Controlling Nanoscale Assembly in Bulk, Thin Film, and Solution Block

Invited Feature Article pubs.acs.org/Langmuir

Interfacial Manipulations: Controlling Nanoscale Assembly in Bulk, Thin Film, and Solution Block Copolymer Systems Sarah E. Mastroianni and Thomas H. Epps, III* Department of Chemical and Biomolecular Engineering, University of Delaware, Newark, Delaware 19716, United States ABSTRACT: Nanostructured soft materials from self-assembled block copolymers (BCP)s and polymer blends can enable the reliable, high-throughput, and cost-effective generation of nanoscale structural motifs for many emerging technologies. Our research group has studied the phase behavior of BCPs in bulk, thin film, and solution environments with a particular focus on using interfacial manipulations to control self-assembly and to access a vast array of nanoscale morphologies and orientations. These interfacial manipulations can be synthetic alterations that are directly incorporated into the BCP chain to modify polymer− polymer interactions, post-polymerization and non-synthetic modifications that affect block interactions, or changes to the polymer specimen’s external surroundings to control selfassembly in a confining environment. Herein, we describe methods that we have employed to manipulate BCP self-assembly for various application targets, and we discuss the key effects of such manipulations on the resulting nanoscale morphologies.

1. INTRODUCTION

2. BULK Linear diblock and triblock copolymers can self-assemble into various nanoscale structures in the bulk, making them desirable for a number of applications. Co-continuous polymer network structures are of particular interest because of their interconnected channels that provide well-defined diffusion pathways. Furthermore, these nanoscale ordered networks normally possess superior mechanical properties relative to their 1D and 2D counterparts (e.g., lamellae and cylinders).3 However, in conventional diblock copolymer systems, these networks typically form only in narrow composition windows. Thus, the ability to obtain a detailed understanding of (and control over) the parameters that influence self-assembly while retaining the ability to achieve various property specifications is desirable. In linear BCPs, bulk self-assembly primarily is governed by the relative volume fractions of the blocks ( f i) and block segregation strength χN, where χ is the Flory− Huggins interaction parameter and N is the degree of polymerization. Much of our work has focused on manipulating the effective interaction parameter, χeff, between polymer blocks. Though this χeff cannot be defined strictly by a Flory−Huggins-type interaction parameter, it provides a general framework for understanding the influence of changes to the copolymer system on the segregation strength (χeffN). Manipulations to χeff can be achieved by altering the traditional BCP structure during copolymer synthesis or via postpolymerization modifications (i.e., through hydrogenation (synthetic means) or the incorporation of additives such as

Increasing interest in nanostructured materials for applications including energy storage and generation, biological and drug delivery systems, data transfer and storage, and separations devices has driven research focus toward the control of smaller and smaller structures in materials systems. The self-assembly of soft materials permits precise control of nanoscale structures and can greatly enhance a material’s mechanical, thermal, and transport properties. One particular focus of our research effort is block copolymers (BCP)s, which can self-assemble into periodic structures on an approximate 10−100 nm size scale. BCPs are comprised of two or more chemically distinct blocks, with a plethora of possible polymer chemistries and molecular architectures that can generate a vast array of nanostructures. Many potential soft materials applications require specific orientations, nanostructures, or functionality. For example, nanoporous and oriented cylindrical morphologies may be desirable for separations or templating applications,1,2 while 3D networks of conducting domains may be desirable for energy generation and storage systems.3 Therefore, the ability to tune the parameters that govern self-assembly is of critical importance. Control of the self-assembly of BCP nanostructures can be achieved by modifying interfaces between polymer blocks through synthetic (or non-synthetic) means or by manipulating the sample environment/surroundings to alter interactions between the polymers and surfaces, substrates, or solvents. Herein, we highlight our work in controlling selfassembly via interfacial manipulations of BCPs in bulk, thin film, and solution systems. © 2013 American Chemical Society

Received: December 5, 2012 Revised: January 25, 2013 Published: February 13, 2013 3864

dx.doi.org/10.1021/la304800t | Langmuir 2013, 29, 3864−3878

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Invited Feature Article

length) and interfacial interactions between the polymer blocks.4 The second example demonstrates our ability to modify the chemical interactions in a tapered BCP system while retaining desirable nanoscale network morphologies.5 In the first example, we employed a semi-batch monomer feed for the anionic polymerization of symmetric and lamellaeforming copolymers containing a linear taper (styrene/ isoprene) region between pure polystyrene (PS) and polyisoprene (PI) blocks.4 We synthesized both normal-tapered [P(I-IS-S)] and inverse-tapered [P(I-SI-S)] BCPs with welldefined taper lengths and compositions. In our naming convention for tapered BCPs, we denote a pure component block with a single letter (I for isoprene, S for styrene) and a linear tapered region with two letters, the order of which designates the direction of the compositional profile (IS tapers from isoprene-rich to styrene-rich compositions and SI tapers from styrene-rich to isoprene-rich compositions). As determined by dynamic mechanical analysis (DMA), tapered BCPs had lower order−disorder transition temperatures (TODT)s than their non-tapered [P(I−S)] counterparts (i.e., copolymers with the same overall composition and similar molecular weight), as shown in Table 1.4 This reduction in TODT was a

complexing salts (non-synthetic means)). We employ both of these methods to control interactions in our BCP systems. In the following subsections, we highlight our major efforts in the synthetic manipulations of diblock copolymers (section 2.1), discuss triblock copolymer phase behavior and synthetic manipulations to triblock copolymer systems (sections 2.2−2.5), and conclude with small-molecule additives for tuning copolymer interactions (section 2.6). 2.1. Tapered Diblock Copolymers.4,5 Typically, interaction parameters between blocks are determined by the specific polymer block chemistry; however, it is often desirable to control BCP interactions independently of the copolymer constituents, especially when the particular chemical, mechanical, thermal, or transport properties afforded by a specific BCP system are required for a given application. Thus, one of our goals has been to decouple polymer interfacial interactions (χeff) from block polymer chemistry and molecular weight. We have accomplished this goal through the introduction of short tapered segments with well-controlled composition profiles between the pure copolymer blocks in our BCPs. For the majority of AB diblock copolymers, the composition profile along the polymer backbone changes abruptly at the junction between A and B copolymer segments. This standard composition profile can be modified easily by introducing a transition region (Figure 1) between the pure A and B blocks that tapers from A to B units (normal-tapered) or from B to A units (inverse-tapered).6

Table 1. Molecular Characteristics and TODT Values of Select Non-tapered and Tapered BCPs4 polymer

taper length (vol %)a

Mn/ g mol−1b

volume fraction ( f I)c

TODT/ °Cd

P(I-S) P(I-IS-S) P(I-IS-S) P(I-SI-S) P(I-SI-S)

0 20 35 20 20

27 900 27 800 34 000 25 900 32 500

0.50 0.49 0.48 0.51 0.49

177 175 171 148 164

a

The taper length is based on the volume fraction of the tapered block compared to the total polymer. bNumber-average molecular weights (Mn) were determined by gel permeation chromatography (GPC) using polystyrene standards. cVolume fractions were determined from peak integrations in 1H NMR using homopolymer densities at 140 °C: ρPI = 0.83 g mL−1 and ρPs = 0.969 g mL−1. dTODT was determined from isochronal temperature ramps where the sudden drop in G′ on second heating indicated the TODT.

direct result of the increased compatibility between blocks afforded by the incorporation of the tapers. Furthermore, the greater drop in TODT for the inverse-tapered materials compared to that for their normal-tapered counterparts was attributed to the long sequences of chemically different repeat units next to the pure blocks, which increased interfacial mixing. The taper length also can be used to fine tune the TODT change. Maintaining a constant molecular weight and increasing the taper volume fraction (decreasing the size of the pure end blocks) resulted in an expected decrease in TODT.4 However, we noted that samples with approximately identical pure end-block (S and I) lengths but with longer tapered middle regions (and thus larger overall BCP molecular weight as a result of the larger taper volume percentage) also had moderately lower TODT values, suggesting that the increased block compatibility was due to an increase in the effective interfacial width, as opposed to a decrease in the absolute length of the pure endblocks.4 Fortunately, as shown in the second example below, the modification of the interaction parameters did not compromise our ability to form desirable complex network structures when

Figure 1. Illustration of the density profile along a polymer chain for a (a) non-tapered, (b) normal-tapered, and (c) inverse-tapered diblock copolymer.

From a synthesis point of view, interfacial modification through tapering constitutes an additional design parameter that permits both the manipulation of interfacial structure and a greater understanding of the influence of that interfacial structure on bulk copolymer self-assembly. Two examples of tapered BCP work from our group are discussed below. The first example illustrates one of the first systematic efforts to establish a direct link between the taper composition (and taper 3865


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Figure 2. SAXS and TEM for (a, b) non-tapered P(I-S), (c, d) normal-tapered P(I-IS-S), and (e, f) inverse-tapered P(I-SI-S). SAXS peaks are indexed according to the Ia3̅d space group, characteristic of the double-gyroid (Q230) morphology. TEM micrographs are representative of the double-gyroid network morphology and depict the “wagon-wheel” pattern characteristic of the [111] projection. (See ref 5 for the [125] projection.) TEM scale bars represent 100 nm in all images. (g) Isochronal storage modulus (G′) vs temperature was obtained using a frequency of 0.5 Hz and a heating rate of 1 °C/min for P(I-S) (●), P(I-IS-S) (⧫), and P(I-SI-S) (▲). Order−disorder transition temperatures (TODT values) are indicated by arrows. An order−order transition (OOT) in the P(I-S) sample is indicated by ▼. Adapted with permission from ref 5. Copyright 2011 American Chemical Society.

governed by two independent volume fractions ( fa and f b, with fc = 1 − fa − f b) and three interaction parameters (χAB, χAC, and χBC). As such, the resulting phase space can be represented by a triangular phase prism with each side defined by the corresponding diblock copolymer phase diagram. This phase space often is simplified by looking at a single slice of the phase prism at nearly constant segregation strength, illustrated similarly to a ternary phase diagram. In triblock copolymer systems, interfacial interactions can have a significant influence on self-assembly. For example, it has been shown that network structures are favorable in ABC systems with nonfrustrated block order, that is, χAC > χBC > χAB, to ensure that the most unfavorable interaction is not forced by the block connectivity.10 Additionally, statistical segment length ratios between the blocks are another set of parameters that can affect the self-assembly of BCPs. Tyler et al. developed a theoretical phase diagram based on self-consistent mean-field theory (SCMFT) for an idealized, nonfrustrated system with equal statistical segment lengths (bA = bB = bC) and equal segregation strengths between adjacent blocks.12 In a follow up to this theoretical work, we explored the bulk morphological behavior of the poly(isoprene-b-styrene-bmethyl methacrylate) (ISM) system, and we experimentally constructed a phase diagram for the ISM system (Figure 3) and compared it to Tyler’s theoretical phase diagram.7,12,13 This

the tapering method was extended to non-symmetric copolymer compositions. We synthesized normal- and inverse-tapered BCPs with overall compositions that fell within the expected network-forming (double-gyroid) region for nontapered P(I-S) BCPs.5 Both tapered materials formed the double-gyroid network structures, as determined by small-angle X-ray scattering (SAXS) and transmission electron microscopy (TEM), and they showed significantly decreased TODT values compared to their non-tapered counterparts (Figure 2).5 (Note that the taper volume percentages were 30% in this example.) Thus, we were able to control the thermal properties and chemical interactions that enhance polymer processability without compromising the ability to form the complex nanoscale networks that are desirable for transport applications. 2.2. Neat Triblock Copolymers.7 Although the introduction of tapered regions is one method of manipulating BCP nanostructures, an alternative approach involves the incorporation of a third block to generate ABC triblock copolymers. This third block adds complexity to the system and allows access to a wider variety of nanostructures,8 including three distinct networks.3,9−11 Furthermore, the network nanostructure phase window is wider in ABC triblock copolymer systems and permits easier access to, and flexibility in, generating these complex morphologies. In contrast to AB diblock copolymer self-assembly, ABC triblock copolymer self-assembly is 3866


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mechanical strength, and the poly(methyl methacrylate) (PMMA) block can be removed easily via etching to produce nanoporous templates. We synthesized ISM triblock copolymers that formed threephase lamellar (LAM3), hexagonally packed cylinder (HEX),13 alternating-gyroid (Q214), and disordered (DIS) morphologies, which were characterized via SAXS, TEM, and DMA.7 The compositions at which these phases were found were consistent with the theoretical phase diagram by Tyler et al.12 Though consistency was noted between the experimental and theoretical diagrams, the synthesis of a large number of polymer specimens to complete the phase diagram would be time-consuming. Thus, as described below, we pursued a different approach to validating the ISM phase behavior. 2.3. Blended Triblock Copolymers.13 To refine the phase boundaries of the ISM system without additional synthesis, constituent homopolymers were blended into the ISM system to change the relative component volume fractions and explore the resulting phase transformations. Judicious copolymer blending enabled the rapid exploration of the triblock copolymer phase space, allowing us to locate regions of specific interest for further study. Through a series of blend studies, we accessed two additional network structures, the double-gyroid (Q230) and orthorhombic network (O70), and we refined the boundaries of the compositional regime for morphologies formed in the neat (nonblended) triblock copolymer specimens (Figure 3).13 Our blended systems were consistent with the phase behavior of neat triblock copolymers of similar compositions. Additionally, the blended systems were internally consistent (i.e., specimens at the same final composition but blended from different neat triblock copolymers and containing different homopolymers exhibited the same final morphology).13 This result supported our blending approach in the exploration of a triblock copolymer system. Though we identified the majority of the structures found in the SCMFT phase diagram,12 we noted slight discrepancies that likely resulted from differences in the energetics when homopolymer was incorporated into the BCP. However, our approach does

Figure 3. Experimentally generated phase diagram for the poly(isoprene-b-styrene-b-methyl methacrylate) (ISM) system near the styrene-rich compositional regime. Component volume fractions, f i (i = I, S, or M), are indicated on the axes. Six nanostructures were identified: LAM, Q214, Q230, O70, HEX, and DIS. Estimated phase boundaries are color coded by nanostructure and paired with corresponding representative TEM micrographs. TEM scale bars represent 20 nm. Colored circles denote the composition of neat triblock copolymers (outlined in black) and blended samples (no outline). Reproduced with permission from ref 13. Copyright 2010 American Chemical Society.

direct comparison was justified because statistical segment length ratios in the ISM system are close to unity (bM/bI = 0.95, bI/bS = 1.11, bS/bM = 0.95),14 and the system has a nonfrustrated block order (χIM ≫ χIS ≥ χSM).13 Additionally, the ISM system is of interest for generating nanoporous membranes for filtration applications1 because the combined PI/PS blocks provide toughness,15 the PS domain provides

Scheme 1. Synthesis of P(I-IS-S-SM-M) via a Combination of Anionic Polymerization, ATRP, and Huisgen 1,3-Dipolar Cycloaddition “Click” Chemistrya

a

The illustrations represent the density profiles along the block copolymer backbone. Adapted with permission from ref 16. Copyright 2012 American Chemical Society. 3867


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Figure 4. (a) SAXS and (b−d) TEM for a tapered P(I-IS-S-SM-M) triblock copolymer. (a) SAXS peaks are indexed according to I4132 symmetry characteristic of the alternating-gyroid (Q214) morphology. TEM micrographs are representative of the alternating-gyroid network morphology and depict the (b) [111] and (c) [001] reflection planes. The inset images are simulated TEM micrographs from Epps et al.8 (d) Reconstructed 3D image from TEM tomography with only PI (blue) and PMMA (green) domains shown for clarity. TEM scale bars represent 100 nm, and TEM samples are stained with OsO4 (PI appears the darkest, PS appears gray, and PMMA appears white). Reproduced with permission from ref 16. Copyright 2012 American Chemical Society.

synthesis of P(I-IS-S) materials and atom-transfer radical polymerization [ATRP] for the synthesis of P(SM-M) materials) and “click” chemistry (to generate tapered P(I-ISS-SM-M) materials) (Scheme 1).16 This method represents a robust framework for generating interfacially modified nanostructured materials. By incorporating tapered regions between blocks, we can control the block interfacial interactions in the ISM triblock copolymer system independently of the molecular weight and block constituents. This manipulation is of particular interest because it is often difficult to generate stable network structures at higher segregation strengths in triblock copolymer systems because of kinetic and thermodynamic restrictions.18 Detailed analysis is ongoing; however, we have demonstrated the ability to retain complex network structures while manipulating the copolymer interfacial profile. As one example, we examined the morphology of an ISM block copolymer with volume fractions of f I = 0.26, f S = 0.59, and f M = 0.15 containing normal I/S and S/M tapers that were 23 vol % and 17 vol %, respectively, of the overall polymer. SAXS and TEM (and TEM tomography) analyses were consistent with a Q214 alternating gyroid assignment (Figure 4). This result indicates that tapering does not completely disrupt the delicate balance of forces necessary to form complex nanoscale network structures, allowing us to use this synthetic modification of interfacial interactions as a handle for the de novo tuning of copolymer properties.

provide detailed information for screening the regions of interest for further study. 2.4. Tapered Triblock Copolymers.16 As described above, triblock copolymers offer advantages over diblock copolymers in terms of the variety of nanostructures and chemical complexity. Although we have demonstrated the ability to control block interactions independently of block chemistry and molecular weight in tapered diblock copolymers,4,5 it is of interest to extend this control to triblock copolymer systems. The understanding of and ability to manipulate self-assembly in triblock copolymer systems provide an opportunity to design functional materials precisely, particularly if the ability to generate various complex nanostructures is achieved while maintaining both block chemistry and block order. Theoretical SCMFT studies12 and Monte Carlo simulations17 that examined BCP morphologies as a function of composition and segregation strength have predicted that ABC morphologies are significantly influenced by the relative strengths of interblock interactions, thus resulting in distinct phase diagrams in which triblock copolymers of the same overall composition and block sequence formed different nanostructures. We are experimentally testing this hypothesis by manipulating the interblock interaction parameters via interfacial modifications (tapering) in our ISM triblock copolymers. As with the diblock copolymer system, we utilized a semi-batch monomer feed protocol. Our synthetic framework was developed using a combination of two living polymerization methods (anionic polymerization for the 3868


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PEO when PEO was doped with salt),20 in which the change in χeff was estimated by correlating the change in domain spacing (d*) to χ in the strong segregation regime (d* ≈ χ1/6).26 For all of the lithium salt studies, χeff increased linearly with salt concentration with the mathematical relationships being PSPEO/LiAsF6, χeff = 5.696Csalt + 0.053; PS-PEO/LiClO4, χeff = 5.532Csalt + 0.053; and PS-PEO/LiCF3SO3, χeff = 3.902Csalt + 0.053, in which the salt concentration, Csalt, is defined by [Li]/ [EO]. However, we note that the slope was dependent on the lithium counterion, with the weaker Lewis bases having a larger slope (PS-PEO/LiAsF6 > PS-PEO/LiClO4 > PS-PEO/ LiCF3SO3). A weaker Lewis base had a weaker interaction with Li+, thus Li+ likely has a stronger association with PEO chains. The strength of these associations resulted in increased chain stretching to accommodate the ions as manifested through the increased domain spacing. In addition to exploring changes in phase behavior induced by salt doping, we examined the influence of these salts on PEO-block crystallization.21,22 In particular, by utilizing a mixed-salt system (i.e., mixtures of two lithium salts in which each salt contained a different counterion) that suppressed conductivity-inhibiting crystallization in the PEO block,25 we found that the room-temperature ionic conductivity increased by at least an order of magnitude in comparison to that of the single-salt counterpart systems.22 Whereas again we found linear relationships between χ eff and the overall salt concentration, the mixed-salt ratio (i.e., counterion composition) did not have a significant effect on the domain spacing.22 Finally, we explored the relationship between ionic conductivities and morphology in PS-PEO systems doped with LiClO4.23 In the neat (undoped) polymer, PS-PEO underwent a transition from hexagonally perforated lamellae (HPL) to HEX when increasing the PEO volume fraction ( f EO) from 0.70 to 0.75, whereas the salt-doped sample underwent transitions from LAM to HPL to HEX over the same volume fractions at a fixed salt concentration. The 3D conducting pathways of the PEO matrix inherent in HPL and HEX morphologies exhibited much higher normalized conductivities (as determined by ac impedance spectroscopy) than the 2D conducting pathways in the LAM system, even after accounting for non-random domain orientations and slight differences in the PEO molecular weight.23 This work demonstrates the importance of accounting for the modification of interfacial interactions upon salt doping in order to design systems with the desired morphologies and, correspondingly, the requisite high ionic conductivities. Furthermore, it highlights the need for the facile design of nanoscale materials containing multiply continuous, network-like, nanoscale domains.

2.5. Post-polymerization Synthetic Modification: Selective Hydrogenation.19 Our exploration of the ISM phase space coupled with our ability to control block interactions gave us wide access to a number of desirable morphologies; however, many of the potential applications for the ISM materials are limited by the poor thermal and oxidative stability of the PI block. Thus, we selectively hydrogenated the PI to form poly(ethylene/propylene) (PEP),19 which is more resistant to thermal and oxidative degradation. The resulting poly(ethylene/propylene-b-styrene-b-methyl methacrylate) (EPSM) triblock copolymers are promising materials for nanoporous membrane applications. Upon hydrogenation, the ratios of statistical segment lengths of the system were more asymmetric than in the ISM precursor system (bM/bEP = 0.85, bEP/bS = 1.25), and the interaction parameters between the blocks changed (χEPS > χIS and χEPM> χIM). These effects were manifested as a shift in the EPSM phase boundaries in comparison to the ISM system. As done for the ISM system, we rapidly explored the EPSM system by blending neat EPSM triblock copolymers with constituent homopolymers. The resulting phase diagram indicated an expansion of the network regions, particularly the Q230 region, at the expense of the HEX region.19 This post-polymerization chemical modification provides another experimental example of how manipulating block interactions, in this case, by changing both χ parameters and statistical segment length ratios, influences self-assembly. In this work, we gained a better understanding of how to account for these changes when designing materials starting from the ISM precursor so that we can target specific morphologies, such as networks, upon hydrogenation to more desirable EPSM copolymers. 2.6. Post-polymerization Modification: Salt Doping of Ion-Conducting BCPs.20−23 Although the hydrogenation of PI domains is one example of an applications-driven postpolymerization modification that affects block interactions and self-assembly, another example is the incorporation of additives such as complexing salts in BCP electrolytes. BCP electrolytes are promising materials for electrochemical devices because of their nanoscale domains, well-defined transport properties, and improved mechanical strength compared to some pure polymer electrolytes.24 Poly(ethylene oxide) (PEO) is one of the most studied polymer electrolytes for lithium ion battery applications as a result of its high concentration of electron pairs, the polarizability of the ether groups, and its low glass-transition temperature (Tg).25 For battery applications, these polymer electrolyte materials are doped with lithium salts to introduce the conducting ions; however, the choice of lithium-salt counterions and salt concentrations has a significant effect on ionic conductivities. Furthermore, the addition of the required lithium salts to BCP electrolytes affects block interactions and thus nanoscale morphology. Because the nanoscale structure also affects the ionic conductivity, it is important to understand the relationship between lithium salt addition and block interactions when designing BCP electrolyte systems. In our work, we systematically examined the effects of lithium salt concentration and counterions on conductivity and nanoscale structure in poly(styrene-b-ethylene oxide) (PS-PEO). In one example, we studied systems of PS-PEO/LiX where X = ClO4−, CF3SO3−, or AsF6−.20,21 In all cases, an increased salt concentration resulted in larger domain spacings, and the percentage increase in the spacing was only partially attributed to the volume of the added salt. The additional increases were largely due to the change in χeff (effective χ between PS and

3. THIN FILM BCP thin films are ideal materials for controlling and templating the nanoscale features required for various applications and devices including electronics, coatings, and biologically active surfaces. Many of these applications require specific morphologies and domain orientations. For example, nanostructures oriented perpendicular to substrate surfaces are often desirable in transport or lithographic applications.1,2 However, in addition to the copolymer volume fractions and segregation strengths that are the key parameters influencing morphologies in bulk systems, the phase behavior in BCP thin films (∼100 nm thickness) also is governed by the film thickness and interfacial interactions with the confining 3869


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environment.27,28 In particular, domain orientations are affected strongly by free and substrate surface interactions, and control over orientation is often limited, or driven, by surface energetics. Through several approaches, we demonstrated the ability to manipulate these interactions with precise free or substrate surface modifications in order to tune phase behavior. We also explored the effects of certain processing conditions that can result in the kinetic trapping of various morphological orientations in BCP thin films. We are interested in highthroughput techniques for the efficient exploration of the vast BCP thin film parameter space. Specifically, we focus on gradient methodologies in which a continuous and positiondependent range of variables (e.g., film thickness, substrate surface energy, free surface energy) is generated on a single surface. This approach provides a rapid means for screening BCP thin film behavior. 3.1. Film Thickness Gradients.29,30 To facilitate the exploration of the thin film parameter space, we developed a flow-coating device that provides exquisite and tunable control of the polymer film thicknesses on a variety of substrates.29 Because the film thickness has a dramatic effect on the BCP morphology and orientation, it is important to have detailed control over this key parameter. We employed the flow-coating technique to fabricate continuous thickness gradients on surfaces of uniform composition when screening the effects of other parameters such as solvent annealing or thermal protocols on polymer morphology. Combined with optical microscopy (OM), these thickness gradients provided a facile means for quickly screening the interfacial processing conditions that produce perpendicular versus parallel structures. For example, island and hole microstructures (which arise to accommodate film thicknesses incommensurate with polymer domain spacing) are indicative of parallel morphologies whereas smooth films are indicative of perpendicular structures (or films with thicknesses commensurate with integer (or half-integer) multiples of the domain spacing). Because island and hole structures can be catalogued via OM, OM on gradient substrates is an efficient screening method that provides insight into a range of BCP thin film morphological transitions on a single substrate.31 3.2. Substrate Surface Modification.31,32 In addition to gradient thickness studies, flow coating also facilitates the production of films of precise and uniform thickness on surface energy gradients (something not afforded by common spincoating techniques), which permits screening of substrate and free surface effects on BCP assembly, as discussed in the following sections. Substrate surfaces have a significant influence on BCP thin film morphology. Surfaces may be “neutral” or preferential for one block such that surface interactions can drive the orientation of BCP domains relative to the surface or even facilitate BCP dewetting.33 We employ surface energy/chemistry gradients on substrates as a highthroughput technique to manipulate the polymer/substrate interface. Previous gradient-generation techniques offered limited tunability, required complex processing, or could be utilized only with specific systems.34−36 Thus, in our work, a different approach was employed where gradient profiles were controlled by a vapor deposition method using functionalized chlorosilanes to modify the surface chemistry (and thus surface energy) of silicon substrates.31,32 Tunable gradients were achieved by utilizing a vapor deposition device (Figure 5) containing liquid chlorosilane reservoirs whose contents were volatile under dynamic vacuum. The reservoir size and position

Figure 5. Schematic of vapor deposition device (a and b). (a) Reservoirs and substrates were placed in a Teflon insert. (b) The Teflon insert was loaded into the deposition chamber, followed by the application of dynamic vacuum for deposition. (c) Representative compositional characterization data shows a linear surface chemistry composition (mole fraction of methacrylsilane) with respect to position along the surface. Reproduced with permission from ref 31. Copyright 2009 American Chemical Society.

relative to the substrate were manipulated to achieve the desired cross-deposition profiles of functional chlorosilanes.32 We used this versatile device to manipulate surface chemistries and surface energies in order to examine BCP thin film behavior rapidly. In one example, we used chlorosilanes with functionalities similar in structure to PS (benzylsilane) and PMMA (methacrylsilane) to generate a nearly linear gradient in substrate surface chemistry as shown in Figure 5.31 The gradients in surface chemistry and surface energy were measured using X-ray photoelectron spectroscopy (XPS) and contact-angle goniometry, respectively. Subsequently, a cylinder-forming PS-PMMA BCP was flow coated at constant film thickness onto the surface gradient. As expected, mixed chlorosilane compositions resulted in relatively neutral surfaces. In these regions, the lack of islands and holes along the gradient in optical micrographs (Figure 6a, ii−iv) corresponded to cylinders oriented perpendicular to the substrate surface as 3870


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Figure 6. (a) OM and (b) AFM of PS-PMMA films (i) on a benzylsilane-functionalized surface, (ii−iv) on three representative positions along the benzyl/methacrylsilane gradient {(ii) 1.5, (iii) 3, and (iv) 5 cm}, and (v) on a methacrylsilane-functionalized surface. Optical micrographs showed island and hole structures on the pure component surfaces and smooth films across a majority of the gradient, which correspond to the parallel and perpendicular cylinders, respectively, as seen in AFM. (c) The composition, reported as the mole fraction of methacrylsilane (xm), and the surface energy (γs) are provided for reference. (d) The lower illustration depicts the chemical functionality gradient along the substrate. Adapted with permission from ref 31. Copyright 2009 American Chemical Society.

Figure 7. Microfluidic device schematic for gradient SVA. Nitrogen is used as a carrier gas for solvent vapor, which enters through ports A and B, mixes in the mixing tree, and passes through annealing chambers 1−6. Chamber 7 is not attached to the mixing tree and serves as a control chamber. Adapted with permission from ref 37. Copyright 2011 American Chemical Society.

direct BCP domain orientation but also to screen quickly the regions of interest for nanoscale investigations. 3.3. Solvent-Vapor Annealing.37,38 Another attractive approach to controlling thin film self-assembly is solvent-vapor annealing (SVA). Traditionally, SVA has been used as an alternative to thermal annealing to impart mobility to kinetically trapped structures, making it a valuable technique for thermally sensitive and higher-molecular-weight systems.39−41 SVA promotes self-assembly by reducing the Tg of the blocks and increasing the chain mobility. Additionally,

confirmed by AFM (Figure 6b, ii−iv). On pure-component substrates and near the ends of the gradient (Figure 6, i and v), the surface was preferential for a specific block resulting in cylinders oriented parallel to the substrate surface as evidenced by the prevalence of island and hole structures in the optical micrographs (Figure 6a, i and v) and also confirmed by AFM (Figure 6b, i and v). Thus, by precisely tuning the surface energy/chemistry of silicon substrates using this vapordeposition device, we demonstrated the ability not only to 3871


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solvent sorption into the film influences the parameters governing self-assembly, including block interactions, volume fractions, and domain spacing. These solvent/polymer interactions can be tuned further by using mixtures of solvent vapors.42,43 Controlling the solvent removal rate also permits the kinetic trapping of desired morphologies.44,45 Furthermore, using SVA at the free surface may mitigate substrate surface preferences and eliminate the need for surface modification in some systems.44 Selective solvents may be used to create an environment in which one block is preferential for the free surface and may induce preferential block segregation, similar to preferential orientations induced by substrate modification.46−48 Such methods are of particular interest for systems that may be inhibited by the chemical modification of substrates, such as materials designed for interfacial transport. A vast parameter space, including solvent choice, solvent concentration/swollen film thickness, and solvent removal rate, governs the BCP morphology in solvent-vapor-annealed systems. Hence, an investigation of BCP film responses to various SVA conditions warrants the development of highthroughput methods to explore thin film self-assembly. To address this issue, we designed a solvent-vapor microfluidic mixing tree (Figure 7) as a high-throughput approach to scan the parameter space using SVA gradients.37 In one example, SVA composition-dependent morphologies were identified in BCP thin films of a 118 kg/mol poly(styrene-b-isoprene-bstyrene) (SIS) (f S = 0.134, f I = 0.732, f S = 0.134) BCP annealed with n-hexane vapor (preferential for the PI block) and tetrahydrofuran (THF) vapor (slightly preferential for the PS block but a good solvent for both blocks). The microfluidic mixing device provided a facile method for determining the effects of solvent-vapor concentration and composition on SIS thin films (Figure 8). High concentrations of n-hexane vapor promoted perpendicularly oriented cylinders whereas increasing THF vapor concentrations resulted in parallel cylinder orientations. Additionally, the swollen film thickness decreased with a decreasing n-hexane-to-THF ratio. Despite this lower overall solvent concentration in the film, the higher THF concentrations resulted in greater chain mobility and longrange order. These explorations examining the effects of solvent selectivity and concentration on the thin film morphology demonstrate the utility of our high-throughput microfluidic mixing device in quickly scanning the parameter space and providing crucial insight into several parameters that govern self-assembly in SVA systems. In addition to morphology control through solvent selection, we also examined the kinetic trapping of thin film morphologies during solvent removal.38 Ultraviolet−ozone (UVO) etching and subsequent AFM imaging were used to determine the mechanism for the reorientation of cylindrical domains in poly(deuterated styrene-b-isoprene-b-deuterated styrene) (dSIdS) BCP thin films annealed with chloroform vapor. A critical result of this work is illustrated in Figure 9, where our solvent-vapor-annealed films formed well-oriented parallel cylinders when swollen to a constant film thickness (near 50 vol % solvent), followed by instantaneous solvent removal, whereas the same films reoriented to perpendicular cylinders upon slow solvent removal. In the case of the perpendicular cylinders, the reorientation propagated from the free surface to the substrate surface with the slowest annealing times, showing complete permeation of the perpendicular orientation through the film thickness. This work helps resolve several seemingly contradictory reports in the literature, where similar SVA

Figure 8. AFM phase images of the SIS BCP annealed in the SVA microfluidic device (Figure 7). THF and n-hexane were the solvent vapors used. The chamber number (Figure 7) and the solvent-vapor composition (% n-hexane) are indicated. Adapted with permission from ref 37. Copyright 2011 American Chemical Society.

conditions led to different final morphologies (or orientations), likely as a result of the solvent removal process. Additionally, our imaging approach, using successive AFM/UVO etching, provided a relatively rapid (as opposed to cross-sectional TEM) and in-house (as opposed to synchrotron grazing-incidence SAXS [GISAXS]) approach to gain insights into the throughfilm effects of solvent-vapor removal on thin film BCP morphologies. 3.4. Raster Solvent-Vapor Annealing.49 Although nanostructure control of BCPs can be achieved through SVA, it is typically limited to small batch processes. Continuous and zone-based annealing methods for morphology control have been adapted from those used in other industries, such as metallurgy and semiconductor processing, but in BCP systems, these approaches typically have been limited to zone casting from solution50 or thermal zone annealing.51,52 Thus, we designed a raster solvent-vapor annealing (RSVA) method that offers the advantages of SVA described previously, as well as faster annealing times and precise spatial morphology control.49 Our RSVA setup used a nozzle to direct solvent vapor, which provided convective vapor transport to the films (as opposed to the diffusive transport in typical SVA experiments) and resulted in significantly reduced annealing times. Furthermore, our method generated a stylus-like point-annealing zone that could 3872


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Figure 9. (a) Thickness profiles for films annealed with chloroform vapor. Films were held at a constant swollen film thickness for 2 h, followed by solvent removal at different rates. (b) AFM phase image of the as-cast film at the free surface. (c) AFM phase images for the free and substrate surfaces at different solvent removal rates. AFM scale bars represent 400 nm. Adapted with permission from ref 38. Copyright 2012 American Chemical Society.

be rastered across a film using a motorized stage to enable localized control over nanoscale structure and orientation.49 In one example, RSVA of SIS thin films (∼100 nm in thickness) was performed by directing a THF-rich vapor stream through a nozzle onto the films (Figure 10a). As-cast, flowcoated films manifested parallel-oriented cylinders with minimal long-range order. As shown in Figure 10, this morphology could be altered by subjecting the films to RSVA at raster speeds ranging from 3 to 500 μm/s, controlled by a motorized stage. Using this setup, we generated parallel cylinders at faster raster speeds (i.e., shorter annealing times per area) and perpendicular cylinder surface structures at the slowest raster speeds (i.e., longest annealing times per area).

The orientation rearrangement to perpendicular cylinders appeared to occur via the gradual breakup of parallel cylinders with longer annealing times, and sufficiently long solvent-vapor annealing times induced this parallel to perpendicular orientation shift in several ways. The driving forces for rearrangement included the decreased surface energy between the blocks and the swelling and deswelling of the polymer thin film, which required the film to cycle through film thicknesses both commensurate and incommensurate with the polymer domain spacing. (The entropic penalty due to chain stretching is better accommodated by perpendicular orientations.) Additionally, the slower raster speeds likely resulted in slower solvent removal rates, which may have induced the 3873


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materials in a high-throughput fashion. Specifically, we employed thermal gradients to examine PS-b-poly(t-butyl acrylate) (PS-PtBA) BCPs, where PtBA undergoes thermal deprotection to generate PS-b-poly(acrylic anhydride) (PSPAH). (PS-PAH is the dehydrated analogue of PS-b-poly(acrylic acid) (PS-PAA).39) Upon deprotection (from PtBA to PAA/PAH), the BCP experienced a mass loss resulting in a decrease in the responsive-block volume fraction. In a parallel lamellar-forming PS-PtBA, this volume fraction change induced a morphological change from the parent lamellar morphology to cylindrical nanostructures. Furthermore, an additional result of the deprotection was a change in the relative block surface energies, which favored a perpendicular nanostructure orientation at the free surface of the PS-PAH films. We further manipulated interfacial interactions in the same polymer system using a combination of SVA and thermal gradients.41 Because the complexity of the deprotection reaction may limit the complete rearrangement of PS at the air interface, SVA processing was used to increase the chain mobility. Morphologies were compared between films as-cast after SVA only, after thermal deprotection only, and after both SVA and thermal deprotection, as shown in Figure 11. The order of the steps was found to have an effect on the final film morphology. SVA on the as-cast films (Figure 11a) using THF or acetone (solvents that are essentially nonpreferential for PS or PtBA) decreased the relative interaction parameters, resulting in phase mixing (Figure 11b). Although the perpendicularly oriented cylindrical morphologies resulting from SVA followed by thermal annealing (Figure 11c) were unaffected by SVA history, films that were thermally deprotected first and subsequently subjected to SVA (Figure 11d) reoriented to parallel cylinders. This rearrangement occurred because THF and acetone were selective for PS over PAH in the deprotected system. We also attempted to hydrolyze the thermally deprotected PAH samples to PAA by exposing the films to water. In the perpendicularly oriented samples, the hydrolysis resulted in swollen mushroom-like tops (Figure 11e), but the parallel cylinders showed no noticeable changes (Figure 11f), suggesting that PAH was either not hydrolyzed to PAA or was too confined to result in swollen structures. Using the gradient methodologies and annealing procedures that we developed previously, we demonstrated the ability to manipulate the parameters governing nanostructure and domain orientation in thermally responsive BCPs and control the self-assembly of these materials in thin film geometries. 3.6. Controlling Particle Location in BCP Thin Films by Tuning Interfacial Energetics.53 In the previous sections, we demonstrated our ability to control the self-assembly of BCP thin films by manipulating the confining environment. However, non-polymer additives also may be manipulated to promote incorporation into specific BCP domains by modifying the interactions between the additives and the polymers.54−56 For example, the segregation of nanoparticles within BCP domains provides a facile method for generating functional nanocomposites for applications including optics and electronics.54,56,57 In one example of designer nanoparticle localization, we controlled the placement of gold nanoparticles (AuNPs) within SIS thin films by manipulating the surface chemistry of the AuNPs with relatively short thiols as opposed to long polymer brushes. By employing the short thiols, we obviated the need for additional polymer synthesis and attachment, which, in addition to being tedious, can impact

Figure 10. (a) Schematic of RSVA apparatus. BCP thin films are zone annealed with directed solvent vapor (left) while being moved on a motorized stage (right) to control the annealing area. (b) Example of film subjected to crossed-path RSVA. (i) The optical micrograph shows the precise spatial control provided by RSVA. (Note that the UD is “stained” to enhance the visual contrast.) Representative AFM micrographs of (ii) single-pass RSVA and (iii) double-pass RSVA show parallel cylinders and perpendicular cylinders, respectively. The scale bars correspond to 1 mm in the optical micrograph and 200 nm in the AFM micrographs. Reproduced with permission from ref 49. Copyright 2012 American Chemical Society.

reorientation from parallel to perpendicular cylinders as discussed in section 3.3.38 Furthermore, grain sizes for the parallel cylinder specimens were tuned using intermediate speeds (intermediate relative to the slowest speeds that induced reorientation to perpendicular cylinders), with larger grains achieved at slower speeds. The precise spatial control afforded by the RSVA apparatus also allowed us to induce orientation shifts in specific areas by utilizing crossed-path RSVA (Figure 10b).49 The resulting films contained distinct regions of either parallel or perpendicular orientations as a result of single-pass RSVA (shorter total annealing times per area (parallel)) or double-pass RSVA (longer total annealing times per area (perpendicular)). Hence, this rapid RSVA method, with its precise position control, allowed us to “write” specific nanoscale structure and orientation on BCP thin films.49 3.5. Nanostructure Manipulation in Thermally Responsive BCP Thin Films.39,41 We used the experience gained from our substrate surface and SVA studies to manipulate the nanoscale morphologies in thermally responsive 3874


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Figure 11. AFM phase images and schematic diagrams of potential chemical structures, morphology, and orientations resulting from combinations of solvent-vapor annealing (SVA), thermal annealing (TA), and hydrolysis with water (H2O). Arrows labeled with TA, SVA, and H2O represent the processing order starting from (a) the as-cast lamellar film of PS-PtBA. The green domains represent PS in all panels, the purple domains represent PtBA in panels a and b, and the pink domains represent the deprotected PAH in panels c and d or its hydrolyzed analogue, PAA, in panel e. It is unclear if PAH is hydrolyzed to PAA in panel f. Adapted with permission from ref 41. Copyright 2011 Wiley Periodicals, Inc.

each block. Moreover, like thin film BCPs, solution structures also are restricted by interactions with their external environment, in this case, the solution.62 In aqueous solution, selfassembly can be manipulated by controlling the solvent quality and interfacial free energy between the hydrophobic core and the hydrophilic (and normally swollen) corona. 4.1. Modified Substrates for High-Contrast Imaging of BCPs: Graphene Oxide.63 Because of its high resolution compared to that of scanning electron microscopy (SEM) and AFM, TEM is a desirable microscopy technique for imaging the nanoscale features of BCP solution assemblies. Although cryogenic-TEM (cryo-TEM) enables the imaging of these systems in their “natural state” within vitrified films of the solutions, there is limited access to the necessary sample preparation and imaging equipment. Thus, samples in which the solution assemblies are dried on a support are still used widely, and the technique can provide meaningful information when taken in context. However, an additional complication of the TEM imaging of BCPs is that heavy metal staining often is required to improve the contrast between polymer blocks and between the polymers and the carbon support. Because artifacts from staining can lead to unreliable results,64 it is of interest to find supports that allow for high specimen/support contrast in the absence of stains. In a growing collaborative effort with the O’Reilly group, we demonstrated improved specimen/substrate contrast by using graphene oxide (GO) supports that are easily prepared and nearly electron-transparent.63 TEM images of solutionassembled specimens prepared on GO-coated grids showed clear resolution of the particle nanostructures. Additionally, the use of GO-coated grids allowed for the direct comparison between TEM, which provides 2D projections of 3D objects, and AFM, which provides complementary 2D and 3D information about 3D structures. Typically, these analyses are performed on samples prepared on different substrates; however, the strength and stability of our GO supports permitted the analysis of the same specimen by multiple imaging techniques (TEM, AFM, and SEM).63 In this work, we compared images of the same individual poly(acrylic acid-bstyrene) (PAA-PS) polymersomes prepared on a GO support on a “finder” TEM grid (Figure 12). Through our analyses, we

transport and material properties in the case of conducting or catalytic nanocomposites. In our work, AuNPs were fabricated with 1-dodecanethiol (C12SH) capping ligands and subsequently ligand exchanged with short thiol-functionalized polystyrene (PS-SH) oligomers. The as-synthesized, C12SHfunctionalized AuNPs preferentially segregated into the PI domains. After ligand exchange, AuNPs with a C12SH:PS-SH molar ratio of 5:1 or lower preferentially segregated into the PS domains. This relatively high ligand ratio for PS-preferential AuNPs was explained by size differences between the ligands because the PS-SH ligand occupied a larger area at the matrix/ ligand interface, resulting in a higher perceived PS-SH surface coverage. Incorporating a mixture of PI-preferential and PSpreferential AuNPs resulted in well-mixed dispersions of AuNPs within the SIS thin film. Thus, by tuning the surface chemistry of AuNPs using relatively short thiol-capping ligands and oligomers, we were able to modify the polymer− nanoparticle interfacial interactions in order to direct the placement of nanoparticles within select BCP thin film domains.

4. SOLUTION Although modifications to block junctions, χeff, and surface interactions affect bulk and thin film BCP behavior as discussed previously, interfacial manipulations can play a distinct role in the assembly and analysis of amphiphilic BCPs in solution environments. Amphiphilic BCPs consisting of a hydrophobic block and a hydrophilic block can self-assemble into a variety of nanostructures in aqueous solutions, including spherical micelles, worm-like micelles, and vesicles.58,59 These solution assemblies are promising candidates for drug delivery applications because they are synthetically versatile, are relatively stable, and have the ability to encapsulate therapeutic molecules.58 Additionally, targeted BCP delivery systems may increase the efficacy and reduce potential side effects of therapeutic agents.60 However, many of the methods used to create such targeted delivery systems involve functionalizing polymer chains, which often disrupts self-assembly.61 As in other BCP systems, the morphological behavior of BCPs in solution is determined by the molecular weight, relative block weight fractions, and chemical composition of 3875


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Figure 12. (a) TEM, (b) SEM, and (c) AFM images of PAA-PS polymersomes prepared on a GO support. Reproduced from ref 63 with the permission of The Royal Society of Chemistry.

Figure 13. (a) Cryo-TEM micrographs of PB-PEO micelles in H2O/THF mixtures at 0, 5, 15, 20, 30, and 40 vol% THF. PB cores appear dark, but the PEO coronas are not visible because of the low electron density contrast between PEO and water. The TEM scale bar represents 100 nm. (b) Illustration representing structural changes in PB-PEO micelles with increasing THF content. Adapted from ref 65 with the permission of The Royal Society of Chemistry.

% THF at constant polymer concentration).65 The detailed effects of this interfacial tension change on the micelle nanostructures were investigated using contrast-variation experiments in SANS. A combination of PB-PEO and deuterated PB-dPEO BCPs was employed to glean information about the micelle core and corona density profiles. SANS data were fit with a form factor model for polydisperse spherical micelles69,70 while also accounting for the presence of free chains in solution at higher THF contents (lower interfacial tensions lead to higher critical micelle concentrations).65 Our analysis indicated that the reduction in interfacial tension led to a broadened core−corona interface, effectively increasing the PB-core solvent accessibility. Additionally, micelle size changes induced by the addition of THF were consistent with theoretical scaling relationships, where the micelle sizes decreased with the lowering of the interfacial tension (Figure 13). Through a combination of techniques, we were able to elucidate key trends in the detailed micelle nanostructures that can have major impacts on the stability and accessibility of functional groups in solution assemblies for drug delivery and other applications.

were able to gain greater insight into the 3D nanostructure of polymersomes upon drying. In this particular example, the measured height from AFM slightly differed from the measured width from TEM, suggesting drying deformations.63 Thus, this relatively simple and accessible sample preparation technique provided a means to obtain accurate structural information about solution-assembled samples without staining, and it indicated some of the pitfalls associated with quantifying solution assemblies from TEM alone. 4.2. Manipulation of Interfacial Energetics in Solution Assemblies: Cosolvent Mixtures.65 Though BCP solution assembly can be controlled by synthesizing polymers of various molecular weights and hydrophobic contents, ultimately it is desirable to tune nanoscale assemblies using a single amphiphilic copolymer. We have approached this task using the manipulation of the interfacial energetics between the hydrophobic core and hydrophilic corona blocks by changing the quality of the solvent for the core block.62,66−68 In one example, we tuned the solvent accessibility, corona thickness, and micelle size of a single parent block copolymer by adding an organic cosolvent, THF, to aqueous solutions of poly(1,2butadiene-b-ethylene oxide) (PB-PEO) micelles.65 We found that the PB/PEO interfacial tension decreased by approximately an order of magnitude with increasing THF content (i.e., the solution conditions changed from pure water to 70 vol

5. SUMMARY We have utilized a number of synthetic and non-synthetic methodologies to manipulate the energetics of nanoscale self3876


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styrene-b-ethylene oxide) Triblock Copolymers. Macromolecules 2002, 35, 7007−7017. (11) Epps, T. H., III; Chatterjee, J.; Bates, F. S. Phase Transformations Involving Network Phases in ISO Triblock CopolymerHomopolymer Blends. Macromolecules 2005, 38, 8775−8784. (12) Tyler, C. A.; Qin, J.; Bates, F. S.; Morse, D. C. SCFT Study of Nonfrustrated ABC Triblock Copolymer Melts. Macromolecules 2007, 40, 4654−4668. (13) Tureau, M. S.; Rong, L. X.; Hsiao, B. S.; Epps, T. H., III. Phase Behavior of Neat Triblock Copolymers and Copolymer/Homopolymer Blends Near Network Phase Windows. Macromolecules 2010, 43, 9039−9048. (14) Fetters, L. J.; Lohse, D. J.; Richter, D.; Witten, T. A.; Zirkel, A. Connection between Polymer Molecular Weight, Density, Chain Dimensions, and Melt Viscoelastic Properties. Macromolecules 1994, 27, 4639−4647. (15) Phillip, W. A.; Dorin, R. M.; Werner, J.; Hoek, E. M. V.; Wiesner, U.; Elimelech, M. Tuning Structure and Properties of Graded Triblock Terpolymer-Based Mesoporous and Hybrid Films. Nano Lett. 2011, 11, 2892−2900. (16) Kuan, W. F.; Roy, R.; Rong, L. X.; Hsiao, B. S.; Epps, T. H., III Design and Synthesis of Network-Forming Triblock Copolymers Using Tapered Block Interfaces. ACS Macro Lett. 2012, 1, 519−523. (17) Nagpal, U.; Detcheverry, F. A.; Nealey, P. F.; de Pablo, J. J. Morphologies of Linear Triblock Copolymers from Monte Carlo Simulations. Macromolecules 2011, 44, 5490−5497. (18) Epps, T. H., III; Bates, F. S. Effect of Molecular Weight on Network Formation in Linear ABC Triblock Copolymers. Macromolecules 2006, 39, 2676−2682. (19) Tureau, M. S.; Epps, T. H., III Effect of Partial Hydrogenation on the Phase Behavior of Poly(isoprene-b-styrene-b-methyl methacrylate) Triblock Copolymers. Macromolecules 2012, 45, 8347−8355. (20) Young, W. S.; Epps, T. H., II.I Salt Doping in PEO-Containing Block Copolymers: Counterion and Concentration Effects. Macromolecules 2009, 42, 2672−2678. (21) Young, W. S.; Brigandi, P. J.; Epps, T. H., III. CrystallizationInduced Lamellar-to-Lamellar Thermal Transition in Salt-Containing Block Copolymer Electrolytes. Macromolecules 2008, 41, 6276−6279. (22) Young, W. S.; Albert, J. N. L.; Schantz, A. B.; Epps, T. H., III. Mixed-Salt Effects on the Ionic Conductivity of Lithium-Doped PEOContaining Block Copolymers. Macromolecules 2011, 44, 8116−8123. (23) Young, W. S.; Epps, T. H., III. Ionic Conductivities of Block Copolymer Electrolytes with Various Conducting Pathways: Sample Preparation and Processing Considerations. Macromolecules 2012, 45, 4689−4697. (24) Panday, A.; Mullin, S.; Gomez, E. D.; Wanakule, N.; Chen, V. L.; Hexemer, A.; Pople, J.; Balsara, N. P. Effect of Molecular Weight and Salt Concentration on Conductivity of Block Copolymer Electrolytes. Macromolecules 2009, 42, 4632−4637. (25) Meyer, W. H. Polymer Electrolytes for Lithium-Ion Batteries. Adv. Mater. 1998, 10, 439−448. (26) Semenov, A. N. Microphase Separation in Diblock Copolymer Melts - Ordering of Micelles. Macromolecules 1989, 22, 2849−2851. (27) Albert, J. N. L.; Epps, T. H., III. Self-Assembly of Block Copolymer Thin Films. Mater. Today 2010, 13, 24−33. (28) Segalman, R. A. Patterning with Block Copolymer Thin Films. Mater. Sci. Eng., R 2005, 48, 191−226. (29) Stafford, C. M.; Roskov, K. E.; Epps, T. H., III; Fasolka, M. J. Generating Thickness Gradients of Thin Polymer Films via Flow Coating. Rev. Sci. Instrum. 2006, 77, 023908. (30) Roskov, K. E.; Epps, T. H., III; Berry, B. C.; Hudson, S. D.; Tureau, M. S.; Fasolka, M. J. Preparation of Combinatorial Arrays of Polymer Thin Films for Transmission Electron Microscopy Analysis. J. Comb. Chem. 2008, 10, 966−973. (31) Albert, J. N. L.; Baney, M. J.; Stafford, C. M.; Kelly, J. Y.; Epps, T. H., III. Generation of Monolayer Gradients in Surface Energy and Surface Chemistry for Block Copolymer Thin Film Studies. ACS Nano 2009, 3, 3977−3986.

assembly in polymeric materials for diverse applications including lithium battery and separation membranes, nanoscale templates and surface coatings, and nanoscale delivery precursor vehicles for targeted drug delivery. Our key advances include the ability to manipulate transition temperatures and nanoscale network formation in bulk materials, the generation of facile methodologies for manipulating and screening thin film phase behavior, and the ability to tune and comprehensively characterize micelle sizes and density profiles using a single amphiphilic BCP. All of these advances are connected by the theme of the manipulation of interfacial interactions, either between domain interfaces or between assembled nanostructures and confining environments. These powerful methodologies that combine designer synthesis with detailed nanostructure characterization allow nearly unlimited access to a wide array of nanostructures in numerous BCP systems, thus providing the means to synthesize materials for utilization in the next generation of structure-dependent nanotechnologies.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The work highlighted herein was funded by the sources described in the acknowledgements sections of the referenced articles. In addition, we also acknowledge a National Science Foundation grant, CBET-0930986, for the financial support of S.E.M. during the writing of this article.

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REFERENCES

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