LiBaF_3−xH_x,¹⁰⁶ have a perovskite structure, in which ionic radii of H⁻ and F⁻ are similar (r_H⁻ = 2.08 Å, r_F⁻ = 1.33 Å). EMF measurements reveal the high conductivity of LiBaF_3−xH_x in H₂ atmosphere, varying from 10⁻⁵–10⁻² S cm⁻¹ at temperatures of 330–650 °C, see Fig. 22. However, such hydride compounds are subject to oxidation and thus suffered stability issues, making them difficult in applications.

Fig. 22 Arrhenius plots of conductivity of LiBaF_3−xH_x in H₂: □x = 0, ◊x = 0.05, ○x = 0.1, Δx = 0.2, ● LiBa_0.9K_0.1F_2.8H_0.1.¹⁰⁶

6. Summary

In summary, the most investigated H⁺ and H⁻ conductors have been reviewed in this article in terms of structures, conduction mechanisms, conductivities, and optimal conducting conditions. The representative conductors for each class are summarized in Table 1.

Table 1 Representative H⁺/H⁻ conductors and their corresponding conductivities and optimal operating conditions

Material	Conductivity (S cm^−1)	Optimal operating conditions	Ref.
Oxides as H^+ conductors
Perovskites
BaCe0.9Y0.1O3−δ (BCY)	1.8 × 10^−2–7 × 10^−2 at 600–1000 °C	500–750 °C	27
BaZr0.9Y0.1O3−δ (BZY)	1.6 × 10^−3–6 × 10^−3 at 600–1000 °C	500–750 °C	28
BaZr0.1Ce0.7Y0.1Yb0.1O3−δ (BZCYYb)	5 × 10^−3–5 × 10^−2 at 400–750 °C	500–750 °C	23
RP oxides
BaLaIn0.9Ti0.1O4.05	7.4 × 10^−8–2.6 × 10^−6 at 250–450 °C		39
Pyrochlores oxides
(La1.95Ca0.05)Zr2O7−δ	6.3 × 10^−5–1.1 × 10^−3 at 400–900 °C	500–800 °C	41
Rare-earth ortho-niobates and ortho-tantalates
1% Ca-doped LaNbO4	2 × 10^−5–9 × 10^−4 at 400–1000 °C	700–1000 °C	45
Perovskites with MC
BZY-(Li0.62K0.38)2CO3	0.01–0.78 at 400–660 °C	500–650 °C	33
Solid polyanionic compounds as H^+ conductors
Phosphates/pyrophosphates
1 mol% Sr-doped LaPO4	8.7 × 10^−6–2.6 × 10^−4 at 500–930 °C	500–900 °C	56
1 mol% Sr-doped CePO4	8.8 × 10^−6–3.8 × 10^−4 at 500–930 °C	500–900 °C	56
Sn0.9In0.1P2O7	7.1 × 10^−2–1.8 × 10^−1 at 75–300 °C	200–300 °C	65
CsHSO4	1.6 × 10^−7–3.5 × 10^−2 at 100–190 °C	140–190 °C	55
Rb3H(SeO4)2	9.6 × 10^−8–8.3 × 10^−3 at 75–235 °C	180–235 °C	55
CsH2PO4	3.1 × 10^−6–2.71 × 10^−2 at 155–252 °C	200–250 °C, pH2O ≥ 0.4 atm	68
Sulphates and nitrates
Li2SO4-45 mol%Al2O3	3.2 × 10^−5–1.8 × 10^−2 at 350–600 °C	450–600 °C	50
RbNO3–Al2O3	2 × 10^−4–6.6 × 10^−2 at 325–650 °C	500–650 °C	52 and 53
CsNO3–Al2O3	1.9 × 10^−4–9.4 × 10^−2 at 400–650 °C	550–650 °C	52 and 53
Hydrates as H^+ conductors
Mesoporous nanocomposites
TiO2–P2O5	6 × 10^−3–1.5 × 10^−2 at 50–200 °C	50–200 °C, 100%RH	72
40% SO3H-MCM-41	1 × 10^−3–2 × 10^−2 at 60–140 °C	60–140 °C, 100%RH	74
Heteropolyacid hydrates
H4SiW12O40·28H2O	2 × 10^−2 at 25 °C	<100 °C, >70%RH	77
H3PW12O40·29H2O	8 × 10^−2 at 25 °C	<100 °C, >70%RH	78
H3PMo12O40·29H2O	1.7 × 10^−1 at 25 °C	<100 °C, >70%RH	79
Layered hydrates
HUO2PO4·4H2O	5 × 10^−3 at 25 °C	<100 °C, high RH	82
Oxide hydrates
V2O5·nH2O, ZrO2·nH2O	1 × 10^−2 at 25 °C	<150 °C, 100%RH	89
H^− conductors
Hydrides
BaH2	5.4 × 10^−6–1.68 × 10^−1 at 100–630 °C	470–630 °C	92
Sr0.925Na0.075LiH2.925	5.0 × 10^−6 at 25 °C	25–150 °C	94
Perovskite oxyhydrides
La2Sr0.1LiH0.9O3	3.86 × 10^−7–2.04 × 10^−5 at 200–300 °C	200–300 °C	95
Ba1.75LiH2.7O0.9	2.24 × 10^−7–1.38 × 10^−2 at 210–350 °C	300–350 °C	98
LaH1.92O0.54	9.24 × 10^−10–2.6 × 10^−2 at 85–342 °C	250–342 °C	100
Nitrohydrides and hydrofluorides
Ba2NH	3.5 × 10^−6–1.1 at 35–435 °C	350–400 °C	101
LiBaF2.8H0.2	4 × 10^−4–1.7 × 10^−2 at 350–600 °C	350–600 °C	106

---

Overall, despite some results shown in electrochemical cells and H₂ separation membranes, no industrial scale devices have been demonstrated for proton and hydride-anion conductors. The main challenges for H⁺/H⁻ based electrochemical cells to be practical are poor chemical, thermal and mechanical stabilities, difficulty to process and chemical reactivity under operating conditions. To advance H⁺/H⁻ based electrochemical cells, further scientific studies and engineering efforts are greatly needed.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

We would like to thank the National Science Foundation for the financial support under award number 1801284 and the Department of Energy, Office Fossil Fuels, National Energy Technology Laboratory for the financial support under award number DE-FE-0032111.

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