Thin Polymer Films with Continuous Vertically Aligned 1 nm Pores

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Thin Polymer Films with Continuous Vertically Aligned 1-nm Pores Fabricated by Soft Confinement Xunda Feng, Siamak Nejati, Matthew G. Cowan, Marissa E. Tousley, Brian R. Wiesenauer , Richard D. Noble, Menachem Elimelech, Douglas L. Gin, and Chinedum O. Osuji ACS Nano, Just Accepted Manuscript • DOI: 10.1021/acsnano.5b06130 • Publication Date (Web): 03 Dec 2015 Downloaded from http://pubs.acs.org on December 4, 2015

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Thin Polymer Films with Continuous Vertically Aligned 1-nm Pores Fabricated by Soft Confinement Xunda Feng,† Siamak Nejati, † Matthew G. Cowan,‡,§ Marissa E. Tousley,† Brian R. Wiesenauer,‡ Richard D. Noble,§ Menachem Elimelech,† Douglas L. Gin ‡,§ and Chinedum O. Osuji†,*

†

Department of Chemical and Environmental Engineering, Yale University, New Haven, CT

06511, USA ‡

Department of Chemistry and Biochemistry, University of Colorado, Boulder, CO 80309,

USA §

Department of Chemical and Biological Engineering, University of Colorado, Boulder, CO

80309, USA

*

Corresponding author: [email protected]

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Abstract Membrane separations are critically important in areas ranging from health care and analytical chemistry, to bioprocessing and water purification. An ideal nanoporous membrane would consist of a thin film with physically continuous and vertically aligned nanopores, with a narrow distribution of pore sizes. However the current state of the art departs considerably from this ideal and is beset by intrinsic tradeoffs between permeability and selectivity. We demonstrate an effective and scalable method to fabricate polymer films with ideal membrane morphologies consisting of sub-micron thickness films with physically continuous and vertically aligned 1 nm pores. The approach is based on soft confinement to control the orientation of a cross-linkable mesophase in which the pores are produced by self-assembly. The scalability, exceptional ease of fabrication, and potential to create a new class of nanofiltration membranes stand out as compelling aspects.

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A long-standing goal in materials science is to provide highly ordered or periodic nanostructures with useful properties over large length scales or technologically relevant dimensions. Bottom-up approaches involve the self-assembly of atomic, molecular and colloidal building blocks as a promising way to achieve this goal at potentially lower cost or higher throughput than top-down strategies1-3 The translation of these ideas to polymeric materials provides potentially low cost pathways to highly compelling applications, such as precisely tailored nanoporous membranes,4,

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photonic band gap materials,6 and high

resolution lithography using self-assembled structures as pattern transfer masks.7 The performance of these ordered or periodic nanostructured polymeric materials often requires thin films to possess a high degree of long-range order (i.e., positional order) and alignment (i.e., orientational order) of the morphology – for the present purposes, the terms orientation and alignment will be used interchangeably to refer simply to the orientational, rather than positional order in the system. In the case of nanoporous membranes, the ideal system is one in which all pores possess the same diameter and are aligned or oriented parallel to the macroscopic transport direction (i.e., in the “thickness”, “through-plane”, or “vertical” direction).8 There have been continuous efforts in developing zeolites, carbon nanotubes, and metal–organic framework membranes with aligned nanopores of molecular size. However, the current state of the art in the area of polymer membranes, departs considerably from this ideal. Ultrafiltration (UF) and nanofiltration (NF) are membrane based separation processes of high commercial importance in varied fields such as food processing, antibiotics manufacturing, and water purification. UF membranes (pore sizes of ~10-100 nm) produced by conventional techniques such as phase inversion have highly tortuous interconnected 3


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pores, and a large variation of pore diameters. The broad pore diameter distribution severely limits membrane selectivity as evident from their molecular-weight cutoff characteristics.9-11 Additionally, pore tortuosity negatively impact membrane permeability as the physical transport distance significantly exceeds the thickness of the membrane.11,

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While UF

operates based on size-exclusion, separation in NF (effective pore sizes ~1-10 nm) is governed to a large extent by solution-diffusion through a dense permselective layer, typically polyamide.11 As a result, there is an inherent tradeoff between permeability and selectivity that sets an upper bound on the performance of NF separations. This limiting phenomenon has been also present in gas separation membranes.

In principle, the aforementioned permeability and selectivity issues can be circumvented using self-assembled materials such as block copolymers (BCPs) or small-molecule liquid crystals (LCs) that feature nanostructures with thermodynamically defined characteristic dimensions that are therefore narrowly distributed.4, 13-15 Leveraging self-assembled materials to fabricate ideal membranes requires both physical continuity and vertical orientation of nanostructures over large areas in thin films. As one might expect, such morphologies in general do not result spontaneously by self-assembly of these systems.16 A concerted effort is therefore required, for example through the use of interfacial engineering17 or external fields,18-20 to control orientation. Although magnetic and electric field alignment can be effective in directing the formation of ordered polymer nanostructures, these approaches are best suited to morphological control in the bulk, i.e., where surface forces do not play a prominent and potentially confounding role, as the field strengths required to overcome 4


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surface induced morphologies are impractically large in most cases. Conversely, physical continuity is determined principally by the propensity of the material to display defects such as dislocations, which represent ‘dead-ends’ in terms of transport.

Our current concern is the development of polymer thin films in which self-assembly provides access to uniform pore sizes in the 1-nm regime. This length scale is of considerable interest in water purification by NF as it permits effective removal of multivalent salts and small molecule solutes by size-exclusion, rather than by the solution-diffusion mechanism. Furthermore, realization of physical continuity and vertical orientation of the uniform 1-nm pores over large areas in thin films with thickness in the sub-µm range has been a non-trivial challenge. Until now, previously reported sub-µm polymer films possessing vertically aligned nanopore structures have been fabricated by BPCs which places a lower bound of ca. 5 nm on the pore diameter.8, 21

Here, we report a strategy to achieve this goal by subjecting sub-µm-thick films of a cross-linkable self-assembled liquid crystalline material to soft confinement using an elastomeric pad of poly(dimethylsiloxane) or PDMS, as schematically illustrated in Figure 1. This method provides a scalable solution for the fabrication of thin polymer films possessing vertically oriented, 1-nm-diameter pores using simple and readily accessible tools. High-resolution transmission electron microscopy (TEM) images provide unambiguous confirmation that the pores are physically continuous with persistent vertical orientation through the entire film thickness. 5


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Figure 1. Schematic illustration of soft confinement-assisted fabrication of polymer thin films with vertically aligned nanopores. (a) Molecular structure of the wedge-shaped amphiphilic LC monomer Na-GA3C11. (b) Self-assembled hexagonal columnar (Colh) phase possessing 1-nm-diameter hydrophilic cylindrical pores. (c) Sub-µm thick LC film is obtained by casting a dilute Na-GA3C11/THF solution onto a silicon substrate and allowing the solvent to evaporate. The as-cast thin film contains supramolecular nanoporous columns lying parallel to the film plane. (d) When a soft PDMS pad is imposed onto this film followed by thermal annealing, the columns adopt a vertical orientation. (e) Photo-cross-linking of the aligned columns results in a sub-µm-thick polymer film with vertically oriented nanopores that can be easily detached from the substrate.

Results and Discussion

Figure 1a shows the molecular structure of the employed amphiphilic LC monomer (Na-GA3C11), the synthesis and characterization of which have been previously reported.22 This type of wedge-shaped amphiphilic molecule possessing a large hydrophobic body and a small hydrophilic head tends to form supramolecular hexagonal columnar (Colh) LC phases with closely-packed, ordered hydrophilic nanochannels.23-27 In its neat state, Na-GA3C11 6


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self-assembles into the Colh phase at room temperature, as confirmed by small angle X-ray scattering (SAXS) studies that show the d(100)-spacing of 3.6 nm and the characteristic diffraction peak location ratios of 1:√3:√4 (Figure S1, Supporting Information). The columnar phase transitions into an isotropic state above 64.9 °C.22 The introduction of reactive acrylate groups at the periphery of Na-GA3C11 enables structural lock-in of the hexagonal columnar order by photo-generated radical cross-linking to afford a mechanically and chemically robust polymer. High-resolution TEM visualization of a microtomed specimen from the magnetically aligned and subsequently photo-cross-linked bulk sample reveals the hexagonally-packed nanopores with an as-estimated diameter of 1 nm (Figure S2, Supporting Information), as previously reported.22 The electron density contrast in the image is provided by chemical staining of the aromatic periphery of the pores and so what is visualized is the outline of the pore, rather than the pore itself.

A thin LC monomer film was prepared by casting a dilute solution of Na-GA3C11 containing a small amount of radical photo-initiator (0.5 wt %) on a silicon substrate and allowing the solvent to evaporate. The unconfined film was then heated into the isotropic phase (75 °C) before being cooled back to room temperature at a rate of 0.1 °C/min. A polymer film (~350 nm thick) was obtained by subsequent photo-cross-linking. Figure 2 shows a direct TEM image of a plan (or “top-down”) view of the cross-linked film. The columnar axes are oriented parallel to the plane of the film with the azimuthal angle being arbitrary, as evidenced by two coexistent grains showing different planar columnar orientations with an observable grain boundary. The Fourier-transform pattern (inset) 7


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exhibiting four-fold contributions further demonstrates the degenerate planar nature of the orientation.

Figure 2. Plan or “top-down” view TEM image of an open polymer film. The film was produced by casting a Na-GA3C11/THF solution onto a glass substrate, followed by thermal treatment and subsequent radical photo-cross-linking. The image shows degenerate planar orientation of the columnar nanochannels. Fourier transform of the image data (inset) indicates the different orientations of two coexistent grains with an observable grain boundary. The nanopores adopted a strikingly different orientation when a soft PDMS elastomer pad was placed in contact with the sample before thermal annealing and cross-linking. Figure 3a shows a plan view TEM image of the thin film prepared by soft confinement. A single-crystal-like array of hexagonally packed nanopores was observed. The Fourier transform (inset) of this image displays sharp six-fold symmetric contributions up to the third order, reflecting the high degree of structural order in the sample. A cross-sectional TEM image (Figure 3b) clearly shows that the nanopores are vertically aligned and physically 8


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continuous through the thickness of the film. The Fourier transform of the TEM image (inset) displays the corresponding two-fold symmetry expected for a cross-sectional view of vertically aligned structures.

Figure 3. TEM images of a cross-linked thin film of Na-GA3C11 produced by soft confinement. The film with a thickness of ~350 nm shows highly-ordered, vertical nanochannels. (a) Direct image of the thin polymer film displayed a single-crystal-like hexagonal pattern. (b) Cross-sectional image of the specimen microtomed parallel to the film normal, demonstrating nanochannel vertical orientation and persistence through the film thickness. Fourier transforms of the two images (insets) show 6- and 2-fold symmetry, respectively, indicating ordered hexagonal morphology. High-magnification images with artificial colors and 3-D models (insets) highlight the two imaging directions. 9


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Figure 4. Room-temperature POM images of thermally annealed Na-GA3C11 films. Films with different thickness and under different surface confinement conditions were prepared. The insets are conoscopy images obtained by a 40× lens. The corresponding 3-D models show the orientations of the superamolecular columns in the films: (a) an 80-µm-thick film sandwiched by two glass slides; (b) a 28-µm-thick film sandwiched by two glass slides; (c) a 28-µm-thick film prepared on a glass slide with one surface exposed to the air; and (d) a 28-µm-thick film on a glass slide covered by a PDMS elastomer pad. Scale bars: all 200 µm.

The different orientations of nanopores found in open vs. confined films originate from different surface anchoring of the columnar structures at the free air surface relative to the PDMS interface. Polarized optical microscopy (POM) and conoscopy investigations were performed to evaluate the role of confinement over larger sample areas and film thicknesses than could be suitably examined in TEM. Figure 4a shows POM and conoscopic images (inset) of an 80-µm-thick film sandwiched by two flat glass slides after thermal treatment. The observed birefringent texture under POM as well as the poorly defined features in the conoscopic image indicate that the optical axes of the domains (i.e., the columnar axes) are not uniformly oriented along the normal to the film surface. That is, there are significant deviations of the columnar axes from the vertical direction in the film. However, when the 10


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film thickness was reduced to 28 µm, the sample exhibited optical extinction under POM and a characteristic Maltese cross in conoscopic imaging, respectively (Figure 4b), which are consistent with the uniform vertical orientation of the nanopores observed in TEM. This behavior indicates that the glass interface induced homeotropic anchoring to the columnar structure and that this effect is more pronounced in thinner samples.

In contrast to the sandwiched sample, an open film of the same thickness (28 µm) with one surface exposed to air displayed randomly oriented nanopores, as evidenced by a birefringent POM texture and poorly defined conoscopic image (Figure 4c). This observation suggests that the free surface of the film induced non-homeotropic anchoring of the system, most likely resulting in a degenerate planar or homogeneous anchoring condition. Air interfaces under ambient conditions are known to favor contact with hydrophobic moieties. We therefore surmise that the anchoring behavior in the 28-µm-thick open sample results from energy minimization at the free air interface by preferentially displaying the hydrophobic alkyl tails of the mesogens, leaving the polar ionic head groups recessed. Thus, the open film was subject to antagonistic or asymmetric boundary conditions as schematically illustrated by the 3-D model in Figure 4c.

When a smooth, soft PDMS pad was placed onto the 28-µm-thick open film, followed by thermal annealing, vertically aligned nanopores were obtained, as confirmed by a dark POM appearance and a characteristic Maltese cross from conoscopy (Figure 4d). This observation indicates that both the glass and the PDMS interfaces induce the same homeotropic anchoring 11


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condition in the system.

The influence of the cooling rate on the quality of vertical orientation of nanopores was also investigated by POM. We found that slow cooling rates favored uniformity of the vertical nanopore orientation within the whole LC film and a typical cooling rate down to 0.1-0.2 °C/min was required for a 28-µm-thick LC film covered by a PDMS stamp (see Figure S3, Supporting Information).

SAXS studies on the polymer films were performed to provide complementary and more quantitative data regarding the orientation of the nanopores. Figure 5a depicts the SAXS experiment, in which the polymer films were stacked to provide sufficiently thick samples for the measurement. Figure 5b shows the 2-D scattering patterns of films of different thicknesses prepared by thermal treatment under confinement between two glass slides. Anisotropic scattering with intensity concentrated along the equatorial line (i.e. along the “left-to-right” direction) confirm the TEM observation that the glass surface produces vertically aligned nanopores. The 1-D integrated data with diffraction peak location ratios of 1:√3:√4 shown in Figure 5c confirms the hexagonal symmetry of the nanostructure in the samples as expected for the Colh phase.

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Figure 5. SAXS measurements on cross-linked polymer films. (a) Schematic illustration of SAXS measurements (b) 2-D SAXS patterns of polymer films of Na-GA3C11 prepared in a glass-slide-sandwiched geometry with thicknesses of 80, 28, and 7 µm, respectively. (c) 1-D integral SAXS data of the polymerized sample demonstrating the hexagonal morphology. The primary reflection peak at q* = 0.179 Å-1 corresponds to the Bragg spacing d100 = 3.5 nm. (d) Azimuthal dependence of scattering intensity of the 2-D SAXS patterns in (b). (e) 2-D SAXS patterns of polymer films of Na-GA3C11 with thicknesses of 28 and 9 µm, respectively, prepared in soft confinement. (f) 2-D SAXS patterns of polymer films of Na-GA3C11 with a thickness of 28 µm prepared in an open-to-air geometry. (g) Room-temperature ionic conductivity data (at 100% relative humidty) measured for cross-linked films (~20 µm thick) of Na-GA3C11 prepared in sandwiched and open-to-air geometries, respectively. 13


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Figure 5d shows the azimuthal intensity distributions for the primary Bragg peak at different film thickness and Gaussian fits of such data. The alignment quality improved on decreasing the film thickness, as characterized by a reduction of the FWHM (full width at half maximum) of the Bragg peak. This thickness dependence is consistent with the POM observations in which the thicker films studied (80 µm) were birefringent due to the presence of columnar axes in the optical plane while thin samples (28 µm) were not, due to their columnar axes being aligned orthogonal to the optical plane. The azimuthal intensity distribution is exceptionally narrow at a thickness of 7 µm, with FWHM=10.7°. Using a Gaussian approximation for the azimuthal intensity distribution, this FWHM value yields an orientational order parameter S of 0.98, where S = 1 corresponds to a perfectly oriented system and S = 0 corresponds to a completely random one.28 The 2-D SAXS patterns (Figure 5e) of the polymer films prepared by soft confinement also display such a dependence of alignment quality on film thickness. The azimuthal intensity distribution at a thickness of 9 µm is much narrower than that of 28 µm with FWHM value of 23.9 vs 13.4°, respectively. Figure 5f shows 2-D SAXS of a 28-µm-thick film prepared with antagonistic boundary conditions in which one interface is exposed to air. The measured scattering is indicative of a morphology with mixed orientations of the columnar structures. Anisotropic scattering with sharp intensity along the meridian suggests the presence of structures with in-plane orientations in the film. The detectable scattering at other azimuthal positions may arise due to degeneracy of parallel orientation at the free surface and distortion of the columnar director under the antagonistic boundary conditions. 14


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The influence of film thickness on mesophase alignment can be readily understood in the context of the finite elasticity of the medium, which specifies the distance over which the memory of a given orientation will decay under thermal forces. That is, the length scale beyond which the elastic energy due to a non-uniform orientation of the LC director is less than or equal to kBT. We can think of this length scale in a dimensionally consistent manner as a persistence length λ defined by the ratio of an effective bending rigidity of the mesophase Keff and kBT, where the effective bending rigidity originates from the elasticity of the mesophase. The quality of the alignment then decays exponentially with distance z from the interface, as captured by the tilt away from the boundary condition, cos ߠ = ݁ ି௭/ఒ . The discernable differences between the 14 and 4.5 µm half-thickness films indicate that the persistence length is similar in magnitude to these dimensions. Therefore, it is reasonable to expect that when the film thickness is reduced to the sub-µm range, highly persistent alignment of nanopores should be achieved. The TEM observations in Figure 3b for a film of 350 nm confirm the validity of this reasoning.

Room-temperature ion conductivity measurements (at 100% relative humidity) on cross-linked polymer films of Na-GA3C11 prepared in sandwiched and open-to-air geometries illustrate the dramatic differences in transport properties associated with differences in alignment of the nanopores (Figure 5g). The sandwiched films show a remarkable 44-fold increase in their conductivity relative to the open films. The enhancement of conductivity in the aligned films observed here is more pronounced than that reported in 15


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aligned block copolymer systems. This difference may be due to the greater persistence of the aligned micro-domains along the ionic transport direction of the present system relative to that seen in block polymers in general.29, 30 The substantial improvement in the conductivity of the system in the sandwiched state highlights the potential utility of the current approach in producing membranes with attractive properties for functional applications.

TEM was employed to obtain a direct visualization of the surface structure and thereby provide insight regarding the orientation and physical continuity of the nanopores at the confining surfaces. The thickness of the investigated unconfined polymer film was about 400 µm, which we expect is sufficient to decouple the surface morphologies and hence exclude the possibility of significant director distortion. Figure 6a shows a cross-sectional TEM image of the nanostructure at the polymer/glass interface. It is clear that the nanopores are perpendicular to the glass interface and that the vertical orientation extends from the surface towards the bulk interior of the 400-µm-thick sample. Accordingly, the Fourier transform of the TEM image (inset) displays two-fold symmetry. Step-imaging at lower magnifications shows that the physical continuity and vertical orientation of the structure persists for a distance of approximately 5 to 10 µm (Figure S4, Supporting Information). This finding is in rough quantitative agreement with both the SAXS and POM data. TEM imaging of the polymer/air interface shows homogeneous anchoring, as distinguished by the occurrence of several areas in which circular hexagonally packed features are clearly visible, consistent with end-on views of the hexagonally packed columnar structures (Figure 6b). The homogeneous anchoring produces a degenerate situation as evident from the various 16


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projections of the parallel-aligned nanopores that are visible in addition to the hexagonally-packed end-on views (Figure S5, Supporting Information).

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Figure 6. Cross-sectional TEM images of surface morphology of the bulk polymer. The nanopores are oriented (a) vertically at the glass interface and (b) parallel to the air interface. On the basis of the data obtained from POM, SAXS, and TEM studies, it is clear that both the glass and PDMS surfaces induce vertical orientation of the nanopores of Na-GA3C11 but the free air interface results in degenerate planar alignment. Films exposed to air were therefore subjected to antagonistic boundary conditions that induce distortion of the structure. In sufficiently thin films, one expects uniform orientation of the nanopores throughout the film because distortion of the LC director on sub-µm thickness length scales in such a film would be precluded by the associated elastic energies.31 Assuming equilibrium conditions, the particular orientation observed would be determined by the surface that has stronger anchoring strength. On this assumption, the TEM image of the open sub-µm thin film (Figure 2) indicates that the free air surface was dominant in controlling the orientation of the nanopores, as a degenerate planar anchoring morphology was observed. It is apparent that the samples conformally contact both PDMS and glass surfaces such that air pockets are not formed, and exposure to air is therefore avoided. In the sandwiched arrangement, both surfaces of the film induce homeotropic anchoring leading to vertical orientation as desired.

The six-fold pattern of the FFT in Figure 3a is a reflection only of the existence of large grains, and not of the formation of a true macroscopic single crystal arrangement of the columns. The confinement of the system controls the orientation of the long axis of the columns, but does not place any constraints on the orientation of the hexagonal lattice on which the columns are positioned. This is apparent from the scattering data of Figure 5. In the context of filtration, the principal concern is for control over the orientation of the 18


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columns as this dictates tortuosity. By contrast the positional order of the columns plays no role in the transport properties.

It is worth noting here that it is decidedly non-trivial to a priori to determine the anchoring condition of a given LC species at a given surface. In the present case we did not seek to identify the underlying molecular mechanisms governing the observed anchoring condition – this is beyond the scope of the current work – but to simply leverage the observed anchoring to control the morphology of the mesophase. Our preliminary studies by POM have indicated that nanopore orientation is insensitive to the surface chemistry of the solid substrates in contact (See Figure S6, Supporting Information). We changed the hydrophobic/hydrophilic properties of the substrate surface using three different silanes with –NH2, –CH3, and –CF3 terminal groups and found that on all these substrates the nanopores displayed in-plane orientation in the open-to-air geometry. Nevertheless, their orientation shifted to vertical once a PDMS pad was placed on top of the film. It is possible that the columnar nanopores exhibit intrinsically vertical orientation under confinement regardless of the surface chemistry of the substrate that may act only as a mechanical barrier. Previous studies have found that systematic variations of surface energy and roughness of solid substrates result in limited influence on the anchoring of thermotropic Colh mesophases formed by discotic mesogens that orient homeotropically under confinement but planarly when exposed to the air.32-34 This may be due to the fact that the difference of the LC/solid interfacial energy between homeotropic and planar anchoring of a Colh phase at a solid substrate ∆γsolid is small. For example, the value of ∆γsolid of a Colh phase formed by a discotic mesogen was estimated to 19


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be ca. 10−5 J/m2.35, 36 One the other hand, the air interfacial tension of the Colh phase with planar anchoring γair has been measured to be 7×10−2 J/m2, being 3 orders of magnitude bigger than ∆γsolid.36, 37

Conclusion

In conclusion, we have devised a simple strategy to produce polymeric thin films with physically continuous and vertically-aligned, 1-nm-diameter pores. Our method relies on confinement and subsequent photo-cross-linking of thin films of a columnar mesophase formed by a wedge-shaped amphiphilic monomer. Although homeotropic alignment has been achieved

in

wedged-shaped

LC

monomers for forming

bulk ordered

polymer

nanostructures,27 it has been a challenge to obtain thin polymer films with vertically oriented 1 nm pores. TEM, SAXS and POM studies provide clear experimental verification of the role of different anchoring conditions in producing the observed morphologies, and of the physical continuity of nanopores through the entire film thickness. Thin films of any desired area can be easily processed by this technique to achieve vertical nanopore orientation, as the only requirement is for the area of the substrate and the confining material, glass or PDMS, to be matched with that of the film. Although both PDMS and glass can be utilized, one immediately recognizes the significant difference in terms of scalability and ease of fabrication afforded by PDMS. It is not difficult to envision a continuous process that involves soft conformal contact with PDMS, followed by a simple heating, cooling, and UV exposure, to rapidly fabricate large quantities of these nanostructured materials with 20


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uniform-diameter, aligned pores spanning the film thickness.

At present the need for mechanical support for these films can be provided by a lift-off technique we have developed which floats the films onto a microporous material which serves as a membrane gutter layer, as schematically illustrated in the Supporting Information (Figure S7). Future efforts will target the elimination of this step by fabrication directly in contact with the support material. Nanostructured films with the 1-nm oriented morphology and channel continuity are highly attractive for use as size-selective separation membranes in a range of important applications including wastewater treatment, water softening, and separation of small organic contaminants. The unique combination of thermodynamically defined nm-scale pore sizes, physical continuity and scalable processing position these materials to overcome the longstanding permeability-selectivity tradeoff that is inherent to the operation of conventional membranes operating by solution-diffusion to reject solutes at the ~1nm length scale.

Experimental Section Materials. The synthesis and characterization of the amphiphilic LC monomer Na-GA3C11 can be found in a previous report.25 The silicone elastomer kit was obtained from Dow Corning. All other chemicals were purchased from Aldrich and used as received. For UV-induced cross-linking, Na-GA3C11 was doped with a small amount of a commercially available radical photo-initiator, diphenyl(2,4,6-trimethylbenzoyl)phosphine oxide (0.5 wt %). Glass slides and silicon wafers were cleaned by acetone and water and then dried by nitrogen gas before use. PDMS elastomer pads were prepared by following the standard procedure. 21


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Briefly, a homogeneous fluid mixture of the silicon monomer and a curing agent with a weight ratio of 9:1 was poured into a Petri dish in which a clean silicon wafer was placed at the bottom. Subsequently, bubbles trapped in the fluid mixture were eliminated by vacuum. Thermal curing at 70 °C resulted in a transparent PDMS elastomer pad (5 mm thick) with a smooth surface.

Preparation of LC and Polymer Films. LC films with a thickness above 20 µm were prepared by sandwiching the monomeric Na-GA3C11 mesophase by two substrates separated by a spacer with the corresponding thickness. LC films with a thickness ranging from 200 nm to 10 µm were obtained by casting of an LC/THF solution onto a glass or silicon substrate followed by solvent evaporation. The thickness of the film was generally controlled by the concentration and the amount of the solution cast on the substrate. Polymer films were prepared by photo-cross-linking of the aforementioned monomeric LC films (with 0.5 wt % radical photo-initiator) obtained by following the above described procedure.

Soft Confinement Fabrication of Polymer Films. A PDMS elastomer pad (1.5 cm × 1.5 cm × 0.5 cm ) was pressed onto a LC film (doped with photo-initiator) lying on a substrate with an additional weight of c.a. 500 g applied to the PDMS pad to ensure the tight contact. The LC film was then thermally annealed by heating up to the isotropic phase (75 °C) and allowing it to cool back to room temperature at a rate of 0.1 °C/min. Subsequently, a polymer film was obtained by cross-linking the LC film immediately through exposure to 365 nm UV light (8W UVL-18 EL lamp at a distance of roughly 10 cm) for 24 h at room temperature. 22


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Polarized Optical Microscopy (POM) and Conoscopy. POM images were obtained by a Zeiss Axiovert 200 M inverted microscope. Conoscopy studies were performed using a Zeiss Axio Imager M2m microscope. The conoscopic images were obtained with a 40× objective and a Bertrand lens introduced between the analyzer and the ocular.

Transmission Electron Microscopy (TEM). To obtain a thin specimen for cross-sectional TEM imaging, a polymer film was embedded in an epoxy and then the epoxy along with the sample was cured at 60 °C overnight to enhance the rigidity for microtoming. The cured epoxy block was then sectioned at room temperature by a diamond knife mounted on a Leica EM UC7 ultramicrotome. Thin sections (~60 nm thick) on top of water were picked up onto a TEM grid and stained in vapor of a 0.5 wt % aqueous solution of RuO4 for 10 min. Specimens were then visualized by an FEI Tecnai Osiris TEM with an accelerating voltage of 200 kV. To visualize a polymer film with a thickness of 350 nm along the through-plane direction, the film was detached from the substrate by sonication in a water/ethanol (9/1 v/v) bath. Specimens were then obtained by placing several drops containing suspension of tiny pieces of films onto a TEM grid.

Small Angle X-ray Scattering (SAXS). SAXS measurements on the Colh mesophase and cross-linked polymer films of Na-GA3C11 were performed using a Rigaku 007 HF+ instrument equipped with a rotating anode Cu Kα X-ray source (λ = 1.542 Å) and a 2-D Saturn 994+ CCD detector. A silver behenate standard (d-spacing of 58.38 Å) was employed 23


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for the calibrations of the resultant 2-D SAXS data. All the 2-D scattering patterns were integrated into 1-D plots of scattering intensity (I) versus q, where q = 4πsin(θ)/λ and the scattering angle is 2θ. Particularly for measurements on a polymer film, as schematically illustrated by Figure 5a, 20 layers were stacked along the film normal to enhance the scattering intensity. The sample was then mounted to the SAXS instrument such that the incident X-ray beam was parallel to the film plane.

Ionic Conductivity Measurements. Open samples prepared on the ITO glass were allowed to equilibrate in a humidity box at a 100 % relative humidity for 12 hours. The sample was then sandwiched by placing a second ITO electrode on the surface and the two electrodes were clamped using small spring clamps on each side and a micrometer gauge in the center. The sandwiched samples were allowed to equilibrate in the humidity box and the impedance responses were continuously monitored over a period of 12 hours using a Solartron 1260 Frequency Response Analyzer connected to the ITO electrodes. The AC voltage amplitude was varied from 10 to 150 mV to ensure that the system response in the range tested was linear, and a 100 mV voltage amplitude was thereafter chosen for all the measurements. The AC voltage frequency was swept from 50 mHz to 1 MHz. All measurements were performed at room temperature and the through-plane resistance was estimated by fitting the frequency response of the system to an appropriate equivalent circuit.22 Careful measurements of the film thicknesses and surface area were made after the resistance measurement to facilitate conversion to accurate conductivity values.

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Supporting Information Available: Supplementary SAXS, TEM, and POM images. This material is available free of charge via the Internet at http://pubs.acs.org.

Acknowledgements. Financial support from NSF (CMMI-1246804) is gratefully acknowledged. Facilities use at Yale was supported by the YINQE and NSF MRSEC program (DMR-1119826). The authors thank Mike Degen (Rigaku, Inc.), Brandon Mercado (Yale CBIC), and Aniko Bezur (Yale IPCH) for technical assistance. C.O.O. acknowledges additional financial support from NSF (DMR-1410568) and from a 3M Nontenured Faculty Award. M.E.T. acknowledges support from the NSG Graduate Research Fellowship DGE-1122492. The monomer synthesis work done at CU boulder was partially supported by the DOE ARPA-e program (D1-AR0000098) and the U.S. Bureau of Reclamation (R13AC80040).

References

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