Perspective on the Influence of Interactions Between Hard and Soft

Perspective pubs.acs.org/cm

Perspective on the Influence of Interactions Between Hard and Soft Templates and Precursors on Morphology of Hierarchically Structured Porous Materials Andreas Stein,* Stephen G. Rudisill, and Nicholas D. Petkovich Department of Chemistry, University of Minnesota, 207 Pleasant Street SE, Minneapolis, Minnesota 55455, United States ABSTRACT: When surfactant templates and hard templates are combined to synthesize silica and other inorganic solids with hierarchical pore structure, the mesopore development and orientation are influenced by interactions between the surfactant templates and the hard templates. Effects such as confinement, surface interactions, and curvature can be used to control the relative mesopore alignment. Such control is important for numerous applications that rely on access to the mesopores and incorporation of multiple components in a hierarchical structure. This review highlights how the interactions between molecular and block-copolymer surfactant micelles and confining surfaces can be used to influence mesopore architecture. After a brief review of block-copolymer systems, more complex systems will be considered that combine block-copolymers and inorganic precursors on a single surface, between two planar surfaces, confined in channels of porous anodic alumina membranes, in colloidal crystal templates, and in three-dimensionally ordered macroporous solids. The influence of confinement on external product morphology and on the distribution of multiple phases in porous materials with complex composition will also be described briefly. The review will conclude with a perspective on developments needed to facilitate further progress and new applications in this field of research. KEYWORDS: surfactant templating, colloidal crystal templating, porous anodic alumina membranes, confinement, interfaces, porous materials

1. INTRODUCTION Throughout much of the history of this journal, surfactant templating has been employed to synthesize mesoporous materials with high surface areas, tunable pore sizes, and numerous pore geometries.1−12 In these syntheses, the major role of the surfactants, whether they are low-molecular-weight ionic or nonionic surfactants or higher-molecular-weight, blockcopolymer based surfactants, is to organize precursor phases within micellar structures. A secondary role of surfactants in templated syntheses includes their ability to compatibilize precursor and solvent components. In subsequent developments, surfactant templates (so-called “soft templates” with structures that vary dynamically, depending on the synthesis conditions) were combined with more rigid, “hard templates”, including colloidal dispersions, colloidal crystals, porous membranes, and lithographic patterns to obtain porous materials with more complex structures (Figure 1). Here, each set of templates acts as a space filler or a structure director at a given length scale, and after removal of the templates, materials with hierarchical pore structure are obtained. Those materials are suitable for applications that require mass or charge transport at multiple length scales or applications that take advantage of the high surface areas and selectivities introduced by smaller pores coupled with better transport properties through larger pores. Materials with hierarchical pore architectures can also provide short diffusion paths inherent to nanostructured features and facilitate interfacial transport. An © XXXX American Chemical Society

illustrative example is given by complex electrode structures for supercapacitors or batteries, in which a mesoporous conductive phase may be filled with a secondary electroactive component and secondary pores remain accessible for penetration with an electrolyte.13 Other applications that can benefit from hierarchical porosity include artificial photosynthesis, photocatalysis, dye-sensitized solar cells, fuel cells, hydrogen storage, solar thermal storage, drug delivery, bioceramics, optical materials, and more.14,15 Several review articles and books on templated solids with hierarchical pore structure have been published in recent years.15−19 In many of these applications, the relative alignment of pores at each length scale is important. In membranes used for separation, for example, mesopore alignment perpendicular to the plane of the membrane rather than parallel to it is needed to facilitate mass transport across the membrane. Similarly, in solids containing mesoporous walls around larger macropores, the relative orientation of mesopores determines their accessibility and the ease with which these pores may be filled with secondary components. In single soft-template systems, Special Issue: Celebrating Twenty-Five Years of Chemistry of Materials Received: July 3, 2013 Revised: August 8, 2013

A

dx.doi.org/10.1021/cm402184h | Chem. Mater. XXXX, XXX, XXX−XXX

Chemistry of Materials

Perspective

2. BLOCK COPOLYMERS UNDER CONFINEMENT Before considering more complex inorganic and hybrid materials templated with surfactants in confinement, it is worth taking a look at confined surfactants by themselves, which have been studied in quite some detail over the past decade.22 In particular, di- and triblock copolymers, in which individual blocks can interact differently with their environment, provide good insight into confinement effects within a hard template. Some of the more common examples of block copolymers used in surfactant templating syntheses include those combining poly(ethylene oxide) (PEO) and poly(propylene oxide) (PPO) blocks, such as the Pluronic surfactants P123 (PEO 2 0 PPO 7 0 PEO 2 0 ) and F127 (PEO100PPO64PEO100), (poly(ethylene-co-butylene)-b-PEO) (KLE), and polystyrene (PS) based block copolymers, such as PS-b-PEO, PS-b-poly(butadiene), and PS-b-poly(methyl methacrylate) (PMMA). Di- and triblock copolymers microphase-separate into periodic, chemically distinct equilibrium domains with characteristic dimensions, L0, and a variety of morphologies (dense or perforated lamellae, cylinders, spheres, etc.).23 The microphase morphologies change when the amphiphilic block copolymers are placed in a confining environment, depending on the relative dimensions of the confining environment compared to L0 (commensurability),24,25 curvature of the host,26 interfacial interactions (surface energies), and entropic effects.27 Lamellar or cylindrical phases often align parallel to a surface, so that curved surfaces will also impose curvature on the mesophases and for very high curvature possibly prevent the formation of micelles. Local curvature of solid surfaces can impose orientational order on micelles, so that the morphology of micelles on curved surfaces may differ from that on planar surfaces.28 Confinement can also break the symmetry of a mesophase.29 Mismatch between dimensions (i.e., if the critical length of confinement is not an integer multiple of L0) introduces structural frustration to which the surfactant responds by compression or expansion and local phase changes.26 For strong confinement, insufficient space is available to accommodate fully developed micelles. The critical micelle concentration increases with smaller confinement size as a consequence of increased entropy loss.30 With weak confinement, the surfactant−wall interactions become the predominant factor. Interactions between surfactant molecules and a confining host control surface energies and are particularly relevant in hosts with high-surface areas, such as the porous solids often used in hard templating. In the case of block copolymers, differences in interfacial interactions for the different blocks can then be exploited to reorient mesophases near a surface. Furthermore, confinement leads to a reduction in entropy because movement of surfactant chains becomes more restricted. All of these parameters can be altered to modify mesophases from more extended micellar structures to spherical micelles or discrete surfactant molecules.27 In addition, kinetic factors, including rates of surfactant assembly and interface formation between surfactants, other system components, and the confining walls influence mesostructure assembly in bulk and confined environments. Thin films supported on a substrate provide a special type of confinement, where a solid support furnishes one interface with the film and air the other interface. Thus, for sufficiently thin films, the two interfaces lead to asymmetric surface boundary energetics. These boundaries can provide attractive and repulsive interactions with specific blocks in a block

Figure 1. Soft and hard templates. (A) Block copolymer phases, including spherical, cylindrical, gyroidal, and lamellar phases. Adapted with permission from ref 20. Copyright 2007 Elsevier. (B) Porous anodic alumina membrane. Reprinted with permission from ref 21. Copyright 2006 Wiley-VCH. (C) Polymeric colloidal crystal template.

the mesostructure of the surfactant-precursor phase is influenced by steric and electrostatic interactions with inorganic precursor components, so that organized phases (micelles) may be obtained even when the surfactant concentrations are below the critical micellar concentration.6 Therefore, phase diagrams of pore geometries in these combined organic−inorganic systems are different from those of the pure surfactant−solvent phases. Yet, a large amount of data has been assembled throughout the last two decades, and phase diagrams of mesostructured/mesoporous materials are quite well understood, certainly for mesoporous silica and carbon, but also for numerous other compositions. On the other hand, when a second template component is added, whether hard or soft, additional interactions arise at interfaces that change the phase behavior of the dynamic components (soft templates and nonsolidified precursors). In particular with hard templates, the abilities of different precursor and surfactant components to wet the surface of the hard template influences the distribution and arrangement of these components across the surface. As a result, both the internal and external morphologies of the products may be changed compared to the single template systems. Furthermore, in systems with complex composition, the distribution of phase components may be influenced by the presence of a hard template. In addition, interactions between hard and soft templates can impact the external morphology of the porous products. This Perspective will highlight examples that demonstrate how such interactions between hard and soft templates, especially confinement of the precursors and soft templates in the restricted space of the hard templates, can be employed to direct the product architecture and phase distribution at multiple length scales. Control over such interactions can help with the design of complex materials with intricate structures, where structural components may provide different functionalities or interact synergistically to enhance materials properties. B



Perspective

copolymer.31 Dynamic density functional theory applied to simulate microphase formation in melts of asymmetric block copolymer films indicated that microphase morphologies depend on the film thickness as well as the preference of each block for the confining surface.32 If surface interactions are neutral with respect to each of the blocks (i.e., the substrate does not have a preference for a particular block), microphase domains may form perpendicular to the surface, as long as the films are sufficiently thin.31,33,34 Equilibrium lattice-based Monte Carlo simulations applied to ternary amphiphile− solvent mixtures demonstrated the effect of surface properties of confining walls.35 Hydrophilic walls formed a water interlayer and produced phases similar to those found in the bulk. However, because the surfactants now occupied diminished space and reached higher local concentrations, phase transitions occurred earlier during sample drying than in the bulk. Hydrophobic walls reversed this trend as surfactants tend to adsorb at the walls, reducing the local concentration in solution and delaying mesophase formation. Neutral walls that have similar attraction for each component in amphiphilic molecules can enforce the formation of hemispherical micelles and alignment of mesopores perpendicular to the wall. Further alignment of block copolymers is possible on physically patterned (grooves)24,36,37 or chemically patterned substrates.38 If surfaces with parallel grooves are modulated, defects are introduced into block copolymer films that resemble dislocations in epitaxial thin film systems.36 Chemical patterns of polar and nonpolar regions can align block copolymers in such a way as to control registry with microphases and the surface patterns. For example, lamellar block-copolymer domains oriented perpendicular to the substrate were obtained if the period of the chemical pattern on the substrate coincided with that of the lamellar spacing of the block copolymer.38 Even defects in the underlying pattern translated into defects in the block copolymer film. The most ordered, substrate-directed films were obtained when the spacings in the substrate and the block copolymer agreed within better than 10%. Other types of confinement (two solid planar interfaces, channels, pores in colloidal crystals or other porous hard templates) will be discussed in later sections for more complex systems that combine surfactant templates with inorganic components.

of silicate oligomers to the block copolymers results in association of silicate with PEO blocks and enhances the incompatibility between PPO segments and PEO−silicate segments.27 In the case of cationic surfactants for mesophase self-assembly, monosilicic acid monomers replace the anionic counterions of the surfactant, which leads to growth of initially small spherical aggregates.39−41 Theoretical considerations suggest that the spontaneous curvature of the surfactant array is reduced as the adsorbed silica neutralizes electrostatic repulsions among surfactant headgroups.42 As silica condensation proceeds, the silicate oligomers begin to interact with multiple micelles, promoting aggregation of the micelles.41,42 Without optimization of surface interactions, mesopores in typical mesostructured silica films with 2D hexagonal channels (e.g., SBA-15 or MCM-41) are aligned parallel to a flat substrate surface, even when the substrate is water.43−46 Such orientation is a consequence of different interactions between the substrate and polar and nonpolar components of the surfactant. Several approaches have been used to enforce an orthogonal orientation of mesopores with respect to the plane of the film, the orientation desired for permeation of molecules through the membrane. One approach uses the application of strong magnetic fields during film formation.47−50 Alternatively, electrochemical potentials at an electrode surface may be used to produce orthogonal channel alignment in mesoporous silica films that were synthesized by potentiostatic electrodeposition on surfaces as varied as gold, copper, glassy carbon, and indium tin oxide.51,52 In this electro-assisted self-assembly process, silica condensation is catalyzed by cathodically generated hydroxide ions. It is likely that the vertical pore orientation is also facilitated by the fact that the growing film is bound by two interfaces (the electrode on one side and a water/ethanol liquid interface on the other side) that are more similar than interfaces with a solid and with air on opposite sites. Partial vertical alignment has also been achieved by shear in confinement using a continuous flow cell.53 With a well-chosen system of a gemini surfactant (N-acyl-Lglutamic acid with chains of 12−18 carbons), a costructure direction agent (3-aminopropyltrimethoxysilane), and silica precursors in the presence of ethanol, vertical channel alignment is possible on hydrophilic and hydrophobic substrates with wide ranges of surface energy.54 However, more common approaches aiming to control mesostructure orientation rely on adjustment of interfacial interactions, because interfaces provide the nucleation sites that direct mesophase growth toward the interior of the film.55 For example, vertical alignment of silica mesochannels was induced through π−π interactions between polyaromatic molecules that formed a discotic lyotropic liquid crystal template and similar π−π interactions with a graphite substrate.56 When interactions between the surfactant system and the interface are not as specific as in the π−π, kinetic control becomes more important. Kinetic control played a role, for example, in mesopore alignment on highly curved surfaces, such as isolated latex spheres. In a polymer-sphere/surfactanttemplated synthesis of hollow mesoporous silica spheres, the silica shells exhibited radial mesopore orientation as long as the condensation of silica was sufficiently slow.57 When the rate of silica condensation was increased, mesopores followed trajectories parallel to the latex surface. Interestingly, the charge on the latex sphere surfaces (positive or negative) had little influence on mesopore orientation.

3. COMBINED SOFT-TEMPLATE/INORGANIC PRECURSOR SYSTEMS UNDER CONFINEMENT 3.1. Interactions with a Single Solid Surface: Mesoporous Silica Films. Because control of the orientation of mesopores first became relevant for silica films intended for use as separation membranes, we will begin the section of confined surfactant-templated mesostructures with thin film geometries. As for the simpler surfactant systems, mesostructures are controlled by relative dimensions of the confining environment (here, film thickness) compared to critical dimensions of the mesophase, interfacial interactions, and entropic effects. In addition, in the more complex systems containing inorganic precursors (e.g., silica), the relative precursor:surfactant ratios, ionic strength, synthesis temperature, and aging time of the precursor influence mesostructure formation through kinetic control. In such systems, the free energy of mesostructure formation is a function of the free energies of the formation of an organic−inorganic interface, the free energies of linking the inorganic components and assembling the organic template, and the free energies of solvation and solvent organization.27 When PEO-PPO-PEO block copolymers are used, the addition C



Perspective

The importance of kinetic control was also demonstrated for a thin film coassembled on a silicon substrate from a PS-b-PEO block copolymer and silsesquioxane (SSQ) as the silica source.58 The silsesquioxane dissolved preferentially in the PEO domains. The volume fractions of the PS and PEO blocks and the composition of the mixture containing the silsesquioxane determined the mesostructure. Under conditions in which cylindrical blocks formed, the orientation of these blocks depended on the annealing time of the films. For short annealing times (15 h), the channels in the 300 nm thick films were oriented perpendicular to the film surface, except for 2−3 layers at the substrate−film and film−air interfaces, where cylindrical pores were oriented parallel to the substrate. This appeared to be a kinetically stabilized state, because, as annealing times increased, the channels started to tilt (64 h) and eventually were aligned parallel to the substrate surface (314 h). These particular mesostructures remained after condensation of silica and calcination to remove the organic template. In this system, the orientation of cylindrical regions was not significantly affected by modifying the interfacial energies at the substrate−film interface by depositing (3aminopropyl)triethoxysilane or gold layers on the silicon substrate. In other cases, though, it has been possible to use chemical substrate modification to alter substrate−precursor interactions and reorient mesopores in templated silica films. When a precursor for a mesoporous silica film with the triblock copolymer Pluronic P123 as structure-directing agent is deposited on a silicon substrate, the PEO blocks of the surfactant interact strongly with the substrate native silica layer on the substrate, and micelles form parallel to the surface. To avoid such interactions, Bardeau and co-workers modified the silicon surface with a hydrophobic monolayer of n-octadecyltrichlorosilane (ca. 2.3 nm thick) and deposited the precursor solution on the more hydrophobic surface by spin-coating or dip-coating.59 On the basis of atomic force microscopy data, they concluded that mesopores now formed normal to the substrate. They observed surface patterns of pits with diameters of 4.2 nm, coincident with the length of P123 extended molecular chains. The ability to employ substrate−precursor interactions for mesostructure alignment sets the stage for photomicropatterning of mesochannels in films. To demonstrate this, Seki and coworkers coated a support with a photo-cross-linkable polymer film containing liquid crystal-forming mesogens and cinnamoyl terminal groups (Figure 2A).60 The film was uniaxially aligned by exposure to linearly polarized UV light. Mesoporous silica thin films prepared on this substrate by dip-coating or static deposition exhibited mesochannels that were in the plane of the film and perpendicularly oriented to the polarized direction of the linearly polarized light. Such alignment of mesopores was observed for both cationic (cetyltrimethylammonium chloride, CTAC) and nonionic (Brij56, C16H33(OCH2CH2)nOH, n ∼ 10) surfactant systems, but not when the triblock copolymer P123 was employed. Chemical surface modification of the substrate need not be limited to molecular groups. Tolbert et al. illustrated epitaxial growth of mesoporous silica films on various substrates (glass, silicon, kapton) that had been modified with a thin layer of cubic mesoporous titania.34 The titania presents a honeycomb pattern (the (111) face) at the top surface to act as a regular pattern for epitaxial growth of block-copolymer templated films. This modified substrate can provide a good match with

Figure 2. (A) Photoalignment procedures for a photocross-linked liquid crystal film and surfactant/silica mesostructured film. Reprinted with permission from ref 60. Copyright 2006 American Chemical Society. (B) Schematic drawing and scanning electron microscopy (SEM) image illustrating the alignment mechanism of mesopores on a polyimide Langmuir−Blodgett film. Reprinted with permission from ref 61. Copyright 2007 Elsevier.

the spacing of mesochannels if a suitable surfactant is chosen (here Pluronic P103 and P123). In fact, the system is relatively forgiving, allowing a lattice mismatch of 15−18% between the cubic titania substrate and the hexagonal silica film grown on this substrate. With poorer lattice match, defect structures are introduced, in which pores trace away from the surface but then curve back toward it. Mesoporous silica is a suitable choice for epitaxial growth, as the hexagonal phase forms sufficiently slowly in this system to interact with the substrate so that patterning can occur. When salts or organic swelling agents are present, thin lamellar layers can form on the top of the film with vertically oriented mesopores. An interesting method of aligning mesopores involves mechanical rubbing of substrates modified with suitable polymeric thin films. Depending on the identities of the thin film and the surfactant template, the alignment process may depend on chemical organization of the surface or on morphological effects. In various studies, planar glass substrates were coated with ca. 10-nm-thick polyimide films, and the surface was rubbed in one direction by moving a fast-rotating buffing wheel across it.61−63 This process is used on an industrial scale for preparing alignment layers in liquid crystal displays. With the appropriate choice of surface coating, hexagonal mesochannels in surfactant-templated silica films grown on the treated substrate were aligned parallel to the plane of the substrate and normal to the rubbing direction and the direction of the polyimide polymer chains (Figure 2B). This behavior resulted from alignment of hemicylindrical micelles on the substrate. Such alignment required sufficiently strong interactions between the oriented polymer chains of the rubbed polyimide and the surfactant tail groups, e.g., hydrophobic interactions between a linear polyimide containing hexamethylene linking groups in the polymer chain and alkyl chains in cationic alkyl ammonium (CTAC) or nonionic polyoxoethylene ether surfactants (C12EO10, C16EO10). In D



Perspective

contrast, polyimide films without matching groups, with rigid chains, or with nonlinear structure did not provide as good uniaxial alignment of mesopores in continuous coatings. In fact, for a rigid polyimide underlayer, islands of mesostructured silica were formed with a common orientation parallel to the rubbing direction. In this case, partial alignment could not be attributed to molecular scale interactions, but rather to microgrooves, along which surfactant micelles assembled.61 3.2. Influence of Additional Side Walls. Stronger confinement by additional surfaces can provide further control over mesostructure. The influence of a hard template on inorganic oxide−surfactant−water interactions was recently modeled in a flat confining space, using lattice Monte Carlo simulations.64 Parameters such as component concentration, wall-to-wall separation, and hydrophilicity of the walls were taken into consideration. In agreement with the experimental observations summarized in this review, several morphologies different from those in the bulk65,66 were predicted, including a perpendicular lamellar arrangement and a bridge morphology. As the surfactant concentration increased, the phase changed from free surfactants to spherical micelles, parallel cylindrical micelles, lamellae, and eventually perforated lamellae. These arrangements formed single layers at small wall-to-wall separations and two or more layers at larger separations. Perpendicular lamellae were found at relatively high concentrations for specific wall separations. With weakly hydrophilic walls, under specific conditions, two layers of lamellae interconnected by short cylinders were produced in the simulations. Another recent simulation based on dissipative particle dynamics showed that, on a flat substrate, cylindrical mesochannels in a block-copolymer/silica mixture can be reoriented from parallel alignment with respect to the substrate to slanted alignment by increasing the difference in interaction parameters between each block and the substrate (Figure 3).67 However, perpendicular orientation was not achieved in this simulation. Instead, for large differences in interaction parameters, a thin layer nearest to the substrate remained parallel to the substrate, while above this layer, the mesochannels were slanted. When side walls were added, full perpendicular alignment was possible for wall separations that would correspond to ca. 100 nm in a real system using P123 as the block copolymer. Experimentally, such guidance by surrounding walls can be realized in mesoporous silica films by employing lithographic or colloidal patterning in two dimensions. Controlled patterning, in which the orientation of cylindrical mesopores can be “selectively engineer[ed]... at designed positions” is possible by lithographic patterning of the substrate.68 In one example, a silicon substrate was spin coated with a 0.5-μm-thick resist film and lithographically patterned over an area of several square millimeters. The patterned surface was then spin coated with a P123-templated silica precursor solution, which was converted to a mesostructured silica film by evaporation-induced self-assembly (EISA). In this process, the surfactant concentration is increased during solvent evaporation, resulting in the formation of micelles. At the same time, silica condenses. Disorder-to-order transitions result in the final lyotropic liquid-crystalline mesophase structure that extends throughout the film. The sides of the lithographically patterned channels had a strong influence on the mesopore alignment if channel widths were sufficiently small (∼100 nm), and consistent alignment was achievable over the whole

Figure 3. Different alignments of columnar mesopores induced by different affinities of blocks with the substrate. (A) εBS − εAS = 0, (B) εBS − εAS = 100, and (C) εBS − εAS = 300. Parallel walls with different interactions with the PPO block. (D) εBW − εAW = 200 and (E) εBW − εAW = 300. Side views, 45° views, and top and bottom views are shown. The air is represented by the green dot layer above the mesophases. Block A is colored olive green, block B blue, and the silica precursor red. The solid green surface is the isosurface defining the boundary between hydrophobic and hydrophilic blocks. Reprinted with permission from ref 67. Copyright 2012 American Chemical Society.

patterned area. When the channel width was comparable to channel height (0.5 μm), mesochannels formed from the bottom and side surfaces. For larger widths, mesopores were randomly oriented.68 As an alternative to lithographic patterns, one may use colloidal arrays for 2D confinement of mesoporous silica films. This approach was demonstrated for a P123-templated mesostructured silica film confined within a periodically spaced polystyrene hemisphere array.67 Monodisperse PS spheres were deposited on a silicon substrate with a thin Al2O3 overlayer by spin coating to form close-packed 2D arrays. Using oxygen plasma etching and annealing, these arrays were transformed into non-close-packed yet periodic hemisphere arrays, in which the spacing between spheres was controlled through the etching time. By spin coating and thermal processing of a precursor for mesostructured silica, hierarchically porous films were obtained with ca. 360 nm macropore arrays surrounding silica walls with 6−8 nm mesopores. To modify interactions between surfactant, precursors, and PS walls, the solvent polarity was changed, using water, methanol, or water/ methanol mixtures. These solvents associate more strongly with the more polar PEO phase of the PPO−PEO−PPO triblock copolymer. As a consequence, the effective polarity of the PEO blocks is modified by the solvents, and therefore, the interactions between these blocks and the confining surfaces are also altered. With methanol as the solvent, mesopores were aligned parallel to the substrate. Interestingly, the mesochannels E



Perspective

maintained a common orientation over domains that encompassed multiple hemispheres, i.e., over a long range they were quite indifferent to the colloids. This may have been a kinetic effect, resulting from rapid drying of solvent during spin-casting. With a mixed water/methanol solvent, transitional mesophases were observed, and with water as the only solvent, twisted mesopores formed with an average direction perpendicular to the substrate. Compression effects resulting from incommensurability between periodicity in the mesostructure and confining dimensions were observed in narrower regions of the membrane walls. Compression also resulted in some spherical mesopores. 3.3. Confinement of Films between Two Planar Surfaces. In most of the cases discussed above, the surfactant/precursor mixture is placed on a solid substrate with air at the top interface. Under typical synthesis conditions for mesoporous silica membranes the layer is therefore formed with asymmetric boundary conditions. Interactions of precursors/surfactant template are then different at the two interfaces. The interfacial energies at the top interface may be modified to better match the interfacial energies of surfactant domains at both interfaces, for example, by exposing the sample to chloroform/octane vapor mixtures58 or by carrying out the film formation process in solution.51,52 Even closer matching is possible by making both interfaces the same, e.g., by confining the precursor mixture between two layers of a secondary surfactant or between two slides with appropriate surface energies. The first approach is exemplified by confining a neutral triblock copolymer (P123) and a silica precursor between bilayers of a mixed cationic/anionic surfactant system (cetyltrimethylammonium bromide, CTAB, and sodium dodecylsulfate, SDS).69,70 The reaction product consists of free-standing mesoporous silica platelets with vertical channels (Figure 4). Interpretation of the product structures led to the conclusion that silica condensation occurred on the micelles of the triblock copolymer, not on aggregations of the ionic

surfactants. While similar platelets with perpendicular mesopores could be prepared when CTAB was replaced by similar surfactants with longer (18 C) or shorter (14 C, 12 C) alkyl chains, a critical ratio of cationic surfactant to SDS between 0.55 and 0.65 had to be maintained.69 For smaller ratios, platelets were not obtained. For larger ratios, mesopores were oriented mostly parallel to the platelets. Sheet thickness depends on temperature and pH, decreasing as the synthesis temperature is increased in the range from 35 to 50 °C and also decreasing as the pH is raised from 4 to 6.70 Pore size increases and wall thickness decreases with increasing synthesis temperature, similar to observations in bulk syntheses of SBA-15. These responses result from the pH and temperature dependence of the association between silicate oligomers and PEO blocks in P123. This association is most pronounced at low pH, when the greater availability of protons results in increased hydrogen bonding. The greatest segregation of PPO and PEO segments occurs at low pH and low temperature, resulting in thicker sheets and thicker walls. Perpendicular orientation is favored in the higher temperature range (40−50 °C) and at the higher pH values (5−6). The slightly negative charge of the confining bilayers ensures exclusion of negatively charged oligosilicate ions from these bilayers.69 Interactions between CTAB/SDS and both blocks in P123 are weak with no strong preference for one block, resulting in the perpendicular alignment of channels in the mesoporous platelets.70 Such neutrality of the substrate toward the specific chemistries of the surfactant template is also important when precursors are sandwiched between two slides,71 in particular for films thicker than ca. 100 nm.72 Rankin et al. controlled the alignment of hexagonally close-packed mesopores in P123templated mesoporous silica films by modifying a glass substrate with a cross-linked random copolymer of PEO− PPO with blocks that are shorter than or the same size as the template copolymer. Complete orthogonal alignment was observed with P123 if the film was sandwiched between two modified slides. Without a top slide (air provides a hydrophobic surface) or with an unmodified top glass slide (hydrophilic surface) the mesopores in large portions of the film were aligned parallel to the substrate. However, silica or titania films with a critical thickness in the range between 70 and 100 nm can form completely perpendicular channels on silicon wafers coated with a chemically neutral copolymer layer, even without a top slide.72

4. CONFINEMENT IN POROUS HOSTS 4.1. Confinement in Porous Membranes. In proceeding from planar substrates to porous hosts, we add curvature to the confining walls and also influence the synthesis kinetics, because solvent evaporation is slowed down by exposing less external surface. In the past decade, much progress has been made in understanding the influence of porous membrane walls on the architecture of confined mesoporous materials. Porous anodic alumina membranes (PAA, also called anodic aluminum oxide, AAO, or anodic alumina membranes, AAM) are most widely used as confining hosts, but polycarbonate (PC) membranes have also been employed. Both types are commercially available with a range of pore sizes. Unlike PAA membranes, PC membranes are stable to acidic reaction conditions. They can be employed for templating a range of mesoporous oxides in combination with ionic or nonionic surfactants as soft templates.73,74 For silica systems, the choice of membrane affects mesophase formation, because silica

Figure 4. Mesoporous silica platelets with perpendicular channels. (A) SEM image of the calcined platelets. (B) TEM image of microtomed sample showing perpendicular channels. (C) Proposed model for the formation of the perpendicular channel structure, showing a 2D hexagonal phase with vertical P123/silicate columns between a bilayer membrane structure of the ionic surfactants. Reprinted with permission from ref 69. Copyright 2004 Wiley-VCH. F



Perspective

Figure 5. Mesophases in PAA membranes. Schematic diagram reprinted with permission from ref 77. Copyright 2011 Wiley-VCH. Transmission electron microscopy (TEM) images of (A) a circular hexagonal phase (reprinted with permission from ref 21, Copyright 2006 Wiley-VCH), (B) a helical phase, (reprinted with permission from ref 81, Copyright 2010 The Royal Society of Chemistry), (C) a columnar hexagonal phase (reprinted with permission from ref 21, Copyright 2006 Wiley-VCH), and (D) a mixture of columnar and lamellar phases (reprinted with permission from ref 82, copyright 2008 Wiley-VCH).

precursors interact more weakly with PC walls than PAA walls. Under conditions when nonionic surfactants, such as P123 or Brij 56, tend to form circular mesochannels in PAA membranes (see below), products in PC membranes showed columnar, circular, and randomly aligned mesochannels.73 Besides mesoporous silica, some other compositions of templated materials grown in PAA membranes include mesoporous titania/silica and zirconia/silica composites75 and mesoporous carbons.76 Confinement in PAA hosts can prevent the structural distortion that occurs during template removal, for example, in thin film structures of mesoporous carbons. Syntheses of mesostructured materials in porous membranes are based on (i) immersion of the porous membranes in a sol− gel precursor with gelation in a hydrothermal reaction, (ii) EISA synthesis that is typically used for thin films and was described above, or (iii) vapor-phase infiltration of an alkoxide vapor into a membrane that is already infiltrated with a surfactant solution above the critical micelle concentration (cmc). These methods have recently been reviewed in detail by Bein and co-workers.77 Mesostructures formed in the confinement of PAA or PC membranes are quite varied. They include mesopores that are aligned parallel to membrane channels (i.e., perpendicular to the membrane surface, columnar hexagonal), perpendicular to the channel axis (tubular lamellar), forming helical channels or donut-shaped circular hexagonal channels (Figure 5). The mesostructure depends on the identity of the precursor, reaction conditions, and confinement dimensions as discussed in more detail below. Pore wall wettability also influences the morphology of mesostructured silica templated in PAA membranes. For example, with F127 as the structure-directing agent, hollow silica nanotubes were formed under specific conditions in unmodified PAA membranes. When the PAA membranes were coated with hydrophobic surface groups, nanofibers formed instead.78 Average diameters of mesopores near the outside of the tubes and fibers were also different, 12.3 nm in the tubes grown in the unmodified PAA and 9.8 nm in the fibers grown in confinement of hydrophobic PAA surfaces.

A systematic study of confined assembly of silica/nonionic block copolymer (P123) composite mesostructures within cylindrical channels of PAA illustrates how the mesopore structure depends on the channel diameter.79 For diameters from 55 to 73 nm, a straight core channel is formed that is surrounded by a double layer of concentric mesochannels. The geometry of the concentric channels varies and can be helical or donut-shaped. Between 49 and 54 nm the central core channel is absent. Between 34 and 45 nm an inner channel surrounded by stacked donuts is formed. With a 31-nm channel diameter, a single-layer helical channel develops. For even smaller channels, stacks of peapods form. From another study, inside 170−220 nm diameter PAA channels, hexagonally ordered mesopores followed the direction of the channels in the porous alumina membrane.80 As the confining diameter increases, a progression from spheres to straight channels to helices to concentric inner shells is predicted by self-consistent field calculations.79 Similar effects of confinement and the high curvature of channels in PAA have been observed for mesoporous silica systems with F127 as the surfactant.81 In this system, when channel diameters in PAA were reduced from 83 to 50 to 34 nm, the mesostructure changed from a multilayer, stacked helix structure to helices with a straight nanochannel core to double helix structures. Small modulations in channel diameter resulted in tilting of mesopore orientation, as a result of changing the commensurability of the mesopores and the host channels. In syntheses of periodic mesoporous organosilicas (PMOs) in confinement of PAA membranes, a cubic Im3̅m mesophase was observed when the nonionic surfactant Brij 56 was used as a structure-directing agent with a ratio of Brij 56/ethylenebridged silsesquioxane (the PMO building block) between 0.06 and 0.08.83 At ratios above 0.09, a circular mesophase was formed. PMOs prepared with P123 as the surfactant by EISA in PAA membranes contained columnar hexagonal or circular hexagonal mesopores, depending on humidity and evaporation rate of volatile components.84 In terms of reaction conditions for PAA-confined mesostructured silica, the mesostructure is sensitive to aging conditions, G



Perspective

in particular the rate of solvent evaporation,85 hydrolysis rate (water content, relative humidity),21,86 and surfactant concentration and identity.21 For example, samples based on tetraethoxysilane (TEOS) and P123 sols and aged with insufficient water for complete hydrolysis displayed longitudinal mesopores (parallel to channel axis) and those aged with excess water had channels perpendicular to the channel axis.86 For samples using Brij 56 or P123 as structure-directing agents, mesophase structure could be tuned between circular and columnar mesostructures through surfactant concentration and relative humidity.21 Circular hexagonal structures are kinetically favored over those with columnar hexagonal channels.82 In contrast, with CTAB as structure directing agent, mesopores are aligned parallel to channels,87 and the surfactant concentration does not have a large influence on the mesophase structure.21 The difference in mesostructure obtained with these structure-directing agents is related to the different mechanisms of micelle formation.88 As the solvent evaporates and mesoporous silica condenses, a contraction occurs that can affect pore morphology and decrease the contact between the confining support and the product phase.84 With P123 and Brij 56, micellar structures are formed abruptly after solvent evaporation.82 It has been proposed that ordered domains are nucleated at the channel walls, but the nucleation sites are not determined by the position along the channel walls. Domains then grow and may be followed by phase transformations (circular to columnar to mixed circular/columnar to lamellar; circular to lamellar).82,88 On the other hand, with CTAB as a structure-directing agent, mesopores form a columnar hexagonal mesostructure directly from the beginning of a synthesis, with pores running parallel to the channels of the PAA membrane after the critical micelle concentration has been reached. Most studies on PAA-membrane-templated mesostructured silica systems have employed water and ethanol-based solvent systems, in which all components were introduced into the template simultaneously. Recently a few variations of this approach have been carried out, which provided further control over mesostructure. In one study, tetrahydrofuran (THF) was used as the main solvent for syntheses of mesoporous silica with P123 or Brij 56-type nonionic surfactants.89 THF is nonselective toward the hydrophobic and hydrophilic segments present in these surfactants and prevents formation of micelles in water-free THF. Water molecules introduced from the aqueous HCl component and from condensation of silica promote micelle formation. In this system, relative humidity, therefore, has a particularly large effect on mesostructure formation. In another study, the channels of PAA or PC membranes were first filled with TEOS. The surfactantcontaining precursor was added in a subsequent step, and the membrane surface was sealed with a thin paraffin layer during the self-assembly process.90 The mesostructure orientation in the resulting silica nanotubes varied with the wall thickness, which was controlled by the relative amount of P123 solution added. Mesopores were oriented perpendicular to 15 nm thick walls, producing perforation in the walls, but circled the channel axis for 40−60 nm thick walls. A final example involves a phasepure columnar mesophase of silica in PAA that was achieved through a salt-induced phase change during evaporationinduced self-assembly.91 Although, as noted above, columnar mesophases can be obtained with the ionic structure-directing agent, CTAB, nonionic block copolymers (P123, Brij 56) normally form circular hexagonal 2D structures. Addition of

certain alkali metal halide salts can induce phase transformation to obtain phase-pure columnar mesophases even with nonionic block copolymers. It is notable that, in the salt-modified systems, a very thin amorphous silica layer formed at the interface between the mesophase and the PAA walls, a sign of increasing interactions between the silica phase and the walls. When walls were modified with hydrophobic surface groups, no ordered mesostructures formed in the 200-nm PAA channels. 4.2. Mesostructures Grown on Top of Porous Membrane Surfaces. Although the direct influence of a porous membrane-based template on mesostructure is largest within the membrane channels, overlayers on the membrane can also be influenced by the substrate. Yamauchi et al. spin coated precursor solutions for mesoporous silica with P123 onto PAA substrates with textured surfaces.92 The PAA substrates were prepared by multistep anodization and leaching procedures to form surfaces with conical holes that were uniformly spaced (∼100 nm), but had different heights (100− 325 nm) and, therefore, different aspect ratios (1.00−3.25). The organization of mesostructures within the voids depended on the cone aspect ratio and influenced the organization on the top surface above the voids (Figure 6). Circularly packed,

Figure 6. Top: cross-sectional SEM image and schematic of a mesoporous silica film with perpendicular mesochannels grown on a PAA substrate with conical holes. Oriented growth is induced by disordered mesopores within the conical holes. Bottom: TEM image and schematic of a mesoporous silica film with parallel mesochannels induced by circularly packed mesochannels within the conical holes. Reprinted with permission from ref 92. Copyright 2008 American Chemical Society.

donut-shaped mesochannels within conical holes were obtained with higher aspect ratio cones and induced mesochannel growth parallel to the surface. In contrast, disordered mesopores within conical holes obtained at lower aspect ratios led to vertically oriented mesochannel growth. In another study, Yamauchi and co-workers grew mesoporous silica in PAA with straight channels and included a continuous silica overlayer region.93 With CTAB as the structure-directing agent, mesochannels were aligned parallel to the long axis of the PAA channels. This orientation translated into the overlayer region, producing a large fraction of mesopores that were oriented perpendicular to the film in addition to some tilted mesochannels. This orientation was ascribed to the weak H



Perspective

silanol groups of silica, using a variety of lattices with different surface functionalities, and concluded that “a certain positive interaction between the template surface and the forming silica network has to be incorporated” in the synthesis system. In these dual-templated systems, polymer spheres acted merely as fillers and did not appear to influence the mesophase structure, perhaps because of the low degree of confinement between nonpacked spheres. Even very close to a polymer sphere surface decorated with PEO chains, the surfactant phase was only weakly aligned parallel to this surface. In a more recent study involving non-close-packed colloidal particles with sphere-tosphere separations that were often greater than the sphere diameter and using a silica precursor containing F127 to produce silica films with hierarchical porosity, mesopores with tetragonal symmetry were observed (space group I4/mmm).103 The tetragonal symmetry was not disrupted by the colloidal particles, even at high particle concentrations. Mesostructural order in dual-templated 3DOM/m silica systems using close-packed (fcc) colloidal crystals has been studied by various research groups, and a variety of mesophases have been observed. Unless otherwise indicated, the systems below employed PS colloidal crystal templates and suitable metal-alkoxide based (e.g., TEOS or tetramethoxysilane, TMOS) silica precursors. With the amphiphilic ionic liquid (1-hexadecyl-3-methylimidazolium-chloride) as the mesostructure-directing agent, a lamellar mesophase surrounded the macropores.104 This was the same phase as that produced in the bulk without CCT confinement.105 The cationic surfactant octadecyltrimethylammonium chloride produced hexagonal mesopores parallel to the surface of the PS spheres.106 Relatively large mesopores in the wall structures are produced using block copolymers of the KLE type with ionic liquids (e.g., 1-hexadecyl-3-methylimidazolium-chloride).107,108 The mesopores formed with these surfactants tend to be spherical both in CCT confinement and in the bulk. The ionic liquids can introduce additional smaller elongated mesopores in the structure, resulting in hierarchical mesoporous silica with a trimodal pore system. Such structural regularity was not achieved with mixtures of the block-copolymer and ionic surfactants, such as CTAB. Importantly, no phase separation was observed between the block copolymer and the ionic liquid, even at relatively large concentrations of the ionic liquid.108 The importance of interactions between polymer spheres and block copolymers on hierarchical structuring of porous silica systems has been noted in several studies.67,109−111 The observed mesostructure depends on the choice of surfactant, cosurfactant, and other processing conditions, such as drying rate (absence or presence of vacuum during drying). Anderson et al. carried out a systematic study of the influence of triblock copolymers F127 or P123 with alcohol cosurfactants on mesostructure in CCT confinement.109 In products using F127 with n-pentanol as cosurfactant, they observed layers parallel to the sphere surface: with P123 and n-pentanol, wormlike mesopores; with P123 and n-butanol, 2D hexagonal cylinders “parallel” to sphere surface; and with F127 and nbutanol, a “hole structure”, which was not further analyzed. Amorphous regions were noted around the windows and attributed to the insufficient space near the contact points between adjacent spheres for micelle formation. With F127 and ethanol, a cubic mesopore arrangement (Fm3̅ m ) was predominant, with some mesopores running parallel to the sphere surface.112

interactions between silica−surfactant assemblies and the substrate. 4.3. Confinement in Three-Dimensional Hosts. In three dimensions, confinement can become even more restrictive. In this section we will focus on confinement of surfactanttemplated mesostructured materials in colloidal crystal and three-dimensionally ordered macroporous (3DOM) hard templates,94 both of which provide a highly symmetric and periodic but curved confining environment with large enough dimensions to accommodate multiple repeat units of the mesostructure. The colloidal crystal templates (CCTs) consist of opal-like structures with face-centered cubic (fcc) arrays of uniform spheres, typically PS, PMMA, or silica spheres in the dual-templating studies reviewed here. When infiltrated with fluid precursors, they contain three distinct environments: larger octahedral interstices in the fcc sphere array, smaller tetrahedral interstices, and very narrow interconnecting regions between these two interstices. 3DOM materials or inverse opals are formed by infiltrating opal templates with precursor material and processing this to form a solid skeleton around the spheres. If the spheres are simultaneously or subsequently removed, an open, interconnected pore structure is obtained with void dimensions similar or slightly smaller than the templating spheres. In most cases, adjacent spherical voids are interconnected at the 12 contact points between spheres in an fcc array. These interconnected, periodic spherical voids provide a single type of environment for infiltrated fluids. 4.3.1. 3DOM/m SiO2. The original motivation for combining CCTs with surfactant templates was to produce materials with hierarchical pore structure, in which larger macropores provide easy access for guest species and, perhaps, furnish the periodic repeat structures desirable for photonic applications, and smaller mesopores introduce high surface areas and sizeselective guest uptake. Here, we will denote the resulting product structures as 3DOM/m materials, where the capital M refers to macropores and the lower-case m to mesopores. Although the first 3DOM/m materials were synthesized soon after the pioneering publications of colloidal crystal templating of macroporous solids appeared,95−99 the relative influence of the hard and soft templates on each other was not yet considered at that stage.100,101 Stucky and co-workers demonstrated the ability to fabricate hierarchically porous silica on three separate length scales, using micromolds for the 1000 nm length scale, colloidal crystals composed of monodisperse polystyrene spheres on the 100 nm length scale, and triblockcopolymer surfactants on the 10 nm length scale.100 With the surfactant F127, the cubic mesopore structure with an ordering length of 11 nm was the same as that found in the bulk. Stein et al. employed an ionic surfactant in a precursor mixture used for synthesizing mesoporous silica MCM-41 together with PSsphere colloidal crystal templating to form hierarchically porous silica in which 30−40 nm thick macropore walls contained ca. 2.3 nm mesopores and provided the material with a surface area >1300 m2/g.101 However, the mesopore geometry was not analyzed in this early study. Interactions between colloidal templates and surfactant templated phases were first discussed by Antonietti et al.102 Although that work was not aimed at producing materials with periodic interconnected macropore structure, it demonstrated a synthesis of hierarchically structured porous silica templated with combinations of nonionic alkyl octaethylene glycol surfactants and polymer latex sphere templates. The authors studied interactions between polymer latex surface groups and I



Perspective

membranes, in which mesopores were oriented perpendicular to the membrane, on average.67 Interesting variations of dual templating have employed polymeric silazanes or long-chain alkyl siloxanes as precursors, where the micelle-forming surfactant is covalently attached to the silicon atoms. As an example of the former, in a system using PS colloidal crystal templating with the precursor poly(ureamethylvinyl)-silazane and the block copolymer poly(isoprene-b-dimethylaminoethyl methacrylate), hexagonal mesopores formed parallel to the sphere surface.114 For the latter approach, colloidal crystal templates infiltrated with long-chain alkyl siloxane precursors (CnH2n+1Si(OSi(OMe)3)3) formed 2D hexagonal (n = 10) or lamellar mesostructures (n = 16) in the walls, with mesostructures being oriented along the spherical surface.115 Notably, the materials with lamellar mesostructures could be reversibly swelled by exposure to decanol. A last example of dual templating of 3DOM/m silica films with PS spheres and the surfactant CTAB involves an electroassisted preparation on a conductive, indium tin oxide coated substrate.116 The precursor was gelled by applying a cathodic potential to the electrode. Through varying the deposition time (on the order of tens of seconds), it was possible to control the extent of filling of the colloidal crystal template. After all template space had been filled, an overlayer continued to grow, in which mesopores were aligned perpendicular to the f ilm surface. Within the film, 2D hexagonal mesopores were aligned perpendicular to the sphere surfaces. Although the mesopore alignment of the overlayer that was 600−1000 nm away from the electrode surface matched that of mesoporous films electrochemically grown directly on the surface (i.e., without a CCT),51,52 it is not clear whether the directionality was induced by the electric field or perhaps by the polymer template surface.67 4.3.2. Other 3DOM/m Materials. Relatively few studies of other 3DOM/m materials prepared by dual templating with surfactants and colloidal crystal templates have considered correlations between the interacting templates and mesostructure. In 3DOM/m γ-alumina templated with polystyrene colloidal crystals and P123 as a soft template, hexagonal channels were aligned mainly parallel to the sphere surface, similar to the patterns for 3DOM/m silica.117 3DOM/m carbon synthesized by polymerization and carbonization of a phenol-formaldehyde (PF) precursor with F127 in a silica CCT contained spherical mesopores arranged in cubic symmetry.118 Because of interactions of the precursor with the silica template, little shrinkage occurred during thermosetting and polymerization, so that large cell parameters and mesopore sizes were obtained (18 and 11 nm, respectively). In another study, in which a PMMA colloidal crystal was used instead for the PF/ F127 system, the resulting mesopores were significantly smaller (ca. 3 nm).119 In addition to different interactions between the precursor mixture and the silica or PMMA templates, a higher pyrolysis temperature used in the latter study (900 °C vs 800 °C) may have contributed to pore shrinkage. By increasing the loading of the F127 surfactant, it was possible to change the phase of confined mesopores from cubic to 2D hexagonal with channels oriented largely parallel to the sphere surface.119 Mesopore diameters could be enlarged in PMMA colloidal crystal templated materials by adding tetraethoxysilane as an additional constituent to the PF/F127 precursor solution.120 With such a triconstituent precursor it was possible to form 3DOM/m C/SiO2 nanocomposites (5.2 nm mesopores) and remove the silica component to obtain 3DOM/m C with a

Li et al. analyzed monolithic 3DOM/m silica materials synthesized with a PMMA colloidal crystal template using the surfactants P123 or Brij 56 with dodecane as an optional swelling agent (Figure 7).111 They observed wormlike

Figure 7. TEM images of 3DOM/m SiO2 fragments taken from monoliths that were synthesized by dual templating with a PMMA CCT and the following surfactant systems: (A) Brij 56, wormlike mesopores. (B) P123, 2D hexagonal mesopores parallel to sphere surface. (C) Brij 56/dodecane, cubic mesopores. (D) Brij 56/ dodecane, 2D hexagonal mesopores perpendicular to sphere surface. Reprinted with permission from ref 111. Copyright 2007 American Chemical Society.

mesopores with Brij 56 and no vacuum drying. When vacuum drying was applied after CCT infiltration, cubic (Pm3n) or 2D hexagonal mesopores perpendicular to the sphere surfaces formed with Brij 56 and various concentrations of dodecane, whereas 2D hexagonal mesopores parallel to the sphere surfaces formed with P123. In syntheses of 3DOM/m silica monoliths the entrance size of macropores can be enlarged by causing mesopores in the walls to shrink by increasing the acid concentration, which alters the extent of hydrolysis in the precursor, and also by increasing the precursor hydrolysis time.113 Several important points regarding the relationship between the colloidal crystal template and micellar arrangements were noted by Li et al.111 For a cubic mesophase, the curvature of the polymer sphere surface can influence the packing order in the interstitial region, resulting in a lattice parameter change or possibly a phase transition. In the larger interstitial regions (octahedral holes), the cubic mesophase is quite regular. However, in the narrow strut-region around windows between adjacent cages, the mesophase is less ordered, caused by space limitations or possibly interfacial effects. For 3DOM/m SiO2 with columnar mesopores, interfacial effects influence micellar alignment more strongly, permitting control over the orientation of 2D hexagonal mesopores via the solvent system. This tunability of mesopore alignment between colloidal particles was later utilized to design hierarchically porous silica J



Perspective

nally to the silica sphere surfaces. The repeat spacing is less in these domains (51−72 nm) than in the bulk (88 nm) and depends on the colloidal crystal sphere size. The lower molecular weight cylinder-forming block copolymer produced a “golf-ball” type morphology. The mesophase architecture could be altered by modifying silica surfaces with PS or PFEMS groups to influence wetting properties of the porous silica host and form preferential interactions with one or the other block (Figure 8).127

bimodal distribution of mesopores (5.4, 7.3 nm) or the carbon component to generate 3DOM/m SiO2 with 6.4 and 7.5 nm mesopores. In analyses of the role of surfactant−hard template interactions, the bulk morphology of the material should be considered.120 Several of the above 3DOM/m SiO2 and 3DOM/m C materials syntheses targeted monolithic materials rather than thin films. Mesostructure development relies on evaporation-induced self-assembly, and the architecture depends on relative humidity, evaporation rates, and thickness of the materials.121−123 In the case of thin films, evaporation occurs relatively quickly, favoring self-assembly of surfactant micelles with the other precursors. In thicker, monolithic systems with longer diffusion paths through the colloidal crystal template, evaporation is less efficient. The generation of more ordered mesostructures then requires longer evaporation times or pumping to remove short chain alcoholic solvents that can destroy the liquid crystalline phase.124 4.3.3. Confinement in 3DOM Hosts. 3DOM materials, in turn, lend themselves as host materials to investigate and control the mesostructure of confined surfactant and silica/ surfactant systems. Pioneering work in this area was carried out by Ozin and co-workers who synthesized mesoporous silica in the confinement of 3DOM PS using the nonionic surfactant C12H25(OCH2CH2)10OH.125 The confined silica exhibited a hexagonal mesochannel structure. In the center of the templated spheres, the mesopore alignment did not appear to be influenced by the polystyrene mold. However, close to the interface between mesoporous silica spheres and PS, channels were aligned along this interface. It was suggested that the formation of mesostructured cylinders began with alignment of micellar cylinders parallel to the PS walls and that this orientation was maintained during silica condensation. Ozin et al. also demonstrated the effectiveness of both direct and inverted silica opals for probing 3D confinement effects of block copolymers.126 The diblock copolymer polystyrene-blockpolyferrocenylethylmethylsilane (PS-b-PFEMS) forms a lamellar structure with 46 and 42 nm thick layers of PFEMS and PS lamellae, respectively, giving a domain spacing, L0, of 88 nm. This diblock copolymer is particularly suited for transmission electron microscopy investigations because of the high electron density contrast between the PS and iron-containing blocks. When confined within inverse opal silica with 210 nm macropores, the diblock copolymer forms layers that follow the contour of macropores, thereby producing concentric shell structures. Some more extended communication with copolymer material in adjacent interstices was implicated in the structure formation mechanism. With even greater confinement, i.e., within a silica colloidal crystal, lamellae orient themselves perpendicular to the sphere surfaces, demonstrating microphase separation at very small dimensions, similar to the 3DOM/m SiO2 system with Brij 56 and dodecane.111 The fact that the thickness of the skeletal walls was similar to the dimensions of each block (42−46 nm) likely contributed to the orthogonal alignment. Follow-up research compared the effects of confinement on block-copolymer phases that form cylinders in the absence of confinement (lower molecular weight PS-bPFEMS with ∼29 nm diameter and ∼43 nm repeat distance) vs those that form lamellae (higher molecular weight PS-b-PFEMS with 88 nm repeat distance).127 For the latter, curved lamellae were found in domains of tetrahedral and octahedral holes of the face-centered cubic colloidal crystal host, and in these locations, the lamellar domains oriented themselves orthogo-

Figure 8. Comparison of PS-b-PFEMS diblock copolymer phases in (A) a silica colloidal crystal, (B) 3DOM silica, and (C) 3DOM silica, surface-modified with PS. The images on the left are dark-field TEM images. Reprinted with permission from ref 127. Copyright 2008 American Chemical Society.

A few recent studies have also considered confinement effects for more complex precursor systems inside 3DOM hosts. When mesoporous titania was grown by a surfactant-directed sol−gel process inside 3DOM titania, the 3DOM host helped to increase the crystallinity of the mesoporous product, mitigated shrinkage of mesopores during calcination, and caused larger mesopores to form in confinement.128 Mesoporous titanium carbide/carbon opals have also been assembled in confinement of 3DOM SiO2, which led to the formation of hexagonal mesopores with p6m symmetry in the templated spheres.129 While the examples in Section 4 illustrate that mesostructure of porous inorganic materials is controllable through confinement in 3D hard templates, predictive control is still limited. The existing literature has identified some predominant factors that influence self-assembly of soft-template/inorganic precursor mixtures in confined systems (e.g., hard template surface moieties, solvent polarity/chain length, condensation kinetics; see also Section 3.2). However, relatively little work has been done on quantifying and characterizing the extent that these factors control self-assembly and what structural features are accessible via confinement. Additional guidance from theory and modeling approaches is clearly needed, particularly for 3D confined systems, to enable further design of materials architecture through dual templating methods. K



Perspective

Figure 9. (a) SEM images of Pechini-derived Ce−Mg oxides with varying ethylene glycol (EG) and citric acid (CA) content and differing microstructure. Images are denoted by the molar ratios of components within the precursor in the form EG:CA:TMI, where TMI is the total metal ion content. Gel compositions with a 1:1 ratio of CA:TMI (x11) form spheres at different levels of growth based on the ethylene glycol content. Gel compositions with higher ratios of CA:TMI (X21, X31) result in bicontinuous networks (3DOM material interspersed with void space) in varying extents of spinodal decomposition. (b) Graph of precursor compositions for Ce0.5Mg0.5O2 materials depicted in (a) by EG and CA content, normalized to TMI. Reprinted with permission from ref 134. Copyright 2013 American Chemical Society.

5. INFLUENCE OF HARD TEMPLATE ON EXTERNAL MORPHOLOGY Beyond regulating the internal mesostructure in confined surfactant−inorganic precursor systems, hard templates can also influence the external morphology of the porous products. In the trivial cases, when monolithic CCTs or opaline films are used as the template, the products may assume the overall shape of the hard template. However, in special situations, interfacial interactions between the evolving precursor system and the confining template may results in other, quite interesting external morphologies, such as uniform spherical particles. To understand such systems, we will need to briefly review the phenomenon of polymerization-induced phase separation (PIPS). Phase separation controls the production of mesostructure from block copolymers. However, a variety of larger-scale features are also accessible via PIPS. This process was first discovered in the production of silica gel from silicon alkoxides, where condensation of silanols leads to the formation of progressively larger and more networked siloxane oligomers.130 These oligomers gradually become more insoluble as they grow and separate out as a distinct phase from the solvent. If the growing silica phase can be stabilized relative to the solvent, nucleation does not occur, and rather, the system undergoes a spinodal decomposition, forming a bicontinuous two-phase system. The structure can then be dried and/or calcined, leaving a macroporous silica material. The solvent phase, once

removed, generates the void spaces between the condensed silica gel. Of course, if we consider only a silica gel formed from water, alcohol, and silicon alkoxide, the silica nucleates, grows, and finally condenses into a solid. Any pores will be disordered, only the result of “missed connections” between silanol oligomers. No spinodal decomposition occurs between solvent and silica, rather, the silica simply forces the solvent out as it condenses. However, if a polymer or organic additive is introduced, macroporous structures arising from phase separation are frequently observed. If the additive readily forms strong hydrogen bonds with the condensing silica oligomers, the mixture separates into a solvent phase and a silica-rich phase containing adsorbed polymer. Phase separation is also possible with poor hydrogen bonding additives, but they must make up a larger constituent part of the precursor, as the predominant phases must be of the additive and a silica-rich phase, with solvent present in both. The effectiveness of a particular additive and the concentrations at which desirable results are obtained are controlled by the functionality and, thus polarity, of the additive.131 One approach to form hierarchical, phase-separated macroporous networks involves the use of block copolymers as the additives in the precursor mixture. The block copolymers in question must be selected for strong hydrogen bonding with silica oligomers, in order for the soft-templating material to be associated the silica-rich phase, as opposed to separating out into the solvent phase. Additionally, micelle formation within L



Perspective

carbon was predominant in the tetrahedral interstices and bridging channels between the two types of interstices (Figure 10A). Furthermore, the carbon-rich regions exhibited concave

this phase must be induced. This is accomplished either through the addition of a strongly hydrophobic substance, such as 1,3,5-trimethylbenzene,132 around which micelles can form, or simply by increasing the water and copolymer concentrations relative to the silica source.133 Returning to the influence of hard confinement, PIPSgenerated macroporous networks have been observed in Pechini-type gelation processes used to synthesize ceria-based mixed oxides within a polymeric colloidal crystal template.134 The most common form of Pechini reactions are polyesterifications of metal-citric acid chelates with ethylene glycol as a cross-linking molecule.135 Whereas in the previous example, a soluble polymer additive was used to mediate the phase separation between the growing gel and the solvent, in this case, an insoluble hard template of PMMA spheres provides the same function. The hard template generates porosity in the gel network, as well as affecting the final morphology of the gel. Additionally, much of the structural control within the phase separation comes from creating or removing stoichiometric imbalance between the monomers in solution, which has a substantial effect on the final average molecular weight of the gel postgelation. As a result, it is possible to adjust the product morphology from continuous 3DOM structures, to a fragmented 3DOM structure with randomly introduced large macropores (from phase separation) and uniform macropores (from the polymer spheres), to uniformly sized 3DOM microspheres (Figure 9). The product morphology depends on the relative ratios of ethylene glycol:citric acid:metal ions used in the precursor, and on the presence of the colloidal crystal template. Without that template, completely dissimilar, disordered morphologies and no microspheres were observed under otherwise identical conditions. The hard-templated PIPS technique has been applied to other oxide compositions but so far produces microspheres only for a few select compositions under well-defined synthesis conditions. Yet for those compositions, it is an inexpensive and probably scalable approach to control the morphology of porous ceramics.

Figure 10. (A) TEM image of 3DOM/m LiFePO4/C samples pyrolyzed at 600 °C. Darker regions correspond to the LiFePO4-rich phase in octahedral interstices, lighter regions to the C-rich phase in tetrahedral interstices. The inset is a SAED pattern for the corresponding sample region, confirming the crystalline nature of LiFePO4 in the composite. Reprinted with permission from ref 136. (B) TEM image of 3DOM/m TiO2/C in which TiO2 nanoparticles (darker spots) are distributed uniformly throughout the carbon phase.137 This material was prepared in a PMMA CCT, using the same PF resol precursor, water and ethanol in a mixed solvent, and HCl components as for 3DOM/m LiFePO4/C. The titania precursor was titanium isopropoxide complexed with trifluoroacetic acid and the surfactant P123.

surfaces following the curvature of the CCT mold, whereas LiFePO4-rich regions were convex, indicating that strong cohesive forces caused the formation of a surface with minimal specific surface area. These observations contrast with 3DOM/m TiO2/C composites also containing nearly 30 wt % carbon (Figure 10B),137 where TiO2 nanoparticles decorate the skeletal walls of the macroporous composite uniformly throughout the whole structure. These materials were prepared by a very similar procedure to that used for 3DOM/m LiFePO4/C, starting with a homogeneous precursor that was infiltrated into the CC template and aged before pyrolysis following the same program. The difference in phase distribution for these two systems is likely to be a consequence of the different interactions with either precursors or products and the template surface. In spite of a small surface charge from the initiator, the PMMA hard template is relatively hydrophobic, often requiring that a cosolvent, such as ethanol, be added to aqueous precursors to permit better surface wetting. With this cosolvent, the homogeneous mixed precursor solution infiltrates the template quite well through capillary forces. However, internal microphase separation is then possible, in which less polar components interact more strongly with the PMMA surface and more polar components reorganize to minimize surface interactions. In confinement, this is possible if the polar components occupy interstitial positions with smaller surfaceto-volume ratios, i.e., the larger octahedral holes. In the 3DOM/m LiFePO4/C composites, LiFePO4 (present in greater proportion) was therefore predominantly found in those positions. This distribution was confirmed by calcination of the product, which produced discrete LiFePO4 spheres.138 These phenomena of microphase separation and preferential site occupancy will likely be found in other systems consisting of hard templates with complex precursors and can influence the materials properties. In the 3DOM LiFePO4/C system

6. INFLUENCE OF TEMPLATES ON PHASE DISTRIBUTION In colloidal crystal templated syntheses that are based on multiple precursors or involve the formation of multiple phases, phase separation of components can occur, resulting in spatial variations of the product composition. An unusual example of such phase separation was found in the “one-pot” synthesis of 3DOM/m LiFePO4/C, a composite material that can be used as a cathode in lithium ion batteries.136 The material, which contained ca. 30 wt % carbon, was synthesized within a PMMA colloidal crystal, using a complex precursor mixture containing a prepolymerized, aqueous phenol−formaldehyde sol, HCl, the surfactant F127, FeCl2, LiCl, ethanol, and H3PO4. The surfactant was intended to introduce mesopores into the skeletal wall structure, but it also served the important role of preventing complexation between phenol and Fe2+, which would raise the viscosity of the precursor to a level that would prevent infiltration into the hard template. At the stage of capillary infiltration into the colloidal crystal, the precursor was a homogeneous solution. However, after aging at elevated temperature to increase the extent of cross-linking in the polymer, followed by pyrolysis to form the LiFePO4/C composite and remove the PMMA template, the solid product was no longer homogeneous. Instead, LiFePO4 had accumulated mainly in the octahedral interstices of the fcc template and M



Perspective

described here, they resulted in an efficiently “carbon-wired” active phase, improved the conductivity of the material by several orders of magnitude compared to the bulk phase, and led to excellent performance as a cathode material for lithium ion batteries.

remains a challenge. Multicomponent composite systems, in particular, are often difficult to synthesize. However, they offer tantalizing prospects for substantially improved functional materials, provided the complex interactions between precursors and the template can be optimized. Third, the use of novel surfactants and hard templates for this approach can be problematic from a synthesis standpoint but can potentially unlock new pore size ranges and geometries. Fourth, extensive characterization is often required to fully analyze the important structure−property relationships in these systems. Characterization methods are needed that can effectively acquire data for the analysis of complex pore structures and spatial variations in composition. Finally, and perhaps the greatest challenge of all, is the development of inventive characterization and computational strategies that can probe the complex interplay of forces that govern self-assembly in systems that contain soft and hard templates. It is likely that any breakthroughs in our understanding of self-assembly will allow for much greater control over meso- and macrostructures in many different systems, allowing us to tailor materials more effectively for specific applications. The next few paragraphs will explore these challenges more in-depth from the standpoint of further developments in synthetic techniques, state-of-the-art characterization techniques, and theoretical and computational modeling of these multicomponent, multiscalar systems. In terms of synthesis, most systems in which confinement effects on mesostructure have been demonstrated involve silica or organosilica. This is not surprising, as mesostructured silica systems are the best-developed compositions even as bulk materials prepared without any confinement. Control of these structures is comparatively easy, because they are noncrystalline, so that competing effects of crystallization and grain growth do not have to be considered. Similarly, such effects are absent in carbon systems, and indeed, as we saw, dual templating is feasible to influence pore alignment in mesoporous carbons. There remains a need then to develop methods that permit better alignment of mesochannels in other oxide and nonoxide systems. As indicated above, some progress has been made using confinement in porous PC membranes for oxides including titania, zirconia, alumina, niobia, and tantala. These compositions are also feasible for dual templating with surfactants in other hard templates, but they will require special synthesis techniques that prevent destruction of the mesopores by excessive grain growth. Given the rapid progress in syntheses of single soft-templated mesostructured oxides and phosphates and other compositions, it is likely that we will also see these techniques applied to more dual-templated systems. New synthetic capabilities are at the horizon, with ongoing developments in both bottom-up and top-down methods. From the top, higher resolution lithographic methods and 3D patterning techniques supply access to new hard templates with smaller confinement dimensions or greater geometric control. From the bottom, new building blocks, such as nanoparticles of various sizes and shapes, are available for assembly by soft and hard templating.139 By extending the range of compositions of hierarchically porous materials, and by developing a better understanding of interactions between multiple templates and complex precursors, the selection of applications for these systems can also be extended. Another target will be to maintain control of mesostructure orientation over larger length scales, for example, in the case of membranes to achieve larger areas over which consistent alignment is possible.

7. CONCLUSIONS AND PERSPECTIVES Over the past 15 years, our ability to control internal mesostructure with respect to an external morphology has increased significantly. Much of the progress has been driven by the target to fabricate membranes with perpendicular mesopore structure by simple, efficient self-assembly methods. But along the way, we have accumulated knowledge that enables us to design and prepare more complex materialsmaterials with hierarchical structure and controlled pore size distributions covering multiple length scales, porous materials with controllable shapes, and selective placement of discrete components. Let us briefly look back and summarize what we have learned. The influence of solid surfaces can be exploited to direct the architecture of mesostructures of silica and other oxides, to obtain mesophases different from those found in the bulk, and to achieve pore orientations in specific directions. In addition to kinetic considerations that are subject to the chemistry of the precursor system, effects of confinement depend on the commensurability between the equilibrium domains in the mesophases and the confinement dimensions of the support or hard template. For block-copolymer surfactants that have been investigated to date, these effects are most pronounced for confinement dimensions up to about 100 nm, although they can reach a little further. For strong confinement, entropic effects predominate, leading to structural frustration and even preventing the formation of micelles. For weak confinement, interfacial effects are most important and can be controlled through both chemical and physical surface modification of the confining walls. In the range from ca. 30 to 100 nm, directional control is strongest, allowing one to synthesize films and membranes with desired mesopore orientations. To achieve thicker membranes, it is convenient either to employ PAA hosts or provide guide-posts (e.g., 2D colloidal patterns or lithographic patterns) at separations in the 100 nm range. Because of the tremendous ongoing progress in patterning and controlling structures at this length scale, there are now new opportunities to study confinement effects in more detail and develop structuring techniques based on selfassembly appropriate for device fabrication. In three dimensions, opals or 3DOM hosts provide suitable confinement to achieve monolithic materials with hierarchical porosity and, in a few cases, porous microspheres through confinementdirected, polymerization-induced phase separation. Looking forward, there are distinct challenges that the field as a whole will actively endeavor to overcome. The innovative solutions developed to tackle these challenges will undoubtedly yield a panoply of novel materials and a much deeper understanding of the processes behind self-assembly. While there are a number of smaller challenges specific to all the various systems described in previous sections, several overarching challenges are present. First, the high degree of structural and compositional sophistication achieved by selfassembly approaches has been demonstrated to date largely on the laboratory scale. Devising ways to move to the pilot plant and industrial scale is a key next step. Second, extending hard and soft templating approaches to include more exotic compositions, such as nonoxide and composite materials, N



Perspective

A better understanding of interactions between multiple fluid phases during self-assembly and processing could facilitate the progression from time/labor intensive multistep processes to single/few step processes. For example, prototypes of 3Dinterpenetrating batteries, needed for on-chip power with little available mounting area, have been demonstrated,147,148 but their fabrication is currently too complex for scale-up. These battery designs require multiple components: two electrodes with current collectors, separated by an electronically insulating, ionically conducting membrane layer. Rather than assembling components one after the other, self-assembly of all components in a single step would be desired.149 Co-assembly of multiple phases is already possible (e.g., 3DOM/m C/SiO2, which contains tricontinuous networks of a conductive component, an insulating component, and a pore system capable of being filled with electrolyte),120 but practical systems will require further development. Other applications that will benefit most from ongoing developments of hierarchically structured porous materials are those that depend on significant interfacial interactions, selective mass transport of guest species through the porous network, effective charge transport through components of the solid skeleton, and/or ready access to high available surface areas. Besides the energy storage applications mentioned above and those listed in the introduction, such applications include controlled drug release, theranostics (i.e., combining therapeutic treatments like drug release with diagnostic capabilities in a single material), separation, and catalysis. The resulting hierarchically structured materials can then be further functionalized with surface groups suitable for each particular application. Because cost-effective scale-up is critical for most of these applications, self-assembly approaches are most attractive. In the near future, scale-up will most likely begin with 2D geometries, i.e., thin films that are prepared in continuous processes by coating techniques. These are most easily implemented with a solid interface at one side and an air or solvent interface at the other side, but they could also involve combined top and bottom supports. Belts or rollers used in these coating processes could be patterned or chemically surface-modified to control the orientation of mesostructure in the film. On a much smaller, lab-on-a-chip scale, electronically addressable confining walls may allow dynamic orientation of mesostructures, which can eventually be frozen in space in the target solid. Ultimately, the most interesting constructs will involve 3D arrays that are currently not easily fabricated by lithographic or printing techniques. For such constructs, controlled manipulation of interactions between hard and soft components may be the programming step for the 3D nanoprinters of the future. Today, this is still a dream, but if the momentum in the field continues, the realization of this dream may not be so far off.

Characterization of hierarchically structured porous materials involves combinations of standard techniques, including SEM to analyze external morphology, TEM and high-resolution SEM for internal structure, gas sorption and mercury intrusion analysis to evaluate pore texture, small-angle X-ray scattering methods to characterize mesostructure geometry, wide-angle Xray diffraction and various spectroscopic techniques for phase determination and compositional analysis, etc. Most of these techniques provide an average description of the material, and even with those that do not (electron microscopy), the assignment of pore symmetries is not trivial with many of these hierarchical structures. In 3DOM/m silica, for example, in which mesoporous silica walls surround ordered macropore arrays, overlap of multiple walls can complicate the interpretation of TEM images. For such samples, it has been helpful to observe thin samples from multiple viewing angles.111 Even better models of pore structure in 3D are obtained by electron tomography, a technique whose use is becoming more widespread and which provides spatially resolved 3D representations of mesopore structures. 140−143 It allows quantitative characterization of local morphology in the types of materials discussed in this review, permitting analyses of very fine details of confinement effects. In 3DOM/m silica materials or surfactant-infiltrated 3DOM hosts, for example, electron tomography would allow one to more easily compare alignment of mesochannels between adjacent cages to analyze longerrange influences of the template. Additional progress in achieving control of structure at multiple length scales in materials with hierarchical porosity will benefit from theoretical and simulation input. A cooperative templating mechanism that relies on the influence of silica on micelle formation was proposed early in the development of the field of mesostructured materials,2,144,145 investigated in detail experimentally (see references in ref 40), and was supported by molecular dynamics simulations and other theoretical studies.42,65,14639−41 Only recently has the influence of confinement been added to computional models of inorganic oxide− surfactant−solvent systems.64,67 For simple systems, commercially available software has made modeling accessible even to the synthetic chemist, providing predictive capabilities.67 Yet, there is a need for expansion to provide further guidance to the synthetic chemist. The increasing structural and compositional complexity at multiple length scales desired in advanced applications will require development of fast, new multiscalar computational tools to address systems with gradient, fractal, and other nonperiodic structures. Some of the issues that need to be addressed include choice of surface modifiers for the confining environments, additional patterning on confining walls, combinations of internal fields with externally applied fields (electric, magnetic, flow) to orient mesostructures, and a better understanding of interaction lengths in various confining geometries. More detailed models need to be developed to address questions of external morphology control and influence of a template on phase distribution for the systems discussed in Sections 5 and 6 of this review. Why are porous microspheres observed with only few compositions, in particular those capable of undergoing changes in oxidation state (e.g., CeO2based materials)? What causes the unique phase distribution in the 3DOM LiFePO4/C structure, and can this be achieved in other materials to control phase distribution locally? Answers to these questions could then enable site-specific placement of nanosized-components in complex architectures.

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest. Biography Andreas Stein is a Distinguished McKnight University Professor at the University of Minnesota. He received his Ph.D. in Physical Chemistry from the University of Toronto in 1991 and carried out postdoctoral research at Bayer A.G., Germany, the University of Texas, and Penn O



Perspective

State University. Professor Stein’s research interests are in the field of solid state chemistry, in particular porous materials, templating methods, and nanocomposites. He is the recipient of several awards, including a Merck Professorship, 3M Faculty Grant, a Dupont Young Professor Grant, an NSF CAREER Award, a McKnight Land-Grant Professorship, and a David & Lucile Packard Fellowship.

(27) Fan, J.; Boettcher, S. W.; Tsung, C.-K.; Shi, Q.; Schierhorn, M.; Stucky, G. D. Chem. Mater. 2008, 20, 909. (28) Tian, F.; Luo, Y.; Zhang, X. J. Chem. Phys. 2010, 133, 144701/1. (29) Xiang, H.; Shin, K.; Kim, T.; Moon, S. I.; McCarthy, T. J.; Russell, T. P. Macromolecules 2005, 38, 1055. (30) Zhang, X.; Chen, G.; Wang, W. J. Chem. Phys. 2007, 127, 034506/1. (31) Fasolka, M. J.; Mayes, A. M. Annu. Rev. Mater. Res. 2001, 31, 323. (32) Huinink, H. P.; Brokken-Zijp, J. C. M.; van Dijk, M. A.; Sevink, G. J. A. J. Chem. Phys. 2000, 112, 2452. (33) Li, H.-W.; Huck, W. T. S. Nano Lett. 2004, 4, 1633. (34) Richman, E. K.; Brezesinski, T.; Tolbert, S. H. Nat. Mater. 2008, 7, 712. (35) Rankin, S. E.; Malonski, A. P.; van Swol, F. Mater. Res. Soc. Symp. Proc. 2001, 636, DI2.1. (36) Cheng, J. Y.; Mayes, A. M.; Ross, C. A. Nat. Mater. 2004, 3, 823. (37) Bang, J.; Jeong, U.; Ryu, D. Y.; Russell, T. P.; Hawker, C. J. Adv. Mater. 2009, 21, 4769. (38) Kim, S. O.; Solak, H. H.; Stoykovich, M. P.; Ferrier, N. J.; de Pablo, J. J.; Nealey, P. F. Nature 2003, 424, 411. (39) Jorge, M.; Gomes, J. R. B.; Cordeiro, M. N. D. S.; Seaton, N. A. J. Am. Chem. Soc. 2007, 129, 15414. (40) Jorge, M.; Gomes, J. R. B.; Cordeiro, M. N. D. S.; Seaton, N. A. J. Phys. Chem. B 2009, 113, 708. (41) Pérez-Sánchez, G.; Gomes, J. R. B.; Jorge, M. Langmuir 2013, 29, 2387. (42) Gov, N.; Borukhov, I.; Goldfarb, D. Langmuir 2006, 22, 605. (43) Yang, H.; Coombs, N.; Ozin, G. A. J. Mater. Chem. 1998, 8, 1205. (44) Zhao, D.; Yang, P.; Melosh, N.; Feng, J.; Chmelka, B. F.; Stucky, G. D. Adv. Mater. 1998, 10, 1380. (45) Klotz, M.; Albouy, P.-A.; Ayral, A.; Menager, C.; Grosso, D.; Van der Lee, A.; Cabuil, V.; Babonneau, F.; Guizard, C. Chem. Mater. 2000, 12, 1721. (46) Alberius, P. C. A.; Frindell, K. L.; Hayward, R. C.; Kramer, E. J.; Stucky, G. D.; Chmelka, B. F. Chem. Mater. 2002, 14, 3284. (47) Firouzi, A.; Schaefer, D. J.; Tolbert, S. H.; Stucky, G. D.; Chmelka, B. F. J. Am. Chem. Soc. 1997, 119, 9466. (48) Yamauchi, Y.; Sawada, M.; Komatsu, M.; Sugiyama, A.; Osaka, T.; Hirota, N.; Sakka, Y.; Kuroda, K. Chem. Asian J. 2007, 2, 1505. (49) Yamauchi, Y.; Sawada, M.; Sugiyama, A.; Osaka, T.; Sakka, Y.; Kuroda, K. J. Mater. Chem. 2006, 16, 3693. (50) Yamauchi, Y.; Sawada, M.; Noma, T.; Ito, H.; Furumi, S.; Sakka, Y.; Kuroda, K. J. Mater. Chem. 2005, 15, 1137. (51) Walcarius, A.; Sibottier, E.; Etienne, M.; Ghanbaja, J. Nat. Mater. 2007, 6, 602. (52) Goux, A.; Etienne, M.; Aubert, E.; Lecomte, C.; Ghanbaja, J.; Walcarius, A. Chem. Mater. 2009, 21, 731. (53) Hillhouse, H. W.; van Egmond, J. W.; Tsapatsis, M. Chem. Mater. 2000, 12, 2888. (54) Ma, C.; Han, L.; Jiang, Z.; Huang, Z.; Feng, J.; Yao, Y.; Che, S. Chem. Mater. 2011, 23, 3583. (55) Brinker, C. J.; Dunphy, D. R. Curr. Opin. Colloid Interface Sci. 2006, 11, 126. (56) Hara, M.; Nagano, S.; Seki, T. J. Am. Chem. Soc. 2010, 132, 13654. (57) Blas, H.; Save, M.; Pasetto, P.; Boissière, C.; Sanchez, C.; Charleux, B. Langmuir 2008, 24, 13132. (58) Freer, E. M.; Krupp, L. E.; Hinsberg, W. D.; Rice, P. M.; Hedrick, J. L.; Cha, J. N.; MIller, R. D.; Kim, H.-C. Nano Lett. 2005, 5, 2014. (59) Dutreilh-Colas, M.; Yan, M.; Labrot, P.; Delorme, N.; Gibaud, A.; Bardeau, J.-F. Surf. Sci. 2008, 602, 829. (60) Fukumoto, H.; Nagano, S.; Kawatsuki, N.; Seki, T. Chem. Mater. 2006, 18, 1226. (61) Miyata, H. Microporous Mesoporous Mater. 2007, 101, 296. (62) Miyata, H.; Kuroda, K. Chem. Mater. 2000, 12, 49.

■

ACKNOWLEDGMENTS A.S. acknowledges support by the Department of Energy Office of Science (DE-SC0008662) for the preparation of this review. S.G.R. and N.D.P. thank the University of Minnesota Initiative for Renewable Energy and the Environment (IREE) and the National Science Foundation (DMR-0704312) for partial support. Portions of the work were carried out in the University of Minnesota Characterization Facility, which receives partial support from the NSF through the MRSEC, ERC, MRI, and NNIN programs.

■

REFERENCES

(1) Yanagisawa, T.; Shimizu, T.; Kuroda, K.; Kato, C. Bull. Chem. Soc. Jpn. 1990, 63, 988. (2) Beck, J. S.; Vartuli, J. C.; Roth, W. J.; Leonowicz, M. E.; Kresge, C. T.; Schmitt, K. D.; Chu, C. T.-W.; Olson, D. H.; Sheppard, E. W.; McCullen, S. B.; Higgins, J. B.; Schlenker, J. L. J. Am. Chem. Soc. 1992, 114, 10834. (3) Bagshaw, S. A.; Prouzet, E.; Pinnavaia, T. J. Science 1995, 269, 1242. (4) Zhao, D.; Feng, J.; Huo, Q.; Melosh, N.; Fredrickson, G. H.; Chmelka, B. F.; Stucky, G. D. Science 1998, 279, 548. (5) Hatton, B.; Landskron, K.; Whitnall, W.; Perovic, D.; Ozin, G. A. Acc. Chem. Res. 2005, 38, 305. (6) Wan, Y.; Zhao, D. Chem. Rev. 2007, 107, 2821. (7) Kimura, T.; Kuroda, K. Adv. Funct. Mater. 2009, 19, 511. (8) Qiu, H.; Che, S. Chem. Soc. Rev. 2011, 40, 1259. (9) Kruk, M. Acc. Chem. Res. 2012, 45, 1678. (10) Ma, T. Y.; Liu, L.; Yuan, Z.-Y. Chem. Soc. Rev. 2013, 42, 3977. (11) Sanchez, C.; Boissière, C.; Grosso, D.; Laberty, C.; Nicole, L. Chem. Mater. 2008, 20, 682. (12) Gu, D.; Schüth, F. Chem. Soc. Rev. 2013, DOI: 10.1039/ C3CS60155B, (in press). (13) Vu, A.; Qian, Y.; Stein, A. Adv. Energy Mater. 2012, 2, 1056. (14) Li, Y.; Fu, Z.-Y.; Su, B. L. Adv. Funct. Mater. 2012, 22, 4634. (15) Su, B.-L.; Sanchez, C.; Yang, X.-Y. Hierarchically Structured Porous Materials: from Nanoscience to Catalysis, Biomedicine, Optics and Energy; Wiley-VCH: Weinheim, 2012. (16) Petkovich, N. D.; Stein, A. In Hierarchically Structured Porous Materials: from Nanoscience to Catalysis, Biomedicine, Optics and Energy; Su, B. L., Sanchez, C., Eds.; Wiley-VCH: Weinheim, Germany, 2012; p 55. (17) Petkovich, N. D.; Stein, A. Chem. Soc. Rev. 2013, 3721. (18) Innocenzi, P.; Malfatti, L.; Soler-Illia, G. J. A. A. Chem. Mater. 2011, 23, 2501. (19) Stein, A.; Li, F.; Denny, N. R. Chem. Mater. 2008, 20, 649. (20) Darling, S. B. Prog. Polym. Sci. 2007, 32, 1152. (21) Platschek, B.; Petkov, N.; Bein, T. Angew. Chem., Int. Ed. 2006, 45, 1134. (22) Xiang, H.; Shin, K.; KIm, T.; Moon, S.; McCarthy, T. J.; Russell, T. P. J. Polym. Sci., Part B: Polym. Phys. 2005, 43, 3377. (23) Schulz, M. F.; Bates, F. S. In Physical Properties of Polymers Handbook; Mark, J. E., Ed.; AIP Press: Woodbury, NY, 1996; p 427. (24) Cheng, J. Y.; Ross, C. A.; Thomas, E. L.; Smith, H. I.; Vansco, G. J. Appl. Phys. Lett. 2002, 81, 3657. (25) Cheng, J. Y.; Ross, C. A.; Thomas, E. L.; Smith, H. I.; Vansco, G. J. Adv. Mater. 2003, 15, 1599. (26) Shin, K.; Xiang, H.; Moon, S. I.; Kim, T.; McCarthy, T. J.; Russell, T. P. Science 2004, 306, 76. P



Perspective

(63) Miyata, H.; Noma, T.; Watanabe, M.; Kuroda, K. Chem. Mater. 2002, 14, 766. (64) Bai, L.; Li, Y.; Chen, B. Mol. Simul. 2012, 38, 554. (65) Siperstein, F. R.; Gubbins, K. E. Langmuir 2003, 19, 2049. (66) Ferreiro-Rangel, C. A.; Lozinska, M. M.; Wright, P. A.; Seaton, N. A.; Düren, T. J. Phys. Chem. C 2012, 116, 20966. (67) Li, F.; Stein, A. Langmuir 2012, 28, 7484. (68) Wu, C.-W.; Ohsuna, T.; Edura, T.; Kuroda, K. Angew. Chem., Int. Ed. 2007, 46, 5364. (69) Chen, B.-C.; Lin, H.-P.; Chao, M.-C.; Mou, C.-Y.; Tang, C.-Y. Adv. Mater. 2004, 16, 1657. (70) Yeh, Y.-Q.; Lin, H.-P.; Tang, C.-Y.; Mou, C.-Y. J. Colloid Interface Sci. 2011, 362, 354. (71) Koganti, V. R.; Rankin, S. E. J. Phys. Chem. B 2005, 109, 3279. (72) Koganti, V. R.; Dunphy, D.; Gowrishankaar, V.; McGehee, M. D.; Li, X.; Wang, J.; Rankin, S. E. Nano Lett. 2006, 6, 2567. (73) Suzuki, N.; Kimura, T.; Yamauchi, Y. J. Mater. Chem. 2010, 20, 5294. (74) Yamauchi, Y.; Suzuki, N.; Kimura, T. Chem. Commun. 2009, 5689. (75) Ku, A. Y.; Taylor, S. T.; Heward, W. J.; Denault, L.; Loureiro, S. M. Microporous Mesoporous Mater. 2006, 88, 214. (76) Schuster, J.; Köhn, R.; Döblinger, M.; Keilbach, A.; Amenitsch, H.; Bein, T. J. Am. Chem. Soc. 2012, 134, 11136. (77) Platschek, B.; Keilbach, A.; Bein, T. Adv. Mater. 2011, 23, 2395. (78) Yang, Z.; Niu, Z.; Cao, X.; Yang, Z.; Lu, Y.; Hu, Z.; Han, C. C. Angew. Chem., Int. Ed. 2003, 42, 4201. (79) Wu, Y.; Cheng, G.; Katsov, K.; Sides, S. W.; Wang, J.; Tang, J.; Fredrickson, G. H.; Moskovits, M.; Stucky, G. D. Nat. Mater. 2004, 3, 816. (80) Lu, Q.; Gao, F.; Komarneni, S.; Mallouk, T. E. J. Am. Chem. Soc. 2004, 8650−8651. (81) Kim, E.-M.; Jung, J.-S.; Chae, W.-S. Chem. Commun. 2010, 46, 1760. (82) Platschek, B.; Köhn, R.; Döblinger, M.; Bein, T. ChemPhysChem 2008, 9, 2059. (83) Keilbach, A.; Döblinger, M.; Köhn, R.; Amenitsch, H.; Bein, T. Chem.Eur. J. 2009, 15, 6645. (84) O’Callaghan, J. M.; Petkov, N.; Copley, M. P.; Arnold, D. C.; Morris, M. A.; Amenitsch, H.; Holmes, J. D. Microporous Mesoporous Mater. 2010, 130, 203. (85) Jin, K.; Yao, B.; Wang, N. Chem. Phys. Lett. 2005, 409, 172. (86) Yao, B.; Fleming, D.; Morris, M. A.; Lawrence, S. E. Chem. Mater. 2004, 16, 4851. (87) Yamaguchi, A.; Uejo, F.; Yoda, T.; Uchida, T.; Tanamura, Y.; Yamashita, T.; Teramae, N. Nat. Mater. 2004, 3, 337. (88) Platschek, B.; Köhn, R.; Döblinger, M.; Bein, T. Langmuir 2008, 24, 5018. (89) Keller, A.; Kimayer, S.; Segal-Peretz, T.; Frey, G. Langmuir 2012, 28, 1506. (90) Zhang, A.; Hou, K.; Gu, L.; Dai, C.; Liu, M.; Song, C.; Guo, X. Chem. Mater. 2012, 24, 1005. (91) Platschek, B.; Petkov, N.; Himsl, D.; Zimdars, S.; Li, Z.; Köhn, R.; Bein, T. J. Am. Chem. Soc. 2008, 130, 17362. (92) Yamauchi, Y.; Nagaura, T.; Ishikawa, A.; Chikyow, T.; Inoue, S. J. Am. Chem. Soc. 2008, 130, 10165. (93) Yamauchi, Y.; Nagaura, T.; Inoue, S. Chem. Asian J. 2009, 4, 1059. (94) Stein, A.; Wilson, B. E.; Rudisill, S. G. Chem. Soc. Rev. 2013, 42, 2763. (95) Velev, O. D.; Jede, T. A.; Lobo, R. F.; Lenhoff, A. M. Nature 1997, 389, 447. (96) Holland, B. T.; Blanford, C. F.; Stein, A. Science 1998, 281, 538. (97) Wijnhoven, J. E. G. J.; Vos, W. L. Science 1998, 281, 802. (98) Zakhidov, A. A.; Baughman, R. H.; Iqbal, Z.; Cui, C.; Khayrullin, I.; Dantas, S. O.; Marti, J.; Ralchenko, V. G. Science 1998, 282, 897. (99) Park, S. H.; Xia, Y. Chem. Mater. 1998, 10, 1745. (100) Yang, P.; Deng, T.; Zhao, D.; Feng, P.; Pine, D.; Chmelka, B. F.; Whitesides, G. M.; Stucky, G. D. Science 1998, 282, 2244.

(101) Holland, B. T.; Blanford, C. F.; Do, T.; Stein, A. Chem. Mater. 1999, 11, 795. (102) Antonietti, M.; Berton, B.; Göltner, C.; Hentze, H. P. Adv. Mater. 1998, 10, 154. (103) Falcaro, P.; Malfatti, L.; Kidchob, T.; Giannini, G.; Falqui, A.; Casula, M. F.; Amenitsch, H.; Marmiroli, B.; Grenci, G.; Innocenzi, P. Chem. Mater. 2009, 21, 2055. (104) Zhou, Y.; Antonietti, M. Chem. Commun. 2003, 2564. (105) Zhou, Y.; Antonietti, M. Adv. Mater. 2003, 15, 1452. (106) Fujita, S.; Nakano, H.; Ishii, M.; Nakamura, H.; Inagaki, S. Microporous Mesoporous Mater. 2006, 96, 205. (107) Sel, O.; Sallard, S.; Brezesinksi, T.; Rathousky, J.; Dunphy, D. R.; Collord, A.; Smarsly, B. M. Adv. Funct. Mater. 2007, 17, 3241. (108) Kuang, D.; Brezesinski, T.; Smarsly, B. J. Am. Chem. Soc. 2004, 126, 10534. (109) Sen, T.; Tiddy, J. T.; Casci, J. L.; Anderson, M. W. Chem. Mater. 2004, 16, 2044. (110) Sel, O.; Kuang, D.; Thommes, M.; Smarsly, B. Langmuir 2006, 22, 2311. (111) Li, F.; Wang, Z.; Ergang, N. S.; Fyfe, C. A.; Stein, A. Langmuir 2007, 23, 3996. (112) Loiola, A. R.; da Silva, L. R. D.; Cubillas, P.; Anderson, M. W. J. Mater. Chem. 2008, 18, 4985. (113) Sun, Z.; Deng, Y.; Wei, J.; Gu, D.; Tu, B.; Zhao, D. Chem. Mater. 2011, 23, 2176. (114) Kamperman, M.; Burns, A.; Weissgraeber, R.; van Vegten, N.; Warren, S. C.; Gruner, S. M.; Baiker, A.; Wiesner, U. Nano Lett. 2009, 9, 2756. (115) Sakurai, M.; Shimojima, A.; Heishi, M.; Kuroda, K. Langmuir 2007, 23, 10788. (116) Etienne, M.; Sallard, S.; Schröder, M.; Guillemin, Y.; Mascotto, S.; Smarsly, B. M.; Walcarius, A. Chem. Mater. 2010, 22, 3426. (117) Bian, S.-W.; Zhang, Y.-L.; Li, H.-L.; Yu, Y.; Song, Y.-L.; Song, W.-G. Microporous Mesoporous Mater. 2010, 131, 289. (118) Deng, Y.; Liu, C.; Yu, T.; Liu, F.; Zhang, F.; Wan, Y.; Zhang, L.; Wang, C.; Tu, B.; Webley, P. A.; Wang, H.; Zhao, D. Chem. Mater. 2007, 19, 3271. (119) Wang, Z.; Kiesel, E. R.; Stein, A. J. Mater. Chem. 2008, 18, 2194. (120) Wang, Z.; Stein, A. Chem. Mater. 2008, 20, 1029. (121) Tanaka, S.; Nishiyama, N.; Egashira, Y.; Ueyama, K. Chem. Commun. 2005, 2125. (122) Meng, Y.; Gu, D.; Zhang, F.; Shi, Y.; Yang, H.; Li, Z.; Yu, C.; Tu, B.; Zhao, D. Angew. Chem., Int. Ed. 2005, 44, 7053. (123) Liang, C.; Dai, S. J. Am. Chem. Soc. 2006, 128, 5316. (124) Attard, G. S.; Glyde, J. C.; Goltner, C. G. Nature 1995, 378, 366. (125) Yang, S. M.; Coombs, N.; Ozin, G. A. Adv. Mater. 2000, 12, 1940. (126) Arsenault, A. C.; Rider, D. A.; Tetreault, N.; Chen, J. I.-L.; Coombs, N.; Ozin, G. A.; Manners, I. J. Am. Chem. Soc. 2005, 127, 9954. (127) Rider, D. A.; Chen, J. I.-L.; Eloi, J.-C.; Arsenault, A. C.; Russell, T. P.; Ozin, G. A.; Manners, I. Macromolecules 2008, 41, 2250. (128) Mandlmeier, B.; Szeifert, J. M.; Fattakhova-Rohlfing, D.; Amenitsch, H.; Bein, T. J. Am. Chem. Soc. 2011, 133, 17274. (129) Zhou, D.-D.; Liu, H.-J.; Wang, Y.-G.; Wang, C.-X.; Xia, Y.-Y. J. Mater. Chem. 2012, 22, 1937. (130) Nakanishi, K.; Soga, N. J. Am. Ceram. Soc. 1991, 74, 2518. (131) Nakanishi, K. J. Porous Mater. 1997, 4, 67. (132) Nakanishi, K.; Kobayashi, Y.; Amatani, T.; Hirao, K.; Kodaira, T. Chem. Mater. 2004, 16, 3652. (133) Nakanishi, K.; Tanaka, N. Acc. Chem. Res. 2007, 40, 863. (134) Rudisill, S. G.; Hein, N. M.; Terzic, D.; Stein, A. Chem. Mater. 2013, 25, 745. (135) Pechini, M. P. U.S. Patent 3,330,697, 1967. (136) Vu, A.; Stein, A. Chem. Mater. 2011, 23, 3237. (137) Petkovich, N. D.; Stein, A. Unpublished results. (138) Vu, A.; Stein, A. J. Power Sources 2014, 245, 48−58. Q



Perspective

(139) Rauda, I. E.; Buonsanti, R.; Saldarriaga-Lopez, L. C.; Benjauthrit, K.; Schelhas, L. T.; Stefik, M.; Augustyn, V.; Ko, J.; Dunn, B.; Wiesner, U.; Milliron, D. J.; Tolbert, S. H. ACS Nano 2012, 6, 6386. (140) Gommes, C. J.; Friedrich, H.; Wolters, M.; de Jongh, P. E.; de Jong, K. P. Chem. Mater. 2009, 21, 1311. (141) Ward, E. P. W.; Yates, T. J. V.; Fernandez, J.-J.; Vaughan, D. E. W.; Midgley, P. A. J. Phys. Chem. C 2007, 111, 11501. (142) Yates, T. J. V.; Thomas, J. M.; Fernandez, J.-J.; Terasaki, O.; Ryoo, R.; Midgley, P. A. Chem. Phys. Lett. 2006, 418, 540. (143) Ersen, O.; Parmentier, J.; Solovyov, L. A.; Drillon, M.; PhamHuu, C.; Werckmann, J.; Schultz, P. J. Am. Chem. Soc. 2008, 130, 16800. (144) Monnier, A.; Schüth, F.; Huo, Q.; Kumar, D.; Margolese, D.; Maxwell, R. S.; Stucky, G. D.; Krishnamurty, M.; Petroff, P.; Firouzi, A.; Janicke, M.; Chmelka, B. F. Science 1993, 261, 1299. (145) Huo, Q.; Margolese, D. I.; Ciesla, U.; Demuth, D. G.; Feng, P.; Gier, T. E.; Sieger, P.; Firouzi, A.; Chmelka, B. F.; Schüth, F.; Stucky, G. D. Chem. Mater. 1994, 6, 1176. (146) Bhattacharya, A.; Mahanti, S. D. J. Phys.: Condens. Matter 2001, 13, 1413. (147) Ergang, N. S.; Fierke, M. A.; Wang, Z.; Smyrl, W. H.; Stein, A. J. Electrochem. Soc. 2007, 154, A1135. (148) Zhang, H.; Yu, X.; Braun, P. V. Nat. Nanotechnol. 2011, 6, 277. (149) Cho, Y. K.; Wartena, R.; Tobias, S. M.; Chiang, Y.-M. Adv. Funct. Mater. 2007, 17, 379.

R


Perspective on the Influence of Interactions Between Hard and Soft

Recommend Documents