Ordered mesoporous metal oxides for electrochemical applications: correlation between structure, electrical properties and device performance

Abstract

Ordered mesoporous metal oxides with a high specific surface area, tailored porosity and engineered interfaces are promising materials for electrochemical applications. In particular, the method of evaporation-induced self-assembly allows the formation of nanocrystalline films of controlled thickness on polar substrates. In general, mesoporous materials have the advantage of benefiting from a unique combination of structural, chemical and physical properties. This Perspective article addresses the structural characteristics and the electrical (charge-transport) properties of mesoporous metal oxides and how these affect their application in energy storage, catalysis and gas sensing.

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1. Introduction

With the introduction of surfactant-templating methods and the evaporation-induced self-assembly (EISA) process in the 1990s, interest in the preparation of ordered mesoporous materials in powder and thin film form has become widespread.^1–6 While initial efforts focused on siliceous materials, soft-templating methods, in general, lend themselves to the formation of mesostructured metal oxides with nanocrystalline walls.^7–16 The preparation usually relies on the solution-phase co-assembly of either sol–gel precursors or preformed (nanocrystal) building blocks with a surfactant or an amphiphilic polymer structure-directing agent to achieve different pore symmetries, followed by some post-treatment.^13–22 Because of their unique structural and compositional diversity, mesoporous materials are currently drawing increasing attention for energy applications (batteries, catalysis, gas storage and sensing, etc.).^23–27 In principle, an open pore–solid architecture provides good accessibility of the material and its internal structure to the surrounding medium (note that gasses or liquids can readily penetrate into the porosity), thereby offering a large number of active surface sites. In addition, it enables effective transport of charge carriers due to short diffusion pathways, among others, and endows the material with mechanical flexibility. Most importantly, because of the small dimensions of the crystallites in the pore walls and the high interface density, defect chemistry and surface space-charge effects strongly affect the device performance and provide the unique possibility of tailoring the materials properties. In this Perspective, we review and describe the interplay between structure and electrical transport properties of mesoporous metal oxides (mainly as thin films), with a particular focus on applications in electrochemical energy storage, catalysis and gas sensing. Challenges and future prospects are also discussed.

2. Mesoporous metal oxide thin films

2.1. Evaporation-induced self-assembly

The principles of self-assembly for sol–gel-derived mesoporous materials have been described in considerable detail elsewhere and will not be further discussed here.^{13–17,28–32} However, in the following, we briefly describe the EISA-based preparation of mesostructured thin films, which affords a high degree of control over the intrinsic (chemistry) and extrinsic (environment) processing parameters. In the past two decades, many important metal oxides, primarily with pores in the size range between 5 and 30 nm, and with different crystalline phases have been produced by EISA and related methods. However, crystallization of the wall structure with retention of the porosity is challenging. This is due, in part, to the lack of broad availability of suitable structure-directing agents.

In general, the processing scheme shown in Fig. 1 can be described as follows: a solution containing the structure-directing agent and the precursor(s), and possibly an organic additive (e.g., swelling or stabilizing agent), is sprayed, spin- or dip-coated onto a planar substrate. Upon solvent evaporation, a mesostructured inorganic–organic hybrid thin film with short- or long-range periodicity is formed. Finally, the material is made porous by removing the structure-directing agent via extraction or combustion. Note that heating at elevated temperatures is typically used to drive condensation of the inorganic framework, crystallize the pore walls and tailor the grain size.

Fig. 1 Schematic illustration of the fabrication steps for ordered mesoporous thin films using an EISA process. (a) Solution consisting of structure-directing agent and inorganic precursors (shown for ZrCl₄ and YCl₃; chemical reactions with the solvent are omitted for clarity), (b) micelle formation, coassembly and condensation of the inorganic reagents, (c) inorganic–organic hybrid thin film formation upon solvent evaporation and (d) thermally-induced crystallization of the inorganic wall structure and combustion of the organic structure-directing agent.

In general, mesoporous materials have the advantage of benefiting from a unique combination of structural, chemical and physical properties, making them interesting for a variety of applications. Especially amphiphilic polymers as structure-directing agents have been shown well suited for the preparation of crystalline metal oxide thin films with cubic or hexagonal pore structures. In this regard, KLE diblock copolymers [KRATON LIQUID-block-poly(ethylene oxide) or poly(ethylene-co-butylene)-block-poly(ethylene oxide)] seem to somewhat stand out due to their favorable templating properties.^11,12,19,33 However, other di- and triblock copolymers, such as poly(ethylene oxide)-block-poly(propylene oxide)-block-poly(ethylene oxide) (referred to as Pluronic), polyisobutylene-block-poly(ethylene oxide), poly(ethylene oxide)-block-polybutadiene-block-poly(ethylene oxide) or poly(ethylene oxide)-block-polystyrene, have also proven to be robust structure-directing agents.^{7,8,13,16,20,34–37}

A variety of binary and multinary metal oxide thin films with cubic mesoporous morphologies have been synthesized from KLE and alkoxide or salt precursors, including TiO₂, Al₂O₃, CeO₂, HfO₂, MoO₃, SnO₂, SrTiO₃, Li₄Ti₅O₁₂, REVO₄, BiFeO₃ and PZT, to name a few.^{11,12,19,38–49} In the following, we use TiO₂, which is often considered as a prototype material for applications in photocatalysis, solar cells or batteries, as an example.^50–53 Smarsly et al. reported on the preparation of cubic mesoporous TiO₂ thin films from EtOH:THF solution containing TiCl₄ and KLE by the dip-coating method. Crystallization into the anatase phase was found to begin at 450 °C and the nanoscale porosity was retained up to about 700 °C.¹¹ In a related study, Fattakhova-Rohlfing et al. examined the Li-storage properties of KLE- and Pluronic P123-templated mesostructured TiO₂ thin films, which allowed quantifying the fraction of crystalline and amorphous phases in the samples.⁵⁴ They showed that the Pluronic P123-derived films display a strong tendency for anatase formation already well below 450 °C (approx. 50% crystallinity at 400 °C) and that full crystallization at temperatures ≥600 °C is accompanied by collapse of the pore network. This behavior is in stark contrast to that of the KLE-templated material, as schematically shown in Fig. 2. Brezesinski et al. also synthesized cubic mesoporous TiO₂ thin films by a conventional sol–gel route using KLE and TiCl₄ as polymer structure-directing agent and precursor, respectively, and by a particle-based route, with anatase nanocrystals serving as preformed building blocks.¹⁹ The authors demonstrated that the nanocrystal films can withstand much higher temperatures (up to 900 °C) before losing nanoscale order. In addition, they showed much improved Li-storage properties, especially greater power density due to high levels of pseudocapacitive charge storage.

Fig. 2 The effect of annealing temperature on the pore structure of Pluronic- and KLE-templated TiO₂ thin films. Adapted from ref. 54 with permission.

Hartmann et al. compared the photoelectrochemical water splitting properties of cubic mesoporous TiO₂ thin films produced using the same synthetic routes described by Brezesinski and coworkers.²⁰ Both materials differed significantly in the pore walls thickness, see scanning electron microscopy (SEM) images in Fig. 3, despite using the same polymer structure-directing agent in the preparation process. Interestingly, the sol–gel material showed a significantly higher efficiency, which the authors explained by insufficient electronic connectivity of the nanocrystal films. Overall, these results indicate that not only the pore accessibility and specific surface area but also the transport properties have a profound effect on the device performance.

Fig. 3 Top-view SEM images of sol–gel-derived (a and b) and preformed nanoparticle-based (c and d) mesoporous TiO₂ thin films. Reproduced from ref. 20 with permission.

3. Transport properties

As briefly discussed in the previous sections, mesoporous oxides are of interest for a broad range of electrochemical applications. For the functionality of such devices, efficient transport of electrical charge carriers (mobile defects of the ideal crystal structure) is required. As mesoporous materials are typically polycrystalline and made from individual crystallites connected in a 3-dimensional framework, as schematically shown in Fig. 4, they contain a high density of interfaces. Solid/solid (grain boundaries), solid/liquid and solid/gas interfaces may significantly alter the mobility of electronic and ionic charge carriers due to the formation of a space-charge region. In general, the electrical conductivity of a charge carrier is given by:

σ = Zeμn, (1)

where Z is the valence, e is the elementary charge, μ is the mobility and n represents the charge-carrier density contributing to the transport. In the presence of an interface, both the charge-carrier density and mobility may differ from the respective bulk values, making detailed understanding of the effect of interfaces on the electrical transport properties indispensable. Apart from elemental doping, optimization of the electronic properties of mesoporous materials, and therefore of the electrochemical device performance, is feasible by tailoring the interface density, for example, by controlling the pore size, the wall thickness or the size of the crystallites (grains) in the wall structure.

Fig. 4 Schematic illustration of the 3-dimensional architecture of mesoporous oxide thin films with efficient transport pathways for electrons and ions and a vast number of interfaces.

In the following, we will give a short description of both the formation of a space-charge layer at interfaces and the characterization of electronic and ionic conductivity using electrochemical impedance spectroscopy (EIS). The influence of a space-charge layer on the electronic and ionic transport properties has been reported in detail for CeO₂- and ZrO₂-based ceramics when the crystallite size becomes comparable to its extension. The findings for these systems will be briefly summarized, followed by reports on the investigation of transport properties in mesoporous materials. Finally, we also discuss the protonic conductivity in nanoscale oxides.

3.1. Defect chemistry at interfaces

As indicated by eqn (1), the electrical conductivity of a material depends on the concentration and mobility of charge carriers. While the mobility is determined by the electronic and ionic structure of the crystal, the carrier concentration is related to the electrochemical potential of the corresponding defect A, which is given by the sum of the chemical potential and the electrical potential:⁵⁵

[small mu, Greek, tilde]_A = μ⁰_A + Z_AFϕ(x). (2)

In the absence of an electrical field, the electrochemical potential is constant throughout the bulk. However, at an interface, the local symmetry is broken. This means the chemical potential (x = 0) differs from that in the bulk (x = ∞):

μ_A(x = 0) ≠ μ_A(x = ∞). (3)

To ensure constant electrochemical potential, the mobile charge carriers redistribute, resulting in charging of the interface and local variation of the electrical potential ϕ(x), the so-called space-charge region, which compensates for the difference in chemical potential. The redistribution of charge carriers affects the carrier concentration at the interface, and therefore the corresponding electrical conductivity.

For an ideal interface, the change in charge-carrier concentration relative to the bulk concentration c_∞is given by (1-dimensional case):^56–58

where Δϕ(x) = ϕ(x) − ϕ_∞ is the potential relative to the bulk potential ϕ_∞. The distribution of charge carriers can then be calculated by solving the Poisson–Boltzmann equation.⁵⁶

Depending on the mobile charge carriers present in the system, different profiles of the resulting charge-carrier distribution are obtained. The two most prominent cases are Gouy–Chapman and Mott–Schottky,⁵⁹ both of which are schematically shown in Fig. 5.

Fig. 5 Schematic illustration of the space-charge region for (a) Gouy–Chapman and (b) Mott–Schottky case. Adapted from ref. 60 with permission.

In the Gouy–Chapman case, two charge carriers of opposite sign are mobile and contribute to the surface charging. In the space-charge region, the carrier with the opposite sign to the surface charge accumulates at the interface, while the other is depleted. The width of the space-charge region is then given by the Debye length:⁵⁹

with ε₀ and ε_∞ being the vacuum permittivity and permittivity of the material, respectively.

If one of the charge carriers is immobile, that is, its concentration profile is independent of the distance to the boundary, redistribution of mobile charge carriers follows a Mott–Schottky profile. Compared with the Gouy–Chapman case, redistribution results in an extended space-charge region of effective width:⁵⁹

As the space-charge profile can only be calculated analytically for two charge carriers of opposite sign, for mixed cases or for low space-charge potentials, approximations are needed. Göbel et al. presented a numerical approach to calculate the concentration profile for these cases, allowing to obtain more precise analytical solutions of the Poisson–Boltzmann equation.^61,62 However, determination of the space-charge profile by solving the Poisson–Boltzmann equation does not account for interactions between the charge carriers, which can become important for high defect concentrations. In 2015, Mebane and De Souza presented a generalized space-charge model (Poisson–Cahn theory) to account for defect–defect interactions at high defect concentrations.^63,64 While in the dilute case, the results of the analytic model described above are well reproduced, the profile of the space charge is modified for higher defect concentration. At the positively charged interface, the oxygen-vacancy concentration increases significantly (segregation of oxygen vacancies at the interface). Interestingly, depletion of oxygen vacancies in the adjacent space-charge region is accompanied by an enrichment of carriers before the bulk value is reached.

3.2. Electronic and ionic conductivity

The effect of a space-charge layer on the electronic and ionic transport is reasonably well understood for polycrystalline CeO₂- and ZrO₂-based ceramics. In both material systems, grain boundaries exhibit a positively charged core, resulting in depletion of oxygen vacancies in the space-charge region, and therefore in a highly resistive grain boundary. The latter hinders the transport of oxygen ions. To investigate the influence of the internal interfaces on the total resistivity, EIS is the most prominent and commonly used tool,^65–70 as this method allows distinguishing to some extent between contributions of the grain and the grain boundary due to differences in the corresponding dielectric relaxation times.^66,68,71

In microcrystalline oxides, the total impedance is typically described using the brick-layer model (BLM),^71–73 where a transport path through the bulk, across and along the grain boundaries is considered. In oxides with highly resistive grain boundaries, such as yttria-stabilized zirconia (YSZ), the transport path parallel to the grain boundary can be neglected.^67,74

The resulting impedance spectra exhibit two semicircles, as shown in Fig. 6, and can be described with an equivalent circuit consisting of two resistance–capacitance elements connected in series, representing the transport along the grain and across the grain boundary. Analysis then allows determining the grain boundary thickness δ_gb:^75,76

where δ_g denotes the grain size and C_gb and C_bulk are the capacitance values from EIS for the grain boundary and bulk, respectively. Furthermore, the space-charge potential at the interface can be calculated according to:⁷⁴A detailed analysis of the impedance of microcrystalline ZrO₂ (Mott–Schottky case) using the BLM confirms that the blocking behavior of grain boundaries is caused by the formation of a space-charge region^74,79,80 and not solely by segregation effects, as previously assumed.^67,81,82 The influence of the space-charge region on the overall impedance becomes even more important when the grain size is comparable to the width of the space-charge region,^83,84 which is typically a few nanometers for oxide ceramics. Then, interface effects are more pronounced and may even dominate the impedance.^57,72 As shown in Fig. 6, the decrease in grain size results in merging of the semicircles representing the bulk and grain boundary transport processes,^56,77,85 making their separation difficult. Furthermore, the accumulated charge carriers may also contribute to the transport through the sample. This is typically observed for the Gouy–Chapman case,⁵⁹ but can also occur for the Mott–Schottky case at low doping concentration.⁸⁶ For large effects, the electrons result in an additional contribution to the conductivity along the grain boundaries:^26,59where n₀ and n_∞ denote the electron concentration at the interface and in the bulk, respectively. An overview of the conductivity/resistivity contributions for different space-charge situations due to accumulation/depletion of charge carriers at the interface is given in ref. 59 and 61.

Fig. 6 (a) Impedance spectra for single-crystal and microcrystalline YSZ. The grain boundaries of the microcrystalline sample lead to the appearance of an additional semicircle. (b) Reducing the grain size to the nanometer level results in merging of semicircles representing the bulk and grain boundary transport processes. Adapted from ref. 77 with permission.

A prominent example, where transport along the grain boundary becomes important, is nanostructured CeO₂.^59,85,87 Here, the space-charge region results in accumulation of electrons at the grain boundary (Fig. 7). While in microcrystalline samples the transport is mainly determined by oxygen ions, nanostructured CeO₂ reveals a dominant electronic conductivity.⁵⁹ As shown in Fig. 8, Tschöpe et al. were capable of describing the change from ionic to electronic transport with decreasing grain size using the space-charge model.^88,89 Furthermore, their numerical calculations revealed that, due to the space-charge layer and segregation effects at the interface, not only the apparent activation energy of the transport process depends on the grain size (Fig. 8) but also the characteristic exponent of the p(O₂)-dependence of conductivity.^88–90 A decrease of the bulk ionic conductivity with decreasing grain size is found for nanocrystalline YSZ due to depletion of oxygen vacancies in the grain, caused by the increasing impact of the space-charge layer.⁷⁹

Fig. 7 Defect concentration for different grain sizes (schematically shown for CeO₂). (a) The space-charge region affects the defect concentration near the interface in microcrystalline material. (b) For grain sizes smaller than four times the Debye length, the space-charge region starts to overlap and (c) finally completely dominates the defect concentration of the whole grain. Adapted from ref. 78 with permission.

Fig. 8 (a) Change in electronic/ionic partial conductivity with the grain size; experimental data taken from ref. 87, 89 and 92–94. (b) Change in apparent activation energy of ionic/electronic transport with the grain size. Adapted from ref. 90 with permission.

However, the analytical description of transport properties using the space-charge model is only valid as long as the grain size is four times larger than the Debye length, δ_g > 4λ_D.^55,75,90 If the grain size is further reduced, the space-charge regions of opposite grain boundaries start to overlap (Fig. 7) and the concentration profiles according to eqn (4) are not valid anymore. In CeO₂ and ZrO₂, this condition is typically fulfilled for grain sizes below approx. 20 nm. If the grain size is further reduced to δ_g ≪ 4λ_D, the space charges completely extend throughout the grain, resulting in an almost constant carrier concentration.⁹¹

For small space-charge potential, Maier derived a so-called nano-factor g to describe the resulting enhancement of conductivity:⁵⁵

Here, c*(δ_g) is the concentration of accumulated charge carriers in the center of the grain, which is given by elliptical integrals depending on the grain size δ_g, the bulk concentration c_bulk and the concentration at the interface c₀.

The influence of the space-charge region on the electrical properties is well reported for nanocrystalline oxides with a high grain-boundary density. In case of mesoporous oxides, not only the grain boundaries act as an interface but also the free surface. Although the same description holds as for grain boundaries,⁵⁵ the surface potential, and therefore the extension of the space-charge layer in the crystallites, is more sensitive to variations in the surrounding atmosphere (gas or liquid penetrating into the pores), thereby allowing to tailor the electrical properties. However, because of their unique architecture, mesoporous thin films may show significant differences compared to nanocrystalline materials.

Hartmann et al. studied the electrical conductivity of mesoporous ceria–zirconia (Ce_xZr_1−xO_2−δ) solid solutions of different composition under varying oxygen partial pressure by means of EIS.⁹⁴ Because of the small crystallite size in the range between 4 and 12 nm, only one semicircle was observed experimentally, as discussed above. Consequently, only the total impedance could be determined. Fig. 9 shows the calculated conductivity values for the different materials. Especially for the pure CeO₂ sample, high conductivity was found. This result can be attributed to the electronic conductivity in the space-charge layer both at the surface and at the grain boundaries, as the obtained value is in good agreement with that expected from the model of Tschöpe et al. (Fig. 8). However, it is worth mentioning that the conductivity was calculated assuming dense films. This means the actual conductivity in the pore walls is expected to be even higher.

Fig. 9 (a) Total conductivity of mesoporous Ce_xZr_1−xO_2−δ solid solutions. Adapted from ref. 94 with permission. (b) Comparison of the oxygen-partial-pressure-dependence of conductivity for mesoporous Ce_0.8Zr_0.2O_2−δ thin films and for CeO₂ and Ce_0.875Zr_0.15O_2−δ single crystals at 500 °C. (c) Oxygen-partial-pressure-dependence of conductivity for mesoporous Ce_xZr_1−xO_2−δ solid solution thin films at 300 °C. Adapted from ref. 94 with permission.

Fig. 9 also shows a comparison of the p(O₂)-dependence of electrical conductivity at 500 °C for mesoporous Ce_0.8Zr_0.2O₂ thin film and single crystals of comparable composition. The latter serve as a kind of reference without the influence of interfaces. Similar to the pure CeO₂, mesoporous Ce_0.8Zr_0.2O₂ showed a significant increase in total conductivity (increased electronic conductivity). The p(O₂)-dependence of conductivity further supports the assumption of dominant electronic transport, as the slope of −1/4 at high oxygen partial pressure is characteristic of electronic transport according to the standard defect chemical description (if the oxygen vacancy concentration can be assumed constant [extrinsic regime]).^85,95,96 For low oxygen partial pressure, standard defect chemistry predicts a −1/6-dependence (intrinsic regime) when the concentration of electrons becomes comparable to the oxygen-vacancy concentration.^85,95–97 The mesoporous thin film showed a maximum in conductivity and a decrease with further decreasing oxygen partial pressure, which can be attributed to a hopping maximum of the electronic transport in the Ce³⁺ sublattice. Because the electronic conductivity in CeO₂ is determined by electron hopping from Ce³⁺ to Ce⁴⁺ sites, the hopping probability depends on the fraction of Ce³⁺ and Ce⁴⁺ ions in the lattice according to:^98,99

with a being the distance between the hopping sites, ν₀ the attempt frequency, n_Ce the number of cerium lattice sites and ΔH_m the migration enthalpy of the hopping process. The product x_Ce³⁺(1 − x_Ce³⁺) takes into account that hopping can only occur from Ce³⁺ to Ce⁴⁺. Hence, at high electron concentration, x_Ce³⁺ ≳ 0.5, the hopping probability decreases again due to the reduced number of available Ce⁴⁺ species, leading to a decrease in electronic conductivity, as recently reported for Pr-doped CeO₂ single crystals, where the electrons hop in the Pr-sublattice.⁹⁹ The increase in electronic conductivity and the occurrence of a hopping maximum suggest high electron concentration, and therefore significantly higher reducibility of the mesoporous material compared to nanocrystalline (dense) material (due to the higher specific surface area).

However, more interesting is the p(O₂)-dependence of conductivity at lower temperatures. As shown in Fig. 9, the mesoporous Ce_0.8Zr_0.2O₂ thin film revealed a much stronger p(O₂)-dependence, with a characteristic slope of −1/2 at 300 °C, while for both mesoporous Ce_0.6Zr_0.4O₂ and Ce_0.4Zr_0.6O₂ thin films, a p(O₂)-dependence with an exponent of approx. −1/3 was observed. The unusual p(O₂)-dependence cannot be explained using the standard defect chemical model. Hartmann et al. proposed that oxygen vacancies with a relative charge lower than +2 are responsible for the behavior. Specifically, two scenarios have been discussed. First, formation of defect associates (due to high electron concentration) may either result in not completely ionized vacancies, , or defect associates formed by vacancies and Ce³⁺ ions.^101,102 Second, O⁻ instead of O²⁻ ions were suggested to exist at the surface to compensate the Ce³⁺ ions. While the origin of the unusual p(O₂)-dependence still remains largely unclear, the results show that the high surface area of mesoporous materials can significantly alter their electrical properties.

3.3. Protonic conductivity

In recent years, especially nanostructured materials are receiving much attention, as significant protonic conductivity may be observed under wet atmospheric conditions for oxides, such as YSZ,^103–106 doped CeO₂,^107–110 Al₂O₃,¹¹¹ SiO₂¹¹² or TiO₂.^113–115 Note that their dense counterparts show no signs of protonic conductivity. The increase in conductivity under wet atmospheric conditions is shown in Fig. 10 for various materials. In contrast to high-temperature proton conductors with a perovskite structure,^116,117 like Y-doped BaZrO₃, where the protons are incorporated into the bulk of the material, the protonic conductivity in nanostructured oxide ceramics at low temperatures relies on the adsorption of water molecules at the surface and grain boundaries (interfaces). Because it is directly related to the nanostructure, this offers an additional degree of freedom to tailor the protonic conductivity of mesoporous oxides, rendering them interesting as water sensors,^118–121 as proton exchange membranes^122–124 or for applications in energy storage and conversion.^106,125

Fig. 10 Arrhenius plot of the temperature-dependent conductivity for different nanostructured oxides. Under wet atmospheric conditions, the adsorption of water molecules at the interfaces leads to significant increases due to surface protonic conductivity. Adapted from ref. 100 with permission.

Although the high interface density is responsible for the increased protonic conductivity, at first glance, it is not obvious if the transport of protons occurs along grain boundaries or the (inner) surface. First studies on nanocrystalline YSZ revealed an increase in protonic conductivity with decreasing grain size,^105,114 suggesting that the transport does occur along the grain boundaries¹²⁶ (or at least is related to the grain-boundary density). However, it is now widely accepted that the increased conductivity arises from protons moving in chemisorbed and physisorbed water layers along the pores and surface of nanostructured oxides.^{100,104,127–130} The strong dependence of the protonic conductivity on relative humidity and temperature is related to the structure of the adsorbed water molecules, as summarized by Stub et al.¹³¹ The structure of the water layer on TiO₂ is schematically shown in Fig. 11.

Fig. 11 Schematic illustration of the structure of adsorbed water molecules on an oxide (TiO₂) surface. The influence of an acceptor A′ or a donor D˙ on the water dissociation is also indicated. Reprinted from ref. 132 with permission.

At high temperature and/or low relative humidity, an “ice-like” layer is formed on the surface, consisting of chemisorbed and physisorbed molecules. In the former case, the oxide surface chemistry determines the amount of adsorbed water molecules and the type of protonic charge carriers. For instance, for an ideal ZrO₂ surface, hydroxyl groups are formed in the chemisorbed layer, which apparently are stable up to 600 °C.^{100,131,133,134} Hydroxyl groups can also be formed due to the presence of surface defects, such as oxygen vacancies, where the water molecule fills the vacancy according to (Kröger–Vink notation):^110,117,135

At higher relative humidity, the number of physisorbed layers increases, resulting in transition from “ice-like” to “liquid-like” water layer on the surface. The increase in water layer thickness is also accompanied by a change in transport mechanism. While in the “ice-like” layer, the protonic charge carriers move via Grotthuss hopping,¹³¹ vehicle-type transport of hydronium ions is found in the physisorbed layers due to the weaker bonding between the water molecules.^100,131,133

Stub et al. probed the protonic conductivity in porous TiO₂ with varying doping levels.¹³² High-valent dopants, which are more acidic, are likely to attract electrons, thereby increasing the dissociation of water. This in turn increases the protonic charge-carrier concentration, and consequently also their mobility. Overall, donor-dopants improve the protonic conductivity, while the opposite effect is expected for acceptor-doped oxides. Hence, higher surface acidity results in increased protonic conductivity, as also reported for Cl-doped, mesoporous Al₂O₃, for example.¹¹¹ The profound effect of pore wall chemistry on the protonic conductivity has also been observed by Vichi et al., who investigated the protonic conductivity of mesoporous TiO₂ pretreated in solutions of different pH and by modifying the surface with phosphate anions.¹³⁷ Similar results have been reported by Marschall et al., who studied the influence of sulfonic acid functionalization on the protonic conductivity in mesoporous silica.¹³⁸

In 2013, Gregori et al. examined the proton transport in porous CeO₂. They discussed about the influence of space-charge region on the proton transport at oxide surfaces.¹⁰⁷ Here, the incorporation of water molecules into the space-charge region according to eqn (12) would result in an increase of hydroxide ions in the water layer at the surface, which may be responsible for the enhanced protonic conductivity at low temperature, too.

The importance of space-charge effects has also been addressed by Stub et al. in 2018.¹³⁶ EIS measurements conducted on nanocrystalline YSZ revealed two transport processes, one representing the proton transport along the grains (intra-grain transport), the other was attributed to a resistive transport across the grain boundaries. The positive core of the grain boundary (space-charge region of positive electric potential) results in accumulation of hydroxide ions, as schematically shown in Fig. 12, hindering the movement of hydronium ions in the physisorbed water layer.

Fig. 12 Schematic illustration of the resistive behavior of a grain boundary with adsorbed water molecules on the surface of an oxide ceramic (MO). Because of the positive core, hydroxide ions accumulate at the grain boundary, impeding the transport of hydronium ions in the physisorbed water layer. Reprinted from ref. 136 with permission.

In porous materials, the adsorption of water not only depends on the surface chemistry but also is strongly determined by the structure. Capillary condensation into the nanoscale pores facilitates water uptake.^{107,124,127,130} As discussed by Vichi et al. for mesoporous TiO₂, the adsorption of water at the interface occurs in two steps.¹³⁷ First, water molecules are adsorbed depending on the surface chemistry of the material, while in the second step the pores are filled due to capillary condensation, independent of the surface properties. After filling of the pores with water, the protonic conductivity in mesoporous thin films can be maintained, even when changing the relative humidity.¹³⁹

However, complete pore filling can also negatively affect the protonic conductivity. Vichi et al. observed an increase in activation energy of the proton transport in completely filled pores, which they attributed to bridges formed between water molecules of opposite walls.¹³⁷ Comparable results were also obtained for mesoporous SiO₂^140–142 and Al₂O₃;¹¹¹ the highest protonic conductivities were achieved with 4 and 12 nm pores, respectively. For smaller pore size, the movement of protonic carriers is hindered by geometrical restrictions.

For cubic mesoporous YSZ thin films with a pore size of about 24 nm, even a gradual decrease in total conductivity with increasing relative humidity was observed above 200 °C, as shown in Fig. 13.¹³⁵ However, porous YSZ thin films may show a significant increase in the same temperature range (Fig. 13).¹²⁹

Fig. 13 Arrhenius plot of the conductivity for (a) mesoporous and (b) porous YSZ thin films under wet atmospheric conditions. Adapted from ref. 129 and 135 with permission.

The decrease in total conductivity in case of the mesoporous films was attributed to the annihilation of oxygen vacancies at the surface according to eqn (12), which are the majority charge carriers in YSZ. A decrease in total conductivity under wet atmospheric conditions has also been reported for microcrystalline CeO₂ by Chueh et al. They attributed this result to changes in the space-charge potential of grain boundaries upon hydration.¹⁴³ Although the influence of pore size on the protonic conductivity in oxide materials is not yet fully understood, the results clearly demonstrate the considerable effect that the pore structure has in general. In contrast to nanocrystalline oxides with an ill-defined porosity, the regular pore network of ordered mesostructured materials seems beneficial to device applications, as the pore size and structure, as well as the wall thickness, can be tailored to some degree to optimize water uptake (adsorption), and therefore the protonic conductivity.

4. Electrochemical applications

As discussed above, mesoporous oxides are characterized by a high specific surface area combined with a unique pore structure and nanocrystalline walls. These properties make them interesting for a broad range of electrochemical applications,^14,144–146 for example, as active electrode material in batteries, for (photo-)electrocatalysis and oxygen storage or as gas sensor. It should be noted that the favorable properties are not only a result of the increased surface area. Instead, the major advantage is the presence of a uniform and continuous pore network, allowing effective transport of electronic and ionic charge carriers and providing good accessibility of the material and its internal structure to the surrounding medium, as gasses or liquids can readily penetrate into the porosity. In the following, we discuss in some more detail the pros and cons of mesoporous metal oxides and which properties are beneficial to the functionality and performance of electrochemical devices.

4.1. Mesoporous metal oxides for batteries

The working principle of an insertion- or intercalation-type material for secondary battery applications relies on both the reversible storage of ions (Li⁺, Na⁺etc.) during electrochemical cycling and the efficient transport of charge-compensating electrons (Fig. 14). In recent years, mesoporous electrode materials have been shown to possess promising properties, allowing for superior energy storage and high power density, among others.¹⁴⁷ The mesoporous morphology ensures equal distribution of electrolyte in the (bulk) structure, minimizing the fraction of inactive material and allowing fast ion diffusion along the pore system.^147–149 Furthermore, the interconnection of nanoscale particles in the pore walls provides short diffusion lengths and efficient transport pathways for both the ions and electrons.

Fig. 14 Schematic illustration of the charge-transport processes accompanying lithium insertion/extraction into/from a mesoporous material, with the nanoscale porosity allowing for short diffusion path lengths.

The positive effect of mesoporosity on the charge-storage properties has been proven for a variety of cathode and anode active materials, mainly for lithium-ion battery (LIB) applications, such as LiCoO₂,¹⁵⁰ LiFePO₄,¹⁵¹ β-MnO₂,¹⁵² LiMn₂O₄,^152,153 Li₄Ti₅O₁₂,³⁸ TiO₂,^19,154–156 α-Fe₂O₃¹⁵⁷ or NiO.¹⁵⁸ For example, Jiao et al. compared the electrochemical performance of nanowire and mesoporous LiCoO₂ to that of “normal” LiCoO₂. The mesoporous material delivered the largest specific discharge capacity and showed the best capacity retention, which was attributed to the unique pore–solid architecture with good connection among adjacent crystallites.¹⁵⁰ Brezesinski et al. came to the same conclusion from studying polymer-templated and nontemplated (nanocrystalline) TiO₂ thin films.¹⁹ These results confirm the beneficial effect of nanoscale porosity on the electrochemical performance.

An important parameter when characterizing battery materials is the amount of stored charge. In general, the total amount consists of three contributions, as summarized in Fig. 15, namely, the faradaic contribution from the insertion of ions into the bulk of the material, the faradaic contribution from charge-transfer processes at the solid/liquid interface (pseudocapacitive charge storage) and the nonfaradaic contribution due to formation of a Helmholtz double layer.^{19,156,159–161} Analysis of cyclic voltammetry (CV) data obtained at different sweep rates v allows distinguishing between diffusion- and nondiffusion-controlled charge storage. Because ion insertion into the bulk is a semi-infinite diffusion process, the current is proportional to v^1/2. In contrast, for capacitive processes, the current depends linearly on the sweep rate. Hence, current as a function of sweep rate is given by:¹⁶²

j(v) = k₁v + k₂v^1/2, (13)

where k₁ depends on the surface area A and the capacitance C of the electrode, while k₂ is given by the Randles–Ševčík equation:¹⁶³with F being the Faraday constant, D the diffusion coefficient and n the number of electrons transferred in the redox reaction. Rearranging eqn (13) then allows separating between both contributions:Especially for nanocrystalline materials, capacitive charge storage becomes significant and plays an important role in the performance and for the kinetics.^{19,156,159,164–166} Analysis of CV data for mesoporous TiO₂ thin films showed that the capacitive contribution to charge storage is virtually independent of the sweep rate, while the bulk insertion significantly decreases with increasing sweep rate.¹⁹ For mesoporous MoS₂, Cook et al. found that 80% of the theoretical specific capacity is accessible within 20 s due to pseudocapacitive charge storage.¹⁶⁷ These results indicate that fast kinetics in mesoporous materials also originates from nondiffusion-controlled surface charge-storage processes.

Fig. 15 Charge-storage mechanisms in nanocrystalline metal oxides. Adapted from ref. 156 with permission.

Significant pseudocapacitive charge storage has also been observed for mesoporous group V transition-metal oxides, including Nb₂O₅ and Ta₂O₅, and for MoO₃.^166,168 Direct comparison of crystalline and amorphous thin films showed that the pseudocapacitance achieved with the crystalline samples exceeds that of the amorphous films by large. The increased pseudocapacitive contribution to charge storage was attributed to the lithium insertion into the interlayer gaps of the layered crystalline materials (referred to as intercalation pseudocapacitance).

An aspect that is rarely being considered is that the amount of charge stored at the interface can be much larger than that stored in the bulk due to formation of the space-charge region.^169–172 This effect may even allow to store, e.g., lithium at the interface between two materials, for which such behavior is not possible in the individual phases.¹⁷³ For example, Yue et al. investigated the pseudocapacitive charge storage in mesoporous Nb-doped TiO₂.¹⁵⁶ Because TiO₂ has a low electronic conductivity, doping with Nb, acting as a donor, can significantly improve the conductivity of the material.^174,175 Apart from the increase in electronic conductivity, the authors found that the pseudocapacitance is also increased up to doping levels of about 5 at%. Higher concentrations led to a decrease in pseudocapacitive charge storage. This behavior was explained by changes in the space-charge layer at the interface, as shown in Fig. 16. For TiO₂ in contact with a Li⁺-containing electrolyte, the space-charge region results in depletion of electrons at the interface and accumulation of lithium ions due to differences in electrical potential between the surface (interface) and the bulk.¹⁷⁰ For low doping concentration, the space-charge region of opposite sites overlaps, resulting in negligible potential differences,^176,177 and therefore in low surface charge storage. With increasing doping level, the space-charge region decreases. Hence, the potential differences increase and more ions can be stored at the interface, leading to increased pseudocapacitance. However, the width of the space-charge region continues to decrease with further increasing doping concentration. For that reason, the pseudocapacitive contribution to charge storage starts to decrease again at some point.

Fig. 16 (a) Electrical conductivity and pseudocapacitive charge storage of mesoporous TiO₂versus the Nb doping level. (b) Schematic illustration of the surface band bending and the corresponding defect concentration for different doping situations. Adapted from ref. 156 with permission.

Although the extension of the space-charge region in oxides is typically less than a few nanometers, this example shows that, in mesoporous materials with crystallites of similar dimensions, the charge storage in the space-charge region can become significant and even dominate the overall behavior.

Another major advantage of mesoporous materials is their mechanical flexibility. The insertion/extraction of ions into/from the electrode during electrochemical cycling is usually accompanied by volume changes,^178–181 often resulting in mechanical degradation (fracturing or even pulverization of the active material).^182–188 In mesoporous materials, such adverse effects can be mitigated, as mechanical strain is accommodated to some degree by pore flexing.^144,189,190 Liu et al. investigated the electrochemical properties and cycling stability of ordered mesoporous NiO produced by hard templating.¹⁵⁸ While NiO as anode in LIBs typically suffers from pulverization and accelerated capacity fading due to volume changes during the conversion reaction, among others,^191,192 the mesoporous material maintained a specific capacity of 680 mA h g⁻¹ at 0.1C rate after 50 cycles. Lee et al. studied ordered mesoporous niobium nitride/N-doped carbon composite as an advanced anode material for application in K-ion batteries.¹⁹³ Structural characterization using in situ X-ray diffraction (XRD) showed that the high cycling stability (stable performance for >2000 cycles) is due, in part, to the negligible mechanical strain that the material experiences during cycling.

The mechanical flexibility also improved the lithium storage in mesoporous CeO₂. Specifically, Brezesinski et al. probed the charge-storage behavior of polymer-templated and nontemplated thin films (Fig. 17). As expected, the authors observed a significant increase in pseudocapacitive charge storage in the mesoporous CeO₂.¹⁶⁵ However, for slow sweep rates or long charging times, that is, if the lithium ions have enough time to diffuse into the lattice, the total amount of stored charge was also increased by a factor of about two compared to the nontemplated material. Using small-angle X-ray scattering (SAXS), they were able to show that the mesoporous CeO₂ thin films have tensile stress imposed by the substrate (note that templated sol–gel films usually undergo unidirectional volume contraction during heating). Upon ion insertion into the lattice, the porous structure expands normal to the substrate, relieving in-plane stress, which in turn facilitates charge storage.

Fig. 17 (a) Charge-storage dependence on sweep rate for nanocrystalline (nontemplated) and mesoporous (KLE-templated) CeO₂ thin films. (b) CV curves for mesoporous CeO₂. The capacitive current is highlighted in gray. (c) Galvanostatic discharge curves at C/2 rate. (d) SAXS patterns for mesoporous CeO₂ and Li_xCeO₂. Adapted from ref. 165 with permission.

The structural stability of ordered mesoporous NiCo₂O₄ during electrochemical cycling has been examined by Bhaway et al.¹⁹⁴ Using grazing-incidence SAXS (GISAXS) and XRD (GIXD), they were able to correlate structural changes of the anode material with the electrochemical performance. Electrodes with approx. 9 nm pores showed severe capacity fading caused by collapse of the nanostructure during the initial cycles. In contrast, for electrodes with larger pores, the nanostructure was retained, which agrees well with the improved cyclability.

Increased structural stability has also been observed by Krins et al. for mesoporous amorphous NbVO₅ thin films with a wormlike pore structure.¹⁹⁵ The authors found that their ability to accommodate the strain from lithium insertion increases with decreasing pore size from 100 to 15 nm, suggesting that the optimal size is closer to 10 nm.

Mesoporous nanocrystalline architectures offer an additional advantage in terms of ion uptake, as the particle size may have a strong impact on the thermodynamics.¹⁹⁶ For example, lithium insertion into bulk TiO₂ is accompanied by a phase transition from anatase TiO₂ to Li_0.55TiO₂.^197–200 With further insertion of lithium, a second phase transition occurs, resulting in the formation of a blocking LiTiO₂ layer at the particle surface. LiTiO₂ exhibits a low lithium-diffusion coefficient, and therefore prevents complete phase transformation (kinetic limitation).¹⁹⁸ However, as discussed by Wagemaker et al., increasing contribution of surface free energy to the total free enthalpy of the system alters the solubility limit of lithium in TiO₂ and can even suppress phase separation.^198,199 For particles ≤7 nm in diameter, both phase transitions are suppressed, allowing full lithiation to form LiTiO₂,¹⁹⁸ which also helps explain the larger specific capacity of nanostructured TiO₂ compared with bulk material.^200–202

Similar behavior has been observed for α-Fe₂O₃ (hematite). While larger particles undergo a phase transition from trigonal α-Fe₂O₃ to cubic Li_xFe₂O₃ at x ≈ 0.05, in nanoparticles, this transition is suppressed and lithium can be inserted into the lattice up to x ≈ 1. As the lithium uptake is reversible for α-Li_xFe₂O₃, nanoparticles exhibit much improved cyclability.^203,204 Note that increased lithium storage and good cycling stability have also been found for mesoporous hematite and other Fe-based oxide thin films.^205,206

4.2. Mesoporous metal oxides in catalysis

Because of their unique structural properties, mesoporous oxides represent a promising class of materials for a variety of catalytic applications, for example, as catalyst or support for CO and NO conversion reactions or for (photo-)catalytic hydrogen production and water splitting. While their high surface area provides a large number of catalytically active sites, the 3-dimensional pore structure ensures good accessibility of the bulk and enables efficient transport of reaction species to/from the material. However, the catalytic properties not only depend on the specific surface area but also the wall structure (grain size, crystallinity etc.).

Ceria–zirconia solid solutions are among the most prominent catalytic materials for CO and NO conversion due to the high redox activity of the Ce³⁺/Ce⁴⁺ couple.²⁰⁷ Furthermore, because of its smaller ionic radius, the incorporation of Zr⁴⁺ ions into the lattice distorts the cubic fluorite structure, resulting in the formation of defects and higher reducibility.^97,208 Overall, ceria–zirconia solid solutions exhibit favorable catalytic properties as a result of their unique oxygen-storage capacity, despite the fact that Zr⁴⁺ is redox inactive.^207,209

There are several reports on the investigation into catalytic properties of mesoporous Ce_xZr_1−xO_2−δ.^210–214 For example, Petkovich et al. probed the catalytic activity of ordered macroporous versions of such materials, prepared using two different synthetic routes, for H₂O splitting and H₂ production.²¹⁵ They found increased H₂ production rates compared to commercial CeO₂ powder, which was attributed to the increased specific surface area. However, the reaction rate was not directly related to the surface area, but instead to the composition and homogeneity of the crystallites. The highest rate was observed for materials with a Zr content between 10 and 20%. Furthermore, samples with compositional heterogeneity, that is, the presence of Ce- and Zr-rich phases, offered better performance than the respective homogeneous (single-phase) materials. Mamontov et al. also observed a correlation between heterogeneity and oxygen-storage capacity in nanocrystalline ceria–zirconia powders.²¹⁶

Ho et al. reported on mesoporous Pd/ceria–zirconia solid solutions with macrochannels,²¹⁷ showing superior activity for CO oxidation over reference samples without macrochannels. The improvement in performance was attributed to two main factors: first, alignment of macrochannels, enabling effective mass transport through the pore system, and second, increased Ce³⁺ and oxygen-vacancy concentrations at the surface (preferred reaction sites).²¹⁸ In contrast to the report by Petkovich et al., crystal structure homogeneity was found to be advantageous, in agreement with other reports,^209,219 as oxygen vacancies are more readily formed in the cubic phase than in mixed-phase structures.²²⁰ Cui et al. examined the influence of La and Pr doping on the catalytic properties of ceria–zirconia solid solutions. They came to the same conclusion that larger pores allow for more effective gas diffusion. They also observed that an increased fraction of Ce³⁺ ions due to doping is beneficial to the catalytic activity and the oxygen-storage capacity.²²¹ Similar results were obtained for mesoporous CeO₂-based catalysts used in methanol decomposition.^210,222

The strong correlation between oxygen-storage capacity and catalytic activity has been theoretically studied by Sayle et al. for nanoscale CeO₂.²²³ Molecular dynamics (MD) simulations confirmed that the catalytic activity is strongly affected by its oxygen content, and therefore directly related to the oxygen-storage capacity. Furthermore, they found that the catalytic activity depends on the morphology of the material, such as wall diameter and particle size. In general, oxygen extraction from nanoparticles is more facile. This is in agreement with density-functional theory (DFT) calculations by Migani et al., predicting the least oxygen-vacancy formation energy for nanoscale CeO₂.^224,225

In 2018, Hao et al. confirmed the considerable effect that the particle size has on the oxygen-storage capacity of CeO₂.²²⁶ The oxygen-storage capacity of 5 nm particles (approx. 260 μmol g⁻¹ at 250 °C) was shown to be two times larger than that of 12 nm particles (Fig. 18). Using scanning tunneling microscopy (STEM) combined with electron energy loss spectroscopy (EELS), they observed an increase in Ce³⁺ concentration at the surface of the nanoparticles or, in other words, the presence of a space-charge layer. The space-charge layer extends over the whole nanoparticles when their size is below 6 nm, resulting in a total Ce³⁺ content of more than 50%. The increase of Ce³⁺ at the surface is accompanied by an increase of the lattice parameter, leading to a lower formation energy for oxygen vacancies. This helps explain the increase in oxygen-storage capacity with decreasing particle size.

Fig. 18 (a) Comparison of the oxygen-storage capacity of CeO₂ nanoparticles of different size and commercially available nanopowder. (b) Ce³⁺ content as function of the particle size and (c) valence-state distribution. Adapted from ref. 226 with permission.

Apart from ceria–zirconia solid solutions, other oxides have also been considered for catalytic applications. Ren et al. studied the CO conversion on ordered mesoporous metal oxides, such as Co₃O₄, Cr₂O₃ and Fe₂O₃, to name a few.²²⁷ These materials exhibited increased CO conversion rates compared to their bulk counterparts, as shown in Fig. 19. The improvements were attributed not only to the increased specific surface area but also to the presence of a regular 3-dimensional pore structure, providing well-defined pathways for gas transport and equally accessible internal surfaces. Similar results were obtained by both Sun et al. and Tüysüz et al. for CO oxidation over ordered mesoporous Co₃O₄.^228,229

Fig. 19 CO conversion on (a) mesoporous and (b) bulk oxide catalysts. Adapted from ref. 227 with permission.

Furthermore, noble-metal nanoparticles as co-catalysts have been shown capable of strongly affecting the activity of oxides by facilitating the conversion reaction (for oxide-supported transition-metal catalysts). The reason is that metallic nanoclusters on the surface increase the oxygen-vacancy concentration and improve the charge transfer.^232,233 Both results in an enhancement of catalytic activity at the metal/oxide interface.^234,235 The beneficial effect of using metallic nanoclusters has been shown for a variety of mesoporous oxides.^210,217,236 However, the reaction rate also depends on the size (smaller particles are usually more active), shape and distribution of nanoparticles inside the mesoporous framework, as demonstrated, for example, by Kapoor et al.²¹⁰ Because the mesoporous support can be post-loaded,²³⁶ this allows for pre-synthesis of tailored nanoparticles. Furthermore, the interaction with the oxide surface prevents the agglomeration of particles.

In addition to their application in conversion reactions, mesoporous metal oxides also gained great interest in the field of photo(electro-)catalysis for water splitting (hydrogen/oxygen production) or carbon dioxide reduction. The basic principle is based on the generation of an electron–hole pair in a semiconducting material under light irradiation, as shown in Fig. 20 for the water-splitting reaction using a particulate photocatalyst.^230,231,237 Both the electron and electron hole then undergo secondary reactions at the surface, which may also be supported by a metal co-catalyst. At the surface of the catalyst, electrons reduce H₂O molecules to form H₂, while H₂O molecules are oxidized by holes to produce O₂. Hence, for photocatalysis, semiconductors with a sufficiently large band gap are needed. Specifically, the valence band should be more positive than the potential for the O₂/H₂O redox reaction, while the conduction band has to be more negative than the potential of H⁺/H₂.

Fig. 20 (a) Schematic of the electron–hole pair generation in semiconducting nanoparticles and possible recombination mechanisms: (A) surface and (B) volume recombination and (C) reduction and (D) oxidation reactions. Adapted from ref. 230 with permission. (b) Schematic illustration of the setup for photoelectrocatalytic water splitting using TiO₂ as photoanode. (c) Energetic position of redox states relative to the bands of the semiconductor. Adapted from ref. 231 with permission.

In photoelectrocatalysis, both reactions are separated by applying an electrical bias (Fig. 20). The catalyst is used as photoelectrode, where only one of the redox reactions occurs. An important aspect for efficient photo(electro-)catalysis is the separation of electron and hole to ensure that both carriers can diffuse to the active surface sites for the redox reactions to occur. High (crystal) quality of the absorber material is advantageous, as defects act as carrier traps, increasing the probability of recombination. Overall, there are many reports on mesoporous semiconducting oxides for photocatalytic applications, including TiO₂,^20,238–240 Ta₂O₅,²⁴¹ LiFe₅O₈,²⁴² ZnFe₂O₄,^243–245 and Rh₂O₃.²⁴⁶

Given that the water-splitting reaction takes place at the surface of the photocatalytic material, it is reasonable to assume that a high specific surface area leads to superior activity. Cherevan et al. analyzed around 200 publications regarding the photocatalytic activity of mesoporous semiconductors, aiming at correlating improved performance with surface area.²⁴⁷ They found that increased specific surface area is not the only reason and may even impede any performance improvements, as a high surface area may also increase the probability of carrier recombination. For example, for Ta₂O₅, an increase of the specific surface area by a factor of approx. 100 (mesoporous vs. nonporous material) only resulted in a relatively small activity rise by a factor of 1.5.²⁴⁸ Instead, the crystallinity, the pore size, the pore connectivity and the wall thickness can significantly affect the performance of the material.

Kirchberg et al. compared the photocurrent density of mesoporous ZnFe₂O₄ photoanodes prepared using two different polymer templates and calcined at various temperatures. In fact, they observed that differences in pore morphology only have a minor influence on the photocurrent. Samples of high crystallinity showed better performance due to the relatively lower concentration of defects acting as recombination centers for the electron–hole pairs.²⁴³ Increased photocatalytic activity because of improved crystallinity has also been reported for mesoporous Nb₂O₅, Mn₂O₃ and NaTaO₃.^249–251

Tüysüz et al. compared mesoporous nanocrystalline NaTaO₃ and amorphous NaTaO_x photoanodes and observed a decrease in photocatalytic activity for the crystalline material.²⁵² Although the crystallinity should reduce the probability of carrier recombination, crystallization of the amorphous matrix resulted in an increase in particle size and a decrease in specific surface area. As discussed by the authors, smaller particles are considered beneficial to the photocatalytic activity. The reason is that the space-charge regions of opposite surfaces overlap, resulting in complete depletion of the particle and negligible band bending at the surface (flat-band condition),^176,177 as mentioned previously. Both photogenerated charge carriers can then easily diffuse to the active surface sites, leading to improved photocatalytic activity, as schematically shown in Fig. 21. For larger particles/crystallites, increasing band bending impedes the diffusion of one kind of charge carrier to the surface, resulting in higher recombination rates in the bulk, and therefore in lower efficiency.

Fig. 21 Effect of particle size on the band bending and photocatalytic activity. (a) Generated carriers show high probability to recombine at boundaries or defects in larger particles. In smaller particles, the carriers can readily diffuse to the catalytically active surface sites. Adapted from ref. 231 with permission. (b) The corresponding band structure. In larger particles, the space-charge region separates the generated carriers, while smaller particles exhibit negligible band bending. Adapted from ref. 176 with permission.

The opposite holds true when the mesoporous material is used for photoelectrocatalytic applications, where the two redox reactions occur at different electrodes. In this case, the band bending in larger crystallites promotes charge separation, thus preventing undesirable recombination. For example, Hartmann et al. investigated the photoelectrocatalytic activity of mesoporous TiO₂ prepared from preformed nanoparticles and a classical sol–gel route.²⁰ They observed higher photocurrents and improved quantum efficiency for the sol–gel-derived thin films with thicker pore walls and larger crystallites. The results were explained, in part, by a reduced probability of electrons to react at the surface (Fig. 22). This can be directly related to the existing band bending, preventing the diffusion of electrons to the surface. Furthermore, the thicker wall structure gave rise to better connectivity between individual crystallites, allowing for more efficient transport of electrons in the pore–solid architecture and contributing to the higher efficiency. The positive effect that a 3-dimensional pore structure has on the photocatalytic performance (due to improved transport of electronic charge carriers) has also been discussed by Hossain et al.²⁵³ and Zhou et al.²⁵⁴

Fig. 22 Incident photon-to-current efficiency for water photoelectrocatalysis over (a) nanoparticle-derived and (b) sol–gel mesoporous TiO₂ thin films. The nanoparticles in the wall structure result in low electronic conductivity and high recombination rate, while the sol–gel films with thicker and continuous pore walls reveal high photocurrents with a relatively lower recombination rate. Adapted from ref. 20 with permission.

As in case of battery applications, the pore connectivity and pore size are critical parameters affecting the photocatalytic activity, as they determine the transport of reactants and products. For example, Weller et al. investigated the influence of pore size of mesoporous CsTaWO₆ (with Rh as co-catalyst) on the photocatalytic performance for H₂ evolution.²⁵⁵ An optimized synthesis route allowed varying the pore size and mesoporous morphology without affecting the crystallinity of the material. Analysis of the photocatalytic activity revealed that the H₂ evolution rate strongly depends on the pore size and the pore size distribution. This result was explained by transport limitations, as the water/methanol mixture used as electrolyte was partly confined in the small pores due to capillary forces, blocking the release of gas bubbles.²⁵⁶ Furthermore, comparison between samples with (large-pore) ordered and disordered mesoporous structures revealed higher H₂ evolution rates for the former material, despite having a lower specific surface area (Fig. 23). The increase in activity was rationalized by better accessibility of the bulk material and improved transport pathways in case of the regular pore structure. Similar results have been reported by Fang et al., for example, who were capable of achieving efficient mass transport by introducing additional macropores.²⁴⁰

Fig. 23 Comparison of steady-state H₂ evolution rates for disordered and ordered mesoporous CsTaWO₆ with different amounts of Rh co-catalyst. Adapted from ref. 256 with permission.

Another interesting aspect is that the regular arrangement of crystallites apparently affects the optical properties. For example, Hossain et al.²⁵³ found that ordered mesoporous anatase TiO₂ with 5 nm thick walls exhibits a smaller band gap than individual 25 nm particles. The decrease in band-gap energy by about 0.2 eV was ascribed to the interconnectivity of nanoparticles, increasing their effective size.²⁵³ Hufnagel et al. also observed a significant improvement in light-absorption properties (light-harvesting efficiency) for nanoscale ZnFe₂O₄ grown on macroporous antimony-doped tin oxide films compared to nonstructured thin films.²⁴⁴ This result was attributed to the presence of extended light paths caused by multiple light reflection and scattering in the porous electrode, as also reported by Fang et al. for hierarchically structured TiO₂ microspheres.²⁴⁰

Taken together, the results indicate that a favorable structure for catalytic applications depends on the interplay between surface area, crystallinity, particle size (often determined by the wall thickness), pore size and pore size distribution. However, the individual parameters can significantly differ among different materials, making it necessary to vary them independently. For mesoporous oxides, unfortunately, this is only feasible to some degree, as e.g., high crystallinity usually requires high-temperature treatment, which in turn results in grain growth and coarsening, and therefore in lowering of the active surface area.

4.3. Mesoporous metal oxides for gas sensors

Most gas sensors rely on the interaction of a specific gas with the surface of the sensing material. There are different kinds of sensing principles, all of which take advantage of changes in the materials properties under varying atmosphere. In capacitive gas sensors, the change of permittivity is measured, while optical gas sensors usually detect changes in photoluminescence, reflection or optical absorption.¹¹⁹

Because oxides are often used as resistive sensors, we only focus on this type in the following. The basic working principle of most resistive gas sensors is schematically shown in Fig. 24. Gas/oxide interactions induce changes of the surface potential, e.g., due to chemisorption of gas molecules (receptor function). The change of the electronic surface structure then results in changes in conductivity, which are detected as the sensing signal (transducer function). In case of an n-type semiconductor, for example, the adsorption of oxygen species (O⁻) typically gives rise to formation of an electron-depleted space-charge layer. In the presence of a gas, such as H₂, O⁻ species are consumed, resulting in a decrease of the space-charge width, and therefore in a decrease of resistance.^258–260 Hence, the working principle of a resistive gas sensor is directly related to the space-charge region at the interfaces (grain boundary or surface) of the oxide material. The largest sensitivity is achieved when the conductivity of the grains is completely determined by the space-charge region, that is, when the grain size is smaller than approx. twice the depletion layer.^261–263

Fig. 24 Schematic illustration of the sensing mechanism in n-type semiconducting metal oxides. (a) The oxide is completely depleted under air exposure due to extension of the space-charge layer. (b) Upon adsorption of reducing gas molecules, the extension of the space-charge layer is reduced, thereby opening conductive channels in the sensor material. Adapted from ref. 257 with permission. (c) Electron-depletion layer in resistive sensors. Granular material with grain size larger than about twice the depletion layer (left), fully depleted granular material (middle) and fully depleted mesoporous material (right). Adapted from ref. 119 with permission.

The large influence of grain size on the resistance and sensitivity of granular SnO₂ sensors is depicted in Fig. 25. Both the sensitivity and response time depend strongly on the morphology and microstructure of the sensing material. To achieve high sensitivity, high surface area is desirable, while efficient and fast diffusion of gas molecules enables good response characteristics of the sensor to changes in the surrounding atmosphere. Consequently, mesoporous materials with tailorable pore and crystallite sizes are of great interest for sensor applications.¹¹⁹

Fig. 25 (a) Electrical resistance at 300 °C in dry air or 800 ppm H₂versus the crystallite size of SnO₂. (b) Sensitivity to 800 ppm H₂ or CO versus the crystallite size of SnO₂. Adapted from ref. 263 with permission.

Similar to granular materials, the crystallite size strongly affects the sensitivity of mesoporous materials. For example, Aqeel et al. probed the NO₂ detection capabilities of mesoporous SnO₂ calcined at 400 and 500 °C.²⁶⁴ They found that the material heated at 400 °C shows significantly better sensing properties, as the crystallite size (4–5 nm) was smaller than twice the depletion layer. The material heated at 500 °C had larger crystallites of approx. 8 nm in diameter, and therefore showed worse sensing performance. Conductivity measurements confirmed the complete depletion of electrons in the smaller crystallites, which exhibited a much larger resistance. In addition, the selectivity for NO₂ was much improved in case of the smaller crystallites. Xu et al. obtained similar findings for SnO₂.^262,263 They noticed an increase in sensitivity and resistance with decreasing grain size (Fig. 25).

Waitz et al. examined the methane sensing properties of mesoporous In₂O₃ as function of the pore size and wall thickness.²⁶⁵ The sensitivity scaled linearly with the specific surface area. The authors also found that especially the wall thickness has a strong effect on sensitivity, confirming that a fully depleted pore wall gives rise to superior performance. However, they pointed out that the structural parameters are coupled to one another, and thus cannot be varied independently, making it difficult to distinguish between the impact of each parameter on sensitivity.

Ghom et al. tested mesoporous ceria–zirconia solid solutions as oxygen sensors at 600 °C.²⁶⁶ In Ce_xZr_1−xO_2−δ, the change in resistance is not caused by changes of the surface space-charge layer due to adsorbed oxygen species, but instead by the significant increase in conductivity because of the release of oxygen under reducing conditions, as described above. As discussed by the authors, the fast response time of such mesoporous sensors is directly related to the oxygen-storage capacity of the material.^209,219 Both the regular arrangement of pores and the pore size not only affect the sensitivity but are also responsible for fast diffusion of gas through the bulk of the material, the latter of which determines the sensor response time.²⁶⁷

Li et al. investigated ordered mesoporous WO₃ for H₂S sensing and observed a fast response of 2 s, with a recovery time of approx. 38 s, which they attributed to the unique (continuous) pore–solid architecture and the resulting fast gas diffusion.²⁶⁹ The similar response time (1–3 s) of mesoporous α-Fe₂O₃ for ethanol sensing was also ascribed to the high diffusivity through the pore network.¹⁵⁷ Li et al. synthesized mesoporous–macroporous SnO₂.²⁶⁸ Compared with conventional granular SnO₂ sensors, the hierarchical pore structure led to significant improvements in response and recovery times, as shown in Fig. 26, which was attributed to facilitated gas diffusion in the macropores.²⁶⁸

Fig. 26 Comparison of recovery and response times to 100 ppm benzene, ethanol or toluene for “traditional” SnO₂ and mesoporous–macroporous SnO₂. Adapted from ref. 268 with permission.

While the sensitivity and response times are determined by structural properties, the selectivity of a sensor depends strongly on the material type and its surface properties. To control or even tailor the selectivity for various gases, the sensing material can be loaded with catalytically active noble-metal nanoparticles, such as Au, Pd or Pt.^260,270,271 As shown in Fig. 27, electrons move from the active oxide to the metal nanoparticles due to different work functions, leading to the formation of a Schottky barrier at the interface, and therefore to an increased electron depletion (electronic sensitization). However, the metal nanoparticles on the surface not only affect the adsorption and dissociation reactions by activating the gas molecules (chemical sensitization)^257,259 but also facilitate oxygen-vacancy formation.^232,233

Fig. 27 Schematic illustration of (a) electronic and (b) chemical sensitization by metal or metal-oxide nanoparticles. Adapted from ref. 259 with permission.

Shimizu et al. studied the gas sensing properties of porous thick films of SnO₂ loaded with 1 wt% Pt or Pd.^260,270 While the sensitivity to H₂ and CH₄ at the surface was improved in the presence of noble metals, it deteriorated in the interior. This result was explained by reduced permeation into the bulk due to increased surface reactivity, emphasizing the necessity of optimizing the gas diffusion pathways.

Ma et al. reported on a positive effect of partially oxidized Pt nanoparticles on the sensitivity and selectivity of mesoporous WO₃ for CO sensing.²⁷² The improvements were attributed to the sensitizing effect of nanoparticles on the WO₃ framework. Using X-ray photoelectron spectroscopy (XPS), they found W⁵⁺ states because of the formation of oxygen vacancies (note that they increase the adsorption of oxygen species at the surface). Furthermore, the good sensitivity was justified by changes in the depletion region at the Pt/WO₃ interface with varying atmosphere. After exposure to air, PtO nanoparticles were detected, leading to the formation of a depletion layer at the interface due to their p-type character. The PtO nanoparticles were reduced back to Pt(0) upon CO exposure, resulting in a vanishing depletion layer and pronounced changes in resistance with varying atmosphere. A positive impact of oxygen vacancies on the NO₂ sensing properties has also been reported for mesoporous WO₃ nanofibers²⁷³ and CeO₂/graphene heterostructures.²⁷⁴

5. Surface modification

Although porous architectures are, without doubt, beneficial to various applications, there are certain constraints in using mesostructured powder and thin film materials. First, they tend to suffer from thermal stability issues. Especially for catalytic applications, elevated temperatures (or high relative humidity) are often required. However, such experimental conditions may result in changes of the pore structure due to phase transitions and/or growth of crystallites (sintering effects) or even complete collapse of the pore network. Second, while some variation in pore size (and wall thickness, pore symmetry etc.) can be achieved, it is mainly determined by the template used. Nevertheless, small changes in pore size can significantly affect the performance of mesoporous materials, as discussed above. Hence, tailoring of structural features on the nanometer level and precise interfacial engineering are desirable.

It has been shown that the properties of mesoporous materials can be altered by coating the surface using atomic layer deposition (ALD). ALD is a variant of the chemical vapor deposition (CVD) technique and is receiving increasing attention in recent years.^275–278 The reason is that ALD is the only method that is suited to produce uniform and conformal coatings on complex surfaces. This is achieved by sequential pulsing of gaseous reactants with intermediate purging steps. The pulses represent self-limiting gas–solid surface reactions, leading to controlled film growth [2-dimensional layer-by-layer (Frank–van der Merwe) growth] and excellent step coverage, unlike other deposition techniques.

The deposition of thin coatings using ALD finds lots of applications, among others, in protection and functionalization of surfaces or stabilization of porous materials.²⁷⁹ For instance, in the battery field, ALD is typically used to modify electrode surfaces.^280–284 The coating acts as an artificial interphase between the electrode and the electrolyte to suppress adverse side reactions, which may result in gas evolution and/or formation of an ion-transport-blocking layer (anode/cathode solid–electrolyte interface). However, ALD also lends itself to the preparation of active electrode materials.^285,286

Apart from electrochemical energy-storage applications, ALD is actively used to improve the catalytic performance of materials by either depositing metal nanoparticles onto their top surface^287,288 or (partial) overcoating them with an oxide layer.^{278,289–293} The growth of single nanoparticles on carbon or oxide surfaces is usually achieved at comparably high deposition temperatures, where the precursor tends to agglomerate depending on the presence of surface defects and the atmospheric conditions instead of forming a conformal coating during the initial ALD cycles.^294–296 In case of partial overcoating, one takes advantage of differences in surface chemistry of the solid (area-selective deposition, also referred to as AS-ALD).^{292,297–300}

In the following, we focus on some selected examples, where ALD has been used to produce a uniform oxide coating in order to modify the structural and/or electrical properties of mesoporous oxides.

Mesoporous oxides are metastable, that is, accelerated grain growth occurs at elevated temperatures, leading to collapse of the pore structure. The deposition of a thin coating using ALD may help to enhance the thermal stability of mesoporous materials. For example, Kraffert et al. synthesized mesoporous ferrihydrite,³⁰¹ showing a phase transition to hematite at 400 °C. By depositing silica onto the pore walls of the mesostructured thin films using a single ALD cycle, the stability of the ferrihydrite phase could be increased to 600 °C. Similar results were achieved for alumina-coated samples. In addition, the onset temperature for sintering, and therefore the thermal stability of the mesoporous thin films, was found to be significantly increased by ALD coating (up to 800 °C for silica). Differences in the stabilizing effect were attributed to different interaction strengths of the coating materials with the mesoporous substrate. Significant improvement in thermal stability of mesoporous Rh₂O₃ thin films by ALD coating (by >300 °C compared to uncoated material) has also been observed by Dubraja et al.²⁴⁶ Pagán-Torres et al. reversed the strategy to stabilize mesoporous materials for applications under harsh conditions. They deposited catalytically active niobia onto the pore walls of mesoporous silica.³⁰² The hybrid material revealed improved hydrothermal stability, with catalytic activity superior to that of commercial niobia.³⁰³

Optimizing the pore size to achieve a high active surface area while providing efficient mass transport through the material and/or efficient charge-transport pathways in the walls is crucial for many applications. Unfortunately, arbitrary variation of pore size on the nanometer level is experimentally not possible. However, ALD can be used to reduce the pore size of pre-synthesized materials in a systematic fashion.³⁰⁴

Dendooven et al. investigated the decrease in pore size in mesoporous TiO₂ thin films with ink-bottle shaped pores upon deposition of HfO₂ using ALD. Achieving a conformal coating in such films is challenging, as there are only small channels within the walls that connect the mesopores and may lead to pore clogging during deposition. Using X-ray fluorescence (XRF) measurements, the authors confirmed that the inner surfaces are uniformly coated until the pore necks are closed, preventing further diffusion of molecular precursor into the architecture, as schematically shown in Fig. 28. Additional ALD cycles then only result in growth of a top (sealing) layer.^304,305 Comparable results were obtained by Cop et al., who investigated TiO₂ ALD coating of mesoporous Ce_xZr_1−xO_2−δ thin films.³⁰⁶

Fig. 28 (a) In situ XRF monitoring of filling of mesoporous thin films with an initial pore size of 4 nm during ALD of TiO₂. After pore filling, the growth rate decreases and TiO₂ only grows as a top layer on the films. (b) Schematic illustration of the pore filling process. Adapted from ref. 304 with permission.

Dendooven et al. also reported on the implementation of ellipsometric porosimetry onto an ALD reactor, allowing to determine the pore size, the film thickness and the mechanical properties in situ during ALD.³⁰⁷ In another work, the authors monitored the pore shrinkage by means of in situ GISAXS, also confirming uniform and conformal coating of the inner surfaces.³⁰⁸ Note that for sub-nm control of the coating thickness, the process parameters must be tailored.³⁰⁹ Too long exposure and/or insufficient purging result in precursor condensation (at the bottom of the pores) and pore plugging, whereas insufficient doses prevent full coverage of the inner surfaces. Especially for materials with a complex porosity, precursor diffusion limitations can lead to anisotropic ALD profiles, as reported for example by Pulinthanathu Sree et al.³¹⁰

Spatial control of deposition has been achieved using plasma-assisted ALD. Here, ALD precursors are employed that need to be activated through plasma irradiation. As shown by Jiang et al., this method allows sealing the pores of mesoporous materials at the immediate surface, as the plasma cannot penetrate into the inner structure.³¹¹

These and other examples demonstrate that conformal surface coating of mesoporous materials can be achieved using ALD.^313–316 However, in order to tailor the surface properties to a certain application, it is necessary to control the crystallinity of the deposited material (note that the lattice structure determines the electrical or catalytic properties, among others). Mitchell et al. characterized ALD-derived TiO₂ thin films on silicon substrates.³¹⁷ They found that native-oxide-free silicon promotes the growth of crystalline TiO₂ islands. In contrast, an initially amorphous TiO₂ film is formed on silicon with an amorphous interfacial SiO_x layer. Nevertheless, after exceeding a critical thickness, the amorphous TiO₂ crystallizes. Zscherp et al. also examined the nanostructure of an ALD oxide coating, namely, 9–20 nm ceria deposited onto porous YSZ thin films.³¹² TEM investigations revealed epitaxial growth of ceria with a columnar structure and excellent conformality (Fig. 29). An amorphous-to-crystalline phase transition with increasing layer thickness has also been reported by Celik et al.¹²⁹ They coated porous YSZ thin films by ALD of TiO₂. While the deposition of 6 nm TiO₂ resulted in the formation of amorphous material, a crystalline coating (anatase TiO₂) was found for a thickness of approx. 18 nm using the same deposition conditions (150 °C substrate temperature). Furthermore, the authors studied the influence of crystallinity on the surface protonic conductivity. The amorphous TiO₂ coating led to a lower protonic conductivity compared with bare YSZ thin films. However, it increased again for the sample with the crystalline surface shell (accompanied by a decrease in proton mobility at high relative humidity, as the reduced pore size resulted in complete filling with water).

Fig. 29 TEM images and mapping results showing the epitaxial growth of ALD-derived ceria on porous YSZ thin films. Reprinted from ref. 312 with permission.

Cop et al. further probed the electrical properties of mesoporous ZrO₂ thin films coated with ceria using different numbers of ALD cycles.³¹⁸ Characterization via ellipsometry confirmed the homogeneity of coating in the pores up to 100 ALD cycles, after which pore plugging occurred. XPS measurements revealed a large Ce³⁺ fraction in the coating because of the space-charge layer at the surface, which decreased with increasing thickness, as shown in Fig. 30. In addition, the total conductivity of the composite thin films was found to be strongly affected by the coating and its thickness due to the non-stoichiometry of the ceria layer. High Ce³⁺ concentration resulted in dominant electronic conductivity because of small polaron hopping of electrons between the Ce³⁺ and Ce⁴⁺ ions.

Fig. 30 (a) Ce³⁺ content of ceria deposited onto mesoporous ZrO₂ thin films as function of the ALD cycle number. The Ce³⁺ content decreases with increasing coating thickness. (b) Temperature-dependence of total conductivity for mesoporous ZrO₂ and CeO₂–ZrO₂ composite thin films (25, 75 and 150 ALD cycles). Adapted from ref. 318 with permission.

Collectively, these studies demonstrate the profound effect that the coating thickness has on the structural and electrical properties. Because the thickness can be controlled precisely, post ALD treatment allows for rational materials design.

6. Conclusions

Herein, developments in the field of ordered mesoporous metal oxides for electrochemical applications are reviewed. Advantages and disadvantages arising from the unique architecture, consisting of open ≤50 nm pores surrounded by interconnected crystallites, are discussed. In the last decades, much research has been carried out into assessing the potential of mesoporous metal oxides for energy applications, ranging from batteries and catalysis to sensors. The presence of an ordered 3-dimensional pore network is advantageous, as it offers a large number of active surface sites and facilitates the accessibility of the material and its internal structure to the surrounding medium. For instance, high-surface-area metal oxide electrodes may increase the charge-storage capacity of batteries due to pseudocapacitive redox reactions, while the porosity is capable of somewhat mediating the transport of reaction species in catalysis. It also endows the materials with a certain degree of mechanical flexibility, which helps accommodate volume changes during electrochemical cycling, among others. However, the favorable properties cannot only be attributed to the increased surface area and uniform pore size distribution, but also to the interconnectivity of crystallites in the wall structure, giving rise to short diffusion pathways for electron and ion transport.

An important aspect affecting the properties of mesoporous metal oxides is the existence of a space-charge region at the surface, which has been shown to exert a profound effect on the charge transport. Resistive oxide sensors are based on this working principle. Considering the surface space-charge region is also critical when evaluating the performance of mesoporous battery electrodes, as it directly affects the interfacial charge storage. In case of catalytic applications, the space-charge potential facilitates the separation of photogenerated charge carriers. This is beneficial for photoelectrocatalytic devices, but counterproductive in photocatalysis, where both redox reactions occur at the surface of the catalyst. However, the structural features of mesoporous metal oxides come at a price. For instance, the high surface area may lead to severe corrosion and adverse side reactions with liquid electrolytes or reactive gases. Furthermore, the nanoscale porosity negatively affects the thermal stability of the material, which in turn limits the temperature range for operation, especially in catalysis. In general, there is a complex interplay between the structural parameters, which makes it difficult to vary them independently (note that this is one of the major drawbacks of mesoporous materials).

Finally, this review article also summarizes recent work on post-surface modification of mesoporous metal oxides using atomic layer deposition (ALD). Because ALD allows conformal coating of complex substrates, it represents a promising technique to tackle the above-mentioned problems. Recent results confirm that surface coating helps to improve the thermal and structural stability. The thickness of the deposited material can also be controlled on the atomic level, since the ALD process relies on self-limiting surface reactions. Hence, the pore size can be adjusted without altering other structural parameters. In addition, the surface potential/chemistry can be manipulated via interfacial engineering, thus directly affecting the space-charge region and offering the possibility to tailor the electrical and electrochemical properties of mesoporous metal oxides for future device applications.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

Financial support by the German Research Foundation (to T. B., grant no. BR 3499/5-1 and to M. T. E., grant no. EL 863/3-1) is gratefully acknowledged. M. T. E. is grateful to the German Federal Ministry of Education and Research (BMBF) for funding the NanoMatFutur project NiKo (03XP0093).

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