50th Anniversary Perspective: Putting the Squeeze on Polymers: A

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50th Anniversary Perspective: Putting the Squeeze on Polymers: A Perspective on Polymer Thin Films and Interfaces Thomas P. Russell*,†,‡,§ and Yu Chai‡ †

Polymer Science and Engineering Department, University of Massachusetts Amherst, Amherst, Massachusetts 01003, United States Materials Sciences Division, Lawrence Berkeley National Laboratory, 1 Cyclotron Road, Berkeley, California 94720, United States § Beijing Advanced Innovation Center for Soft Matter Science and Engineering, Beijing University of Chemical Technology, Beijing 100029, China ‡

ABSTRACT: Polymeric materials, used predominantly in their bulk form, are now finding increasing application as thin films or where the end-use dimensions are approaching the size of the individual polymer molecules. This introduces questions concerning the perturbations to the fundamental static and dynamic properties arising when the pervaded volume of the polymer must conform to the finite size limitations. These properties include the single chain dimensions, entanglement molecular weight, anisotropy of polymer molecules, and wetting. For block copolymers, where the disparity in the chemical nature of the constituent blocks gives rise to microphase-separated morphologies, confinement can introduce fundamental changes in the ordering transitions of the copolymer, changes in the fundamental length scales and morphologies typically seen in the bulk, and changes in the orientation of the microdomains. As the size scale decreases, the surface-to-volume ratio increases, and as such, interfacial interactions and surface energies become increasingly important. The dramatic density at a surface and preferential interactions at surfaces and interfaces can have a profound impact on the chain dynamics and in multicomponent systems and block copolymers the preferential segregations of one component to the surface or interface. While the orientation of morphologies relative to the interface will occur, the extent to which surface effects propagate into the film and the adhesion of one material to another are important to understand for end-use applications. For interfaces between disparate materials, as for example between an organic polymer and a metal, transport of electrons or holes from one medium to the next, where there is a disparity in the work function, reduces to the behavior and nature of the polymer immediately at the interface. Here we summarize some of the significant advances that have been made on polymer thin films and interfaces, where the polymer chains can be squeezed, and provide some insight into future and emerging areas in the field.

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the film thickness approaches a radius of gyration, Rg, the Laplace pressure drives the surface of the film to remain flat, so the dimensions of the polymer chains normal to the surface must be compressed. But, what happens to the dimensions of the polymer chain in the plane of the film? There must be an anisotropy in the configuration of the polymer chain that will lead to anisotropic dielectric and mechanical properties that are different in and out of the plane of the film. At the surface of the film, where the density undergoes a change from the bulk value to zero, what ramifications does this have on chain configuration and chain dynamics, and how far do these surface effects propagate into the film? If the system of interest is a polymer blend or a block copolymer, the chemical difference between the constituents translates into differences in the surface and interfacial energies that lead to the preferential segregation of one component to the substrate and, by default,

INTRODUCTION Since their discovery in the past century, the use of polymers has become pervasive in our everyday lives. For the most part, applications have focused on the use of polymers where the bulk characteristics of the macromolecules have been of primary interest. However, the explosive growth of microelectronics has brought attention to the use of polymeric materials in thin films, as for example dielectric insulators in chips, alignment layers in flat panel displays, hole transport layers in photovoltaic devices, and coatings in and for devices ranging from touch screens to encapsulants to no-stick or wearresistant coatings. As the thickness of the films decreases, the surface and interfacial behavior of the polymer chains becomes increasingly important. Whether interfacial interactions induce a well-defined orientation to the polymer chain at an interface by specific interactions with segments of the polymer or promotes a specific orientation of a morphology, i.e., a collective ensemble of polymer chains, the long chain nature of the molecules can result in the propagation of surface effects into the film that could dominate the properties of the film. As © XXXX American Chemical Society

Received: February 24, 2017 Revised: April 28, 2017

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question that small-angle neutron scattering (SANS) has quantitatively shown that the polymer chain assumes an unperturbed configuration in the bulk, for thin films the results in the literature using SANS are contradictory. The first smallangle neutron scattering measurements on thin films10,11 indicated that the lateral dimensions of the chains, i.e., the Rg of the polymer in the plane of the film, remain constant, as shown in the upper part of Figure 1. In the opposite extreme,

an orientation of the morphology that will dramatically impact the end-use of these materials. The film thickness also introduces confinement effects on the morphology that can result in changes to the fundamental lengths scales of the morphology or even changes in the phase behavior, e.g., the order-to-disorder transition (ODT) in block copolymers or the lower critical solution temperature in polymer blends. As the degrees of confinement increase, totally new morphologies can emerge along with changes in the chain dynamics. Significant advances have been made in our understanding of thin polymer films and the behavior of polymers at surfaces and interfaces due, in large part, to major developments in instrumentation that can probe polymeric systems on the segmental level. There have been tremendous advances in scanning probe microscopies, electron microscopies, neutron and X-ray reflectivity, hard and soft X-ray scattering at synchrotron and free electron light sources, optical and confocal microscopy, and spectroscopies. These instrumentation developments have revealed structural, morphological, and spectroscopic details on the behavior of polymers at interfaces and in thin films over a broad range of length scales, from the segmental to mesoscale levels. Paralleling these are advances in simulation and computations that have provided even more detail. It is appropriate, therefore, to take stock of where we stand in our understanding of the behavior of polymers at surfaces and interfaces and in thin films and how interfacial interactions have played a key role. More importantly, we need to ask what lays ahead, what questions are still open, and what potential instrumental developments can further our understanding.

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HOMPOLYMERS In the bulk, polymer chains pervade a spatial volume that is defined by the radius of gyration, Rg, of the polymer.1 However, in thin films, as the thickness of the film approaches 2Rg, the polymer chains must compress normal to the film surface. Yet, what happens to the chain dimensions in the plane of the film is still an open question. Will the chains deform biaxially to preserve the pervaded volume of the chain, or will the chains retain the bulk Rg? The pervaded volume relative to the occupied volume is related to the entanglement molecular weight, which is of consequence to the mechanical and rheological properties of the polymer thin films. Will the constraints imposed by the film thickness cause changes in the local configuration of the polymer? If the chains biaxially stretch, will the chains become more rigid or will they retain the characteristics of the flexible chain in the bulk? This, of course, is important in defining the moduli of the thin films. These different properties are important, ultimately, in the end use of thin polymers films across a broad range of applications, including thin film substrates, insulating layers, and encapsulation layers. Computer simulations of polymer melts between impenetrable walls indicate that chain dimensions parallel to the surfaces are only slightly larger than in the bulk and that the chain conformation in this direction remains Gaussian.2−6 In contrast to these ideas, self-diffusion measurements for confined polymeric systems are consistent with the notion that chain conformation is strongly modifed.7,8 Others9 have suggested that chain structure can be significantly affected in ultrathin films (that is, thickness D < 100 nm). Neutron scattering is the logical method to determine the single chain configuration in the bulk and in thin films. While there is no

Figure 1. (A) Radius of gyration as a function of film thickness (in nm) for PS with an Rg of ∼10 nm. Reproduced with with permission from ref 10. Copyright 1999 Macmillan Publishers Limited. (B) Kratky plots of two PS films (fill circles are bulk data and open circles confined by film thickness). Reproduced with permission from ref 12.

another study,12 again using SANS on films spin coated onto silicon substrates, indicated that the Rg of the polymer nearly doubles in the plane of the film and the persistence length, the crossover between coil-like and rod-like behavior, has shifted to smaller scattering vectors or has increased significantly, indicating that the chain has become more rigid. Quantitatively addressing these differences will require neutron sources that have higher flux and are more brilliant. Thermally driven collective dynamics are important for many macroscopic properties of polymers.13−19 However, our understanding of these dynamics, which are cooperative and heterogeneous, remains limited.20−22 Dynamic heterogeneities are accompanied by a broad spectrum of temporal fluctuations, but this spectrum has not been studied at short length scales where the behavior must break down. The structural, dynamic, and materials properties must deviate from their bulk values B

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There has been a tremendous effort to study the Tg reduction in thin polymer films and whether it is caused by the presence of a liquid-like layer at the surface of the polymer. A reduction in the Tg of organic liquids confined in small pores preceded the observations on polymers, where Tg, measured by differential scanning calorimetry, was found to be a function of pore size.28 This suggests that Tg reduction is a general phenomenon of amorphous materials, occurring when the volume at the surface of the polymer or at the interface with another material constitutes a significant fraction of the total volume of the material. By using physical vapor deposition (PVD), for both organic29 and polymer30 glasses researchers were able to form glassy materials with higher stability (stable glasses). Stable glasses could be prepared by the PVD technique with the temperature of the deposited substrate lower than the bulk Tg. This suggests that, with a higher mobility of the free surface, the deposited layer can rearrange and form a stable structure before the next layer was deposited. If there is a mobile surface layer, an increase in the free surface area should lead a stronger Tg reduction. Brillouin light scattering on free-standing films confirmed this hypothesis.31 It would stand to reason that removing the free surface should remove this effect. Ellipsometric measurements of the Tg of thin films covered with a metal layer indicated that the bulk Tg value was obtained in films as thin as 7 nm, while a 11 K reduction of Tg was found for a film of the same thickness without the capping metal layer.32 Similar types of results were found for free-standing films transferred to a solid substrate where a significant Tg reduction measured by ellipsometry in the freestanding films was removed after transferring the films to a substrate.33 Time-dependent power spectral density measurements with atomic force microscopy on lower molecular weight polystyrene films showed a lower viscosity of thin films in comparison to the bulk.34 The reduction of viscosity was explained by a two layer-model, where there was a thin mobile layer with a thickness less than 2.3 nm. Arrhenius-type dynamics was found rather than typical Vogel−Fulcher− Tammann dynamics. Atomic force microscopy measurements on the relaxation of nanoindentations at the surface of glassy polystyrene films indicated that the surface relaxed much more rapidly than the bulk and the surface relaxation had a weaker temperature dependence in comparison to the bulk relaxation.35 To minimize surface energy, rough surfaces are disfavored and flat surfaces are thermodynamically more stable; i.e., the surface energy is minimized. The relaxation of artificially made steps in polystyrene was measured above and below Tg by atomic force microscopy.36 The planarization of the film surface above Tg was explained by the flow of the entire film, while below Tg only the thin layer near the free surface could flow. X-ray photon correlation spectroscopy (XPCS) measurement on the surface dynamics of polymer films at temperatures well above Tg showed that the dynamics at the surface are the same as those measured in bulk polystyrene rheologically.37,38 Entropy favors the segregation of polymer chain ends to the surface, as found experimentally39,40 and by simulations.41,42 However, the magnitude of the Tg reduction cannot be simply explained by chain-end effects. There is a second school of thought where a reduction in Tg at free surfaces was not observed. Dielectric spectroscopy measurements on ultrathin films,43−45 studies on polymeric nanoparticles,46 surface relaxation measurements by near edge X-ray absorption fine structure analysis on deformed surfaces,47,48 and scanning probe microscopy measurements

when the system is confined below a characteristic length scale, whether this should be Rg or the persistence length, at the interface between two materials or at the free surface. Ionconductive polymers are increasingly applied as nanometerthick thin films and at interfaces with other materials to form multifunctional composite structures. Inorganic particles are frequently incorporated into porous polymer electrodes to facilitate reactions and provide additional functionalities, as seen with carbon/platinum nanoparticles in fuel-cell electrodes or silicon in vapor-fed solar-fuel devices. Polymer thin films are frequently used as capacitors or dielectric insulators, as resists for lithography, and as lubrication layers. In all of these applications, the orientation and the dynamics of the polymers in confined geometries or at interfaces are important. Yet, obtaining information on the configuration of the polymer chain, the average orientation of the polymer, and the dynamics of the polymers in these confining geometries and at interfaces has been difficult. Some advances have been made, but we still have a long way to go before we get a complete understanding. An important physical property of bulk polymers is the glass transition temperature, Tg, where, upon cooling, the volume contraction of the polymer deviates from equilibrium.23−25 However, as the film thickness decreases, approaching the dimensions of a single polymer chain, or at an interface, where there is an abrupt change in the density, does Tg, a fundamental reflection of chain segmental dynamics, change? This seemingly simple question has drawn a significant amount of attention from laboratories worldwide. Initial studies on the thickness dependence of Tg measured by ellipsometry (shown in Figure 2) indicated that, when the film thickness was less than 40 nm,

Figure 2. Glass transition temperature of polystyrene (120 000 (triangles), 500 800 (circles), and 2 900 000 (diamonds)) on a silicon substrate as a function of the film thickness. Reproduced with permission from ref 26. Copyright 1994 European Physical Society.

Tg decreased with decreasing film thickness.26 In addition, no molecular weight dependence of the Tg reduction was observed, leading to the conclusion that the Tg reduction was not caused by the geometric confinement of the chains but, rather, that there was a liquid-like layer at the surface of the glassy polymer that can affect the apparent Tg of the thin film. According to the Williams−Landel−Ferry (WLF) equation, the relaxation time of polymers changes dramatically when the temperature approaches Tg. Taking polystyrene as an example, a 10 K decrease of Tg implies that the relaxation time decreases by a factor of 1000.27 Such a large change of relaxation time must affect how we can use polymeric materials, particularly in nanotechnology applications. C

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Macromolecules on the surface roughness as a function of temperature49 did not show a significant reduction in the Tg or enhanced surface mobility. This raises questions concerning the sample history,50,51 the precise information that one is measuring with a given probe technique, the total volume and dynamics with the different time scales52 being probed by a given technique, and the depth or length scale over which surface effects propagate into the bulk of the material. As the film thickness decreases, the interactions of the polymer with the underlying substrate must be taken into account. Similarly, in polymer composites, the interfacial interactions of the polymer with the solid additive and the concentration of the additive must be considered. For systems where the interactions are weak, as for example poly(methyl methacrylate), PMMA, on gold, Tg was found to decrease with decreasing film thickness.26 In composite materials, the dynamics and thermodynamics of the matrix polymer were found to be significantly influenced by the presence of nanoparticles.53 When the interactions are strong, as for example PMMA on a silicon substrate, Tg was found to be markedly increased.26 A more comprehensive study of the relation between Tg and polymer−substrate interfacial energy was conducted by X-ray reflectivity and ellipsometry.54 Tg was found to vary almost linearly as a function of the interfacial energy, irrespective of the chemical composition of polymers. However, as evidenced by the behavior of confined colloidal polystyrene dispersed in different solvents,55 our understanding of the role of interfacial interactions on dynamics is far from being quantitatively understood. Because of the strong correlation between dynamics and T − Tg, any impact on Tg should also be manifest in in the mobility of the polymer. Nanoindentation recovery experiments of isotactic PMMA films supported on different substrates silicon and aluminumindicated that the surface relaxation process was significantly influence by the substrate.56 With strong interfacial interactions, mobility was slowed, while with weak interactions, the mobility was increased. These effects were observed even for films as thick as 180 nm, many times the pervaded volume of the polymer, reflecting the cooperativity of motions in polymeric materials. Fluorescence measurements of a dye-labeled layer in a PS film57−59 and neutron reflectivity measurements with a deuterium-layer labeled layer of PS,60 both located at different distances from the interfaces, showed a similar result. Markedly enhanced diffusion of PS in the pores of an alumina membrane, a weakly interacting system, was also found, showing a significant increase in the diffusion of the PS in the pores.61 The physical aging rate, measured by the fluorescence of dye molecules, showed that the relaxation was reduced by a factor of 15 at the polymer−silica substrate and the relaxation process was nearly arrested.62 Taking confinement and interfacial interactions to an extreme, the Tg of PS nanoparticles and PS nanoparticles coated with a shell of silicon oxide was measured by differential scanning calorimetry.63 As the diameter of polystyrene nanoparticles decreased, the PS nanoparticles showed a pronounced reduction in the Tg, while the silicon oxide-coated PS nanoparticles remained at the bulk value of Tg. A physical aging experiment with the same materials showed a pronounced physical aging rate for the bare PS nanoparticles, while the silicon oxide-coated nanoparticles showed a reduced aging rate.64 So, while there has been a tremendous amount of attention paid to role of interfaces and confinement effects on the

dynamics of polymers, an absolutely definitive answer has not yet emerged. There is no question that confinement and interfacial interactions have an impact, but a quantitative measure of the effect is still lacking. This, in part, may arise from the fact that it is difficult to obtain data on thin films where the polymer is brought from the molten state, where the polymer is at equilibrium, into the glassy state, where the polymer deviates from equilibrium. This introduces artifacts from the sample processing history. The lack of clarity may also arise from the fact that different techniques probe different characteristics of the polymer that may or may not respond in the same manner when the polymer is brought into a nonequilibrium state. It may also be inherent to the heterogeneity in the dynamics of glassy polymers where, as the films are made thinner or as the confinement volume decreases, the dynamics may become more uniform, echoing the sentiments of Feynman where uniformity of nanostructured materials (and he points to polymers) can only be achieved with disordered materials.65 A very basic ramification of interfacial interactions on polymers, particularly polymer thin films, is wetting.66−71 If the interaction energy, interfacial energy between a polymer and a substrate, multiplied by the interfacial area, plus the surface energy of the polymer over the same area, is greater than the surface energy of the substrate over the same area, the polymer film will wet or spread over the surface.72 If it is energetically more favorable to expose the substrate surface, the polymer will dewet the substrate. This will, of course, depend on the film thickness where, if the film is thick enough, dewetting will not occur but, as the film is made thinner, the film will become unstable, producing height fluctuations at the film surface that will grow with time, eventually exposing the substrate and nucleating the dewetting of the polymer film from the substrate. Dewetting studies have provided a very rich area to understand critical phenomena in thin films and, also, the rheological characteristics of thin, molten polymer films73−77 and even of glassy polymer films.78,79 Directly related to the wetting characteristics of polymer films is the adhesion of a polymer films to a substrate or to another polymer film.80,81 Obviously, if the interfacial interactions are highly nonfavorable, the adhesion will be poor and adhesive failure will occur at the interface. If, though, the interactions are favorable, as for example with the use of a surface modification layer or other adhesion promoter, failure will occur cohesively when the layers are pulled apart, where a crack will propagate within one of the layers when the layers are separated. The area of adhesion opened an interesting application for thin copolymer films, both random and block copolymers.82−84 By taking two immiscible polymers and placing an A−B block copolymer at the interface, where the thickness of this layer can be varied and each block is miscible with one of the polymers comprising the bilayer, then the adhesion between the two layers will increase, as the thickness of the interlayer is increased. This trend continues up to a thickness of L0/2, where L0 is the repeat period of the block copolymer. For interlayer films thicker than L0/2, the block copolymer will form a multilayered structure at the interface and the fracture strength will be limited by that of the block copolymer or the blocks with each layer, whichever is less.85 As the complexity of the chain topology increases, from multiblock copolymers to random copolymers, strengthening of the interface is also observed where the copolymer makes multiple traverses from one layer to the other, effectively stitching the layers together D

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and, like CPEs, are soluble in polar solvents and are emerging as an alternative to CPEs.108 The pendent zwitterions are dipoles, with the distinct advantage of having no mobile counterions and have been found to markedly alter the electrode work function, even for films as thin as 1 nm. For both CPEs and CPZs, the packing of the chains at the interface and the orientation of the backbone chains and side chain dipoles are critical where both the chemical structure and composition of the polymers and their processing are key to their performance. While CPEs and CPZs have provided a novel means to manipulate the electrode work function for OPVs, they perform a similar role for perovskite-based photovoltaic active layers. Significant promise for CPEs and CPZs has been demonstrated; further modification of the chemical structure of the main chain, uncovering new zwitterionic materials, and fine-tuning the morphology of the interlayer will markedly expand their potential impact.

by the trajectory of the copolymer and entanglement of the segments of the copolymer that loop into and entangle with each layer. Even if, as with a random copolymer or strictly alternating copolymer, the sequences of like segments are short, adhesion promotion is observed, since interactions between the polymers comprising the layers are mediated, the interface is broadened, and the entanglement of the polymers at the interface is enhanced.86,87 Yet, if the copolymer could be made responsive to an external trigger, as for example light, such that, upon exposure, the copolymer chain was degraded, then a simple release mechanism could be designed.88 Conversely, if upon exposure to the trigger the copolymer chain was chemically altered, such that one block was converted from being hydrophobic to hydrophilic or from being to polar to even more polar, then the adhesion between the layers could be tuned in a very simple manner. The promise of organic electronics is found in the combination of desirable properties: electronic functionality, lightweight, mechanical flexibility, and facile solution-processing over large areas.89−93 Interfaces play a critical role in achieving high efficiency in organic electronic devices. Charge transport across the interface between an electrode, generally a conductive inorganic material, and an organic active layer is the ability to exchange charge between conductive electrodes and organic active layers.94 An energy level mismatch presents energy barriers that that can prevent or severely hamper charge injection or extraction processes.95 Polymeric interlayers at the electrode interfaces are emerging as an effective route to overcome the energy level mismatch, leading to potential applications in organic field-effect transistors (OFETs),96 lightemitting diodes (OLEDs), 97 and photovoltaic devices (OPVs).98 For example, in OPVs, interlayers increase the built-in electric field (Vbi) across the active layer, ensuring efficient extraction of photogenerated charge carriers.99,100 Such interface engineering is manifest in placing an interlayer between the photoactive layer and a high work function cathode,101,102 which has enabled power conversion efficiencies (PCEs) to exceed 10% for single-junction devices103 (Figure 3).

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BLOCK COPOLYMERS The chemical bonding of two dissimilar polymers together introduces, by default, variations in the surface energies and interfacial interactions within a single chain. For the simplest case of an A−B diblock copolymer on a solid substrate, the preferential interactions of one block with the substrate will cause a segregation of that block to the substrate and the lower surface energy block to the free surface.109−112 The connectivity of the blocks will cause the microdomains to orient parallel to the substrate or the free surface, for cylindrical or lamellar microdomains,113 while for spherical microdomains,114 depending on film thickness, either body-centered cubic symmetry or hexagonally close-packed structures are observed parallel to the surface. This orientation will propagate into the film by amounts that depend on the strength of the interfacial interactions and the segmental interactions between the blocks. There have been numerous studies on thin films of microphaseseparated diblock copolymers where the orientation of the microdomains leads to a condition where the film thickness of the block copolymer films at any point on the sample is, in general, dictated by the period of the microdomain morphology. If one block segregates to both the substrate and air interfaces (symmetric wetting), the films thickness is given by nL0, where L0 is the period of the microdomain morphology in the bulk and n is an integer, while for asymmetric wetting, the film thickness is given by (n + 1/2)L0. If the initial film thickness does not satisfy these conditions, then the surface of the film is decorated with islands or holes, topographic features with a step height of L0, where the fraction of the surface covered with these features is governed by the fraction of L0 the film thickness is between successive values of n. Observation of this behavior was first made with reflection interference optical microscopy where, using white light, each discrete change in the film thickness resulted in a distinct change in the interference color of the film.109−112 At the same time this observation was made, developments in the use of neutron reflectivity,113 dynamic secondary ion mass spectrometry,101 and scanning probe microscopies115 on polymers had occurred that, when coupled with these preferential interactions to generate highly oriented morphologies on planar substrates, like silicon wafers, enabled the dissection of block copolymer morphologies, from the distribution of segments within the domains to details on the interface between the microdomains with sub-nanometer spatial resolution.

Figure 3. Schematic of an organic photovoltaic devices with the active layer on a PEDOT:PSS conducting polymer electrode and a conjugated polymer zwitterion interlayer at a silver cathode. Two conjugated polymers with sulfobetaine zwitterion side chains where the separation distance was varied with different alkane spacer lengths. Reproduced with permission from ref 160. Copyright 2014 AAAS.

Conjugated polyelectrolytes (CPEs), for example, are one class of polymers that are proving to be an effective interlayer material.97,100,104−106 Their charged side chains afford solubility in solvents that are orthogonal to solvents for the active layer, allowing sequential deposition. While effective, CPEs have the disadvantage of mobile counterions that can impact function;107 conjugated polymer zwitterions (CPZs), on the other hand, are charge neutral, hydrophilic, polymer semiconductors E

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segments, then, by varying x in the synthesis, the interfacial interactions can be changed from being strongly favorable to A or strongly favorable to B or equally favorable or unfavorable to A or B. In the final case, where interfacial interactions are balanced, preferential interactions a removed and both A and B blocks of the copolymer can be located at the interface. It should be noted that this advance in interfacial engineering occurred in parallel with advances in living free radical polymerization, where random copolymers could be prepared with well-defined molecular weights and narrow molecular weight distributions.118−120 Having removed the enthalpic penalty at the interface, the entropy of the block copolymer favors placing both blocks at the interface, resulting in an orientation of the microdomains normal to the interface.117 The exact conditions of neutrality will depend on the chemical nature of the segments and the volume fraction of the copolymer.117 It is, of course, not mandatory that the constituents of the random copolymer are identical to the segments of the block copolymer, only that the interactions of the blocks with this surface layer can be varied in a controlled manner. Water contact angle measurements are a simple, yet powerful method to determine the surface energy of the modified surfaces. To circumvent specific surface chemistries required to anchor the random copolymer chains to the substrate surface, cross-linking a thin random copolymer film on the substrate surface has been shown to be a very versatile and flexible route to control interfacial interactions.121 Since the copolymer will penetrate into this surface modification layer, it must be thick enough to prevent the block copolymer from interacting with the underlying substrate, or else a parallel orientation will result. While anchoring a random copolymer to the substrate can be used to control interactions with the substrate, at the opposite interface, interfacial interactions can be controlled in a similar manner. Evaporation of silicon oxide onto the surface of a copolymer film produces a solid surface with preferential interactions with one of the blocks and confines the copolymer between two hard surfaces.122 Alternatively, films of very high molecular weight Ax-r-B1−x can be placed on the surface to control interactions, where interdiffusion between the Ax-r-B1−x is minimal123 or an immiscible, high glass transition temperature polymer film can also be effective.124 Strategies have been developed for sacrificial surface layers to control interfacial interactions that can easily be removed subsequent to thermal annealing of the block copolymer films to control the orientation of the block copolymer microdomains on effectively free surfaces.125 Consequently, routes have been developed to manipulate interfacial interactions uniformly at both interfaces. While static control of interfacial interactions has been described, developing routes to control interfacial interactions actively, while the system is actively changing, represents an interesting challenge. For example, to generate block copolymer microdomains where the period is exceptionally small, i.e., on the several nanometer level, strong, nonfavorable interactions between the segments of the blocks are required, introducing large differences in the interfacial and surface energies. An alternative strategy is to produce uniform, phase mixed films of a block copolymer where each block is hydrophobic, on a surface modified with a random copolymer having the same chemical constituents as the bock copolymer. Chemically or by irradiation, one of the blocks can be converted into a hydrophilic block while the anchored random copolymer is also changed, continuously maintaining neutral interfacial

If the block copolymer is above the order-to-disorder transition, then the blocks are mixed with essentially a random distribution of segments of both blocks throughout a bulk film. However, at the substrate or free surface, preferential segregation of one of the blocks to these interfaces leads to a pinning of one of the blocks at the interfaces. This pins a concentration fluctuation at the interface that will decay to the bulk concentration with a decay rate of e−x/ξ, where x is the distance from the substrate interface and ξ is the characteristic decay length (shown in Figure 4). The magnitude of ξ will

Figure 4. (A) Neutron reflectivity measurements on a PS-b-PMMA symmetric diblock copolymer confined between two silicon oxide surfaces as the temperature is changed bring the bock copolymer into the ordered state. (B) Variation in the effective ODT as a function confinement thickness. Adapted with permission from ref 116. Copyright 1992 American Physical Society.

depend on how far the block copolymer is from the ODT. The closer to the ODT, the larger will be ξ. Neutron reflectivity studies on thin films have clearly shown that the preferential interfacial interactions at both interfaces can lead to a change in the ODT.116 Whether the ODT is increased or decreased will depend on the commensurability of the film thickness with the period and whether the two interfaces are hard or soft. As the complexity of the block copolymers increases to triblocks, multiblocks, star-shaped, and miktoarms, the varieties and complexities of the morphologies will also increase. Yet, the preferential interactions at the interfaces will, ultimately, dictate the orientation of the morphology and phase behavior at the interface, and the extent to which this orientation propagates into the film will depend on the energetic cost, both enthalpic and entropic, to introduce defects that will negate interfacial effects leading to bulk morphologies, characterized by randomly oriented grains of the microdomain morphology. Controlling the interfacial interactions opens the possibility to control the orientation of the microdomain morphology. Perhaps the simplest strategy by which the interfacial interactions can be varied in a controlled manner is by anchoring random copolymers to the interfaces where the chemical constituents of the random copolymer are identical to those in the block copolymers.117 In the case of A-b-B diblock copolymers, by anchoring end-functionalized random copolymers, Ax-r-B1−x, to the substrate, where x is the fraction of A F

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Figure 5. SFM of PS±PVP film on top of a mesa. When the mesa edges are less than 5 μm apart, a single crystal is formed. Reproduced with permission from ref 126. Copyright 2001 Wiley-VCH.

Figure 6. Films of asymmetric PS-b-PEO having cylindrical microdomains solvent annealed on a faceted sapphire substrate (AFM height images shown in A and B) where the molecular weight of the PS-b-PEO was varied (shown in insets). Highly ordered cylindrical microdomains oriented normal to the films surface were obtained where the Fourier transform of the AFM images are shown in the insets to characterize the lateral order. The areal densities obtained are also specified in the insets. Reproduced with permission from ref 136. Copyright 2009 AAAS. (B) Spin-coated film of PS-b-PEO asymmetric diblock copolymer having cylindrical microdomains on a faceted sapphire substrate that was thermally annealed. The direction of the ridges on the sapphire substrate is indicated along with a Fourier transform of the AFM image. Reproduced from ref 137. Copyright 2012 PNAS.

interaction conditions as the block copolymer is converted. Consequently, even though the block copolymer and segmental interactions have significantly changed, to the point where the block copolymer microphase separates into ultrasmall domains, the interfacial interactions have also changed to maintain balanced interfacial interactions with the copolymer. By varying the chemical composition of the surface laterally, while maintaining a planar geometry, the preferential interfacial interactions can be patterned, directing the spatial location and orientation of the block copolymer microdomains. Graphoepitaxy, for example, has been used to propagate ordering of spherical (Figure 5), cylindrical, and lamellar microdomains in the plane of the film.126−129 Shear has also been shown as an effective and promising route to achieve long-range lateral order.130 Significant advances have been made in the refinement of the surface patterning,131,132 in minimizing the extent of surface patterning through pattern amplification,133,134 and in achieving long-range lateral ordering of the microdomains. This directed self-assembly of block copolymers may, in fact, be the only route by which perfection in the lateral ordering of the copolymer microdomains can be achieved, where strong interfacial interactions can be used to suppress defects routinely encountered in the self-assembly of the

microdomains. Even if one could nucleate the growth of a single grain of microdomains at one point or line on the surface, or use gradient fields to propagate order across an entire surface, surface patterning would still be needed to suppress defects routinely encountered in the self-assembly process. The forces that drive self-assembly are relatively weak, leading to fluctuations in the structures and inherent defects. With blocks, where segmental interactions are strongly nonfavorable, as with hydrophobic/hydrophilic interactions, or with blocks that are rigid, the costs of defects increase but so does the difficulty in producing uniform films from a common solvent or mixtures of solvents. Producing perfection in lateral order still remains a challenge, though the combination of topdown lithographic process with bottom-up self-assembly remains as the most viable route to produce addressable media to take storage to the next generation and interfacial interactions remain central. Varying the surface topography opens numerous routes to manipulate the interfacial behavior of polymers and block copolymers. Simply changing the surface from being perfectly smooth where the strength of interfacial interaction is given by γ, where γ can preferentially interact with one block forcing an orientation of the copolymer microdomains parallel to the G

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Macromolecules surface, to a surface with a roughness less than a period of the microdomain morphology but with sufficient amplitude comparable to the period, the preferential interactions will be mediated by the copolymer filling the pores such that in the plane defining the interface γ is effectively reduced to the point where the block copolymer microdomains orient normal to the surface.135 This is akin to Cassie’s law, where surface roughness modifies the contact angle of a water droplet on a surface due to the incomplete penetration of the water into the pores and the interfacial energy depends on the fraction of the surface in contact with the droplet. If the amplitude of the roughness is large and the lateral size of the roughness is larger than the copolymer period, as in lithographically generated trench patterns, the entropic constraints placed on the block copolymer chains coupled with preferential interfacial interactions can lead to an orientation of the microdomains parallel to or orthogonal to the trenches.136 If the surface roughness is small in amplitude but with a well-defined period, the packing of the microdomains within the surface features, as in the reconstructed surface of sapphire or silicon with a periodic sawtooth pattern (Figure 6), will guide the lateral orientation of the microdomains. For cylinders oriented normal to the surface, as can easily be produced with solvent annealing where interfacial and surface energies are mediated, large area arrays with superb in-plane orientational order can be achieved, though the translational order of the microdomains is only slightly better than a hexatic. For cylindrical microdomains oriented parallel to the surface, taking advantage of preferential interactions of the blocks with the substrate and the entropic constraints in packing the microdomains at the interface, perfect line patterns can be produced over the entire surface of the film, however large it may be.137 Surface topography can also be lithographically patterned into well-defined shapes, or one could use pores in alumina membranes to confine block copolymers where curvature can be introduced, forcing the block copolymer to assumed morphologies that are far removed from their bulk morphologies. For example, confining cylinder-forming copolymers within the cylindrical pores of an alumina membrane (Figure 7), depending on the diameter of the pores relative to the period of the block copolymer, cylinders oriented in the direction of the pores can be found or the copolymer can be forced into forming single or multiple helices that wrap around the central axis of the pores.138,139 Gyroid-forming block copolymers are less studied but can also form exotic structures when cylindrically confined140 or in thin film.141 Lamellae can be forced into concentric cylinders within the pores, where the period varies radially.138 Spherical confinement has also been achieved where the competition between interfacial interactions with the medium in which the spherical particles are dispersed and the blocks of the copolymer and the deformation imposed on the copolymer chains leads to yet more novel morphologies.142−144 While similar types of morphologies can be found for circular depressions etched into a substrate, lithography offers the opportunity to place well-defined features within the circular depression where, by placing a small step feature in the otherwise circular pattern, the interactions of the copolymer with the walls of the pattern force the cylinderforming copolymer into a spiral pattern (Figure 7). This opens the very exciting possibility of combining cylindrical confinement with a guiding notch at one end such that helical morphologies are formed that are chiral, i.e., all right-handed or all left-handed, where the notch is used to guide the

Figure 7. (A) PS-b-PBD diblock copolymer having cylindrical microdomains confined within the pores of an aluminum oxide membrane where the ratio of the pore diameter to the period of the copolymer is 1.27. Both a TEM image and a TEM tomographic image are shown of the copolymer released from the membrane. Reproduced with permission from ref 138. (B) SEM images of PS-b-PDMS diblock copolymer having cylindrical microdomains confined to a circular trench (top) and to trenches where divots have been lithographically placed in the trenches to guide the direction of the copolymer assembly (left-handed spiral (middle) and right-handed spiral (bottom)). Adapted with permission from ref 161.

handedness and, also, act as a nucleation sites to initiate the formation of the chiral, helical microdomains. Monte Carlo simulations of block copolymers have indicated that defect-free arrays of microdomains, a holy grail in the selfassembly of block copolymers, can be achieved on chemically patterned substrates;145 significant strides have been made in chemically patterning substrates to guide the assembly of block copolymers over macroscopic areas.146,147 Determining whether the arrays of the microdomains are perfect requires methods to characterize thin films on the nanoscopic size scale over macroscopic distances. While grazing incidence scattering methods do the job, the results represent a statistical averaging of the surface and cannot isolate or locate individual defects. Scanning probe methods can provide this information, but only over a limited area. Sequential scanning of multiple areas with proscribed area of overlap (as shown in Figure 5) can be done, but are time-consuming and laborious. In addition, not only is the structure at the free surface of interest but also the variation in the morphology as a function of depth, arising from slight mismatches in the chemical patterning and the natural morphology of the block copolymer. This can give rise to distortions or tilting of the morphology. Ar ion milling coupled with scanning electron microscopy can yield an unambiguous description of the morphology over a very limited area.148 Similarly, transmission electron microscopy tomography can be used to elucidate the three-dimensional morphology of the block copolymer over a limited volume and cannot be used practically over macroscopic areas. Theoretical and computational studies have described the morphology arising from the directed self-assembly of block copolymers under different confinements.149−153 Consequently, theory can predict, but H

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level resolution where, for crystalline polymers, individual atoms comprising the polymer can be resolved. Extending this to disordered materials, while a challenge, is possible, opening the opportunity to obtain atomic resolution on disordered glassy materials, which would be a stunning advance, revolutionizing our understanding of disordered materials. Impediments in probing interfaces from the relatively low interfacial volume, in obtaining information relevant to the interface independent of the surrounding material, and the dynamics occurring on the second to nanosecond time scale have limited existing probes in revealing either local or statistical information on the behavior of materials at interfaces. With many of the instrumentation advances the development of accessible and informed software packages for data processing and analysis is essential. As it stands currently, the amount of data that is produced in experiments at advanced sources is enormous, and particularly with time-resolved or dynamic experiments, handling the data in an intelligent and efficient manner cannot be overlooked. With that aid, the advances in instrumentation that are on the horizon will provide pathways to overcome these challenges, opening broad areas of research on polymer thin films and polymers at interfaces.

experimentally realizing and demonstrating perfection in lateral order still remains a challenge.

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PROSPECTS Some of the advances that have been made in our understanding of polymer thin films and at surfaces and interfaces have been realized by advances in simulations and by extraordinary advances in experimental methods to probe polymeric materials. Scanning probe microscopies, discovered only a few decades ago, with height and lateral spatial resolution of several segment lengths, are now routine tools available in most laboratories. The spatial resolution of electron microscopy and the imaging software to perform tomography have opened three-dimensional imaging of polymer structures and morphologies. Initial hard and soft synchrotron X-rays sources were only first used on polymers some 30 years ago, and now grazing incidence X-ray scattering to elucidate the surface and thin film structure and morphology of polymers is routine,154 while soft X-rays, with chemical bond sensitivity provided by variation of the incident X-ray energy, allow the absolute definition of a structure without the phase problem inherent to single wavelength X-ray measurement.155 Utilization of the polarized nature of the synchrotron source provides direct access to interfaces where, by default, there is an anisotropic spatial arrangement of polymer chains or segments and, therefore, the chemical bonding of the constituents.156,157 Neutron reflectivity has provided information on interfaces, both buried and at surfaces, with a spatial resolution of a nanometer or less.158,159 There has been a host of different dynamic fluorescence optical microscopy methods that have emerged providing unprecedented lateral spatial resolution. All of these developments have been used to gain insight into polymers at surfaces and interfaces. Yet, there are still many questions and opportunities in the study of polymer in thin films and at interfaces, particularly in the dynamics. Chemical reactions, charge transport, and much device function occur at interfaces of solids, liquids, and gases. Both the chemistry and morphology of the interfaces and how they change over the second to microsecond time scale are critical to our understanding and ability to predict their function of this important class of materials. The complex structure and dynamics of the solid−electrolyte interphases, electrical double layers, photosynthesis, charge transport in organic electronic materials, glassy dynamics over all spatial and time scales (filling in the missing gaps), dielectric breakdown and other rare events, and the dynamics of polymers immediately at interfaces are all topics that are open. To gain insight into these problems, innovative experimental techniques will be needed. Synchrotron X-ray sources are now moving to fully coherent sources where X-ray photon correlation spectroscopy is enabled, where X-ray ptychography will enable real space imaging of materials with near atomic resolution, and standing wave methods (coupled with the coherence and energy resolution) will enable the interrogation of interface dynamics without the dynamics of the material surrounding the interface provided unwanted background. Neutron sources, particularly spallation neutron sources, are being made more brilliant with the promise of a nearly 30-fold increase in the brilliance being developed. This would enable the investigations of single chain conformations in single thin films, the kinetic processes at interfaces (like interdiffusion and reaction polymerization) in real time, and a range of different kinetic processes. Transmission electron microscopy now has atomic-

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected]. ORCID

Thomas P. Russell: 0000-0001-6384-5826 Notes

The authors declare no competing financial interest. Biographies

Thomas P. Russell, the Silvio O. Conte Distinguished Professor of Polymer Science and Engineering at the University of Massachusetts in Amherst, received his PhD in 1979 from the same institution. He was a Research Associate at the University of Mainz (1979−1981), a Research Staff Member at the IBM Almaden Research Center in San Jose, CA (1981−1996), and a Professor of Polymer Science and Engineering at the University of Massachusetts Amherst (1997). He is also a Visiting Faculty at the Materials Sciences Division in the Lawrence Berkeley National Laboratory, an Adjunct Professor at the Beijing University of Chemical Technology, and a lead PI at the Advanced Institute of Materials Research at Tohoku University. He was the Director of the Materials Research Science and Engineering Center from (1996−2009) and the Director of the Energy Frontier Research Center on Polymer-Based Materials for Harvesting Solar Energy (2009−2014) and is a lead PI in the WPI-Advanced Institute I

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(10) Jones, R.; Kumar, S.; Ho, D.; Briber, R.; Russell, T. Chain Conformation in Ultrathin Polymer Films. Nature 1999, 400, 146− 149. (11) Jones, R. L.; Kumar, S. K.; Ho, D. L.; Briber, R. M.; Russell, T. P. Chain Conformation in Ultrathin Polymer Films Using Small-Angle Neutron Scattering. Macromolecules 2001, 34, 559−567. (12) Brûlet, A.; Boué, F.; Menelle, A.; Cotton, J. P. Conformation of Polystyrene Chain in Ultrathin Films Obtained by Spin Coating. Macromolecules 2000, 33, 997−1001. (13) Ediger, M. D.; Angell, C. A.; Nagel, S. R. Supercooled Liquids and Glasses. J. Phys. Chem. 1996, 100, 13200−13212. (14) Debenedetti, P. G.; Stillinger, F. H. Supercooled Liquids and the Glass Transition. Nature 2001, 410, 259−267. (15) Gebremichael, Y.; Schrøder, T. B.; Starr, F. W.; Glotzer, S. C. Spatially Correlated Dynamics in a Simulated Glass-Forming Polymer Melt: Analysis of Clustering Phenomena. Phys. Rev. E: Stat. Phys., Plasmas, Fluids, Relat. Interdiscip. Top. 2001, 64, 51503. (16) Pan, A. C.; Garrahan, J. P.; Chandler, D. Heterogeneity and Growing Length Scales in the Dynamics of Kinetically Constrained Lattice Gases in Two Dimensions. Phys. Rev. E 2005, 72, 41106. (17) Chandler, D. Liquids: Condensed, Disordered, and Sometimes Complex. Proc. Natl. Acad. Sci. U. S. A. 2009, 106, 15111−15112. (18) Chandler, D.; Garrahan, J. P. Dynamics on the Way to Forming Glass: Bubbles in Space-Time. Annu. Rev. Phys. Chem. 2010, 61, 191− 217. (19) Garrahan, J. P. Dynamic Heterogeneity Comes to Life. Proc. Natl. Acad. Sci. U. S. A. 2011, 108, 4701−4702. (20) Liu, A. J.; Nagel, S. R. The Jamming Transition and the Marginally Jammed Solid. Annu. Rev. Condens. Matter Phys. 2010, 1, 347−369. (21) Jaeger, H. M.; Nagel, S. R.; Behringer, R. P. Granular Solids, Liquids, and Gases. Rev. Mod. Phys. 1996, 68, 1259−1273. (22) Berthier, L.; Biroli, G.; Bouchaud, J.-P.; Cipelletti, L.; van Saarloos, W. Dynamical Heterogeneities in Glasses, Colloids, and Granular Media; OUP: Oxford, 2011; Vol. 150. (23) Kovacs, A. J. Glass Transition in Amorphous Polymers: A Phenomenological Study. Adv. Polym. Sci. 1963, 3, 394−508. (24) Ferry, J. D. Viscoelastic Properties of Polymers; John Wiley & Sons: 1980. (25) McCrum, N. G.; Read, B. E.; Williams, G. Anelastic and Dielectric Effects in Polymeric Solids; Wiley: 1967. (26) Keddie, J. L.; Jones, R.; Cory, R. Size-Dependent Depression of the Glass Transition Temperature in Polymer Films. Europhys. Lett. 1994, 27, 59−64. (27) Ediger, M. D.; Forrest, J. A. Dynamics near Free Surfaces and the Glass Transition in Thin Polymer Films: A View to the Future. Macromolecules 2014, 47, 471−478. (28) Jackson, C. L.; McKenna, G. B. The Glass Transition of Organic Liquids Confined to Small Pores. J. Non-Cryst. Solids 1991, 131−133, 221−224. (29) Swallen, S. F.; Kearns, K. L.; Mapes, M. K.; Kim, Y. S.; McMahon, R. J.; Ediger, M. D.; Wu, T.; Yu, L.; Satija, S. Organic Glasses with Exceptional Thermodynamic and Kinetic Stability. Science 2007, 315, 353−356. (30) Guo, Y.; Morozov, A.; Schneider, D.; Chung, J. W.; Zhang, C.; Waldmann, M.; Yao, N.; Fytas, G.; Arnold, C. B.; Priestley, R. D. Ultrastable Nanostructured Polymer Glasses. Nat. Mater. 2012, 11, 337−343. (31) Forrest, J. A.; Dalnoki-Veress, K.; Stevens, J. R.; Dutcher, J. R. Effect of Free Surfaces on the Glass Transition Temperature of Thin Polymer Films. Phys. Rev. Lett. 1996, 77, 2002−2005. (32) Sharp, J. S.; Forrest, J. A. Free Surfaces Cause Reductions in the Glass Transition Temperature of Thin Polystyrene Films. Phys. Rev. Lett. 2003, 91, 235701. (33) Bäumchen, O.; McGraw, J. D.; Forrest, J. A.; Dalnoki-Veress, K. Reduced Glass Transition Temperatures in Thin Polymer Films: Surface Effect or Artifact? Phys. Rev. Lett. 2012, 109, 55701.

of Materials Research at Tohoku University (2006−present), the Global Research Laboratory at Seoul National University (2005− 2015), and the Beijing Advanced Innovation Center on Soft Matter (2016−present). He is a Fellow of the American Physical Society, Materials Research Society, Neutron Scattering Society of America, American Association for the Advancement of Science, and the American Chemical Society, Polymer Materials Science and Engineering Division. He has received the Polymer Physic Prize of the APS, the Cooperative Research Award of the ACS, the Dutch Polymer Award, the ACS Award in Applied Polymer Science, and Society of Polymer Science Japan International Award and is an elected member of the National Academy of Engineering. He is also an Associate Editor of Macromolecules.

Yu Chai, a postdoc at UC Berkeley and Lawrence Berkeley National Laboratory with Professor Thomas P. Russell, received both his BSc (2011) and PhD (2016) in physics from the University of Waterloo under the supervision of Professor James A. Forrest. His main research interests are fundamental physics of polymers, colloids, and other soft matters, particularly the dynamics of polymers and colloids at surfaces and interfaces.

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REFERENCES

(1) Flory, P. J. Principles of Polymer Chemistry; Cornell University Press: 1953. (2) Bitsanis, I. A.; ten Brinke, G. A Lattice Monte Carlo Study of Long Chain Conformations at Solid−polymer Melt Interfaces. J. Chem. Phys. 1993, 99, 3100−3111. (3) Kumar, S. K.; Vacatello, M.; Yoon, D. Y. Off-lattice Monte Carlo Simulations of Polymer Melts Confined between Two Plates. J. Chem. Phys. 1988, 89, 5206−5215. (4) Yethiraj, A.; Hall, C. K. Square-well Chains: Bulk Equation of State Using Perturbation Theory and Monte Carlo Simulations of the Bulk Pressure and of the Density Profiles near Walls. J. Chem. Phys. 1991, 95, 1999−2005. (5) Theodorou, D. N. Structure and Thermodynamics of Bulk Homopolymer Solid Interfaces: A Site Lattice Model Approach. Macromolecules 1988, 21, 1400−1410. (6) Pakula, T. Computer Simulation of Polymers in Thin Layers. I. Polymer Melt between Neutral Walls − Static Properties. J. Chem. Phys. 1991, 95, 4685−4690. (7) Zheng, X.; Rafailovich, M. H.; Sokolov, J.; Strzhemechny, Y.; Schwarz, S. A.; Sauer, B. B.; Rubinstein, M. Long-Range Effects on Polymer Diffusion Induced by a Bounding Interface. Phys. Rev. Lett. 1997, 79, 241−244. (8) Calvert, P. Rough Guide to the Nanoworld. Nature 1996, 383, 300−301. (9) Frank, C. W.; Rao, V.; Despotopoulou, M. M.; Pease, R. F. W.; Hinsberg, W. D.; Miller, R. D.; Rabolt, J. F. Structure in Thin and Ultrathin Spin-Cast Polymer Films. Science 1996, 273, 912−915. J

DOI: 10.1021/acs.macromol.7b00418 Macromolecules XXXX, XXX, XXX−XXX

Perspective

Macromolecules (34) Yang, Z.; Fujii, Y.; Lee, F. K.; Lam, C.-H.; Tsui, O. K. C. Glass Transition Dynamics and Surface Layer Mobility in Unentangled Polystyrene Films. Science 2010, 328, 1676−1679. (35) Fakhraai, Z.; Forrest, J. A. Measuring the Surface Dynamics of Glassy Polymers. Science 2008, 319, 600−604. (36) Chai, Y.; Salez, T.; McGraw, J. D.; Benzaquen, M.; DalnokiVeress, K.; Raphael, E.; Forrest, J. A. A Direct Quantitative Measure of Surface Mobility in a Glassy Polymer. Science 2014, 343, 994−999. (37) Kim, H.; Rühm, A.; Lurio, L. B.; Basu, J. K.; Lal, J.; Lumma, D.; Mochrie, S. G. J.; Sinha, S. K. Surface Dynamics of Polymer Films. Phys. Rev. Lett. 2003, 90, 68302. (38) Sinha, S. K.; Jiang, Z.; Lurio, L. B. X-Ray Photon Correlation Spectroscopy Studies of Surfaces and Thin Films. Adv. Mater. 2014, 26, 7764−7785. (39) Zhao, W.; Zhao, X.; Rafailovich, M. H.; Sokolov, J.; Composto, R. J.; Smith, S. D.; Russell, T. P.; Dozier, W. D.; Mansfield, T.; Satkowski, M. Segregation of Chain Ends to Polymer Melt Surfaces and Interfaces. Macromolecules 1993, 26, 561−562. (40) Jalbert, C.; Koberstein, J. T.; Yilgor, I.; Gallagher, P.; Krukonis, V. Molecular Weight Dependence and End-Group Effects on the Surface Tension of Poly(dimethylsiloxane). Macromolecules 1993, 26, 3069−3074. (41) Wu, D. T.; Fredrickson, G. H.; Carton, J.-P.; Ajdari, A.; Leibler, L. Distribution of Chain Ends at the Surface of a Polymer Melt: Compensation Effects and Surface Tension. J. Polym. Sci., Part B: Polym. Phys. 1995, 33, 2373−2389. (42) Matsen, M. W.; Mahmoudi, P. Segregation of Chain Ends to the Surface of a Polymer Melt. Eur. Phys. J. E: Soft Matter Biol. Phys. 2014, 37, 78. (43) Serghei, A.; Tress, M.; Kremer, F. The Glass Transition of Thin Polymer Films in Relation to the Interfacial Dynamics. J. Chem. Phys. 2009, 131, 154904. (44) Serghei, a.; Huth, H.; Schick, C.; Kremer, F. Glassy Dynamics in Thin Polymer Layers Having a Free Upper Interface. Macromolecules 2008, 41, 3636−3639. (45) Tress, M.; Mapesa, E. U.; Kossack, W.; Kipnusu, W. K.; Reiche, M.; Kremer, F. Glassy Dynamics in Condensed Isolated Polymer Chains. Science 2013, 341, 1371−1374. (46) Sasaki, T.; Shimizu, A.; Mourey, T. H.; Thurau, C. T.; Ediger, M. D. Glass Transition of Small Polystyrene Spheres in Aqueous Suspensions. J. Chem. Phys. 2003, 119, 8730−8735. (47) Samant, M. G.; Stöhr, J.; Brown, H. R.; Russell, T. P.; Sands, J. M.; Kumar, S. K. NEXAFS Studies on the Surface Orientation of Buffed Polyimides. Macromolecules 1996, 29, 8334−8342. (48) Liu, Y.; Russell, T. P.; Samant, M. G.; Stöhr, J.; Brown, H. R.; Cossy-Favre, A.; Diaz, J. Surface Relaxations in Polymers. Macromolecules 1997, 30, 7768−7771. (49) Kerle, T.; Lin, Z.; Kim, H.-C.; Russell, T. P. Mobility of Polymers at the Air/Polymer Interface. Macromolecules 2001, 34, 3484−3492. (50) Riggleman, R. A.; Yoshimoto, K.; Douglas, J. F.; de Pablo, J. J. Influence of Confinement on the Fragility of Antiplasticized and Pure Polymer Films. Phys. Rev. Lett. 2006, 97, 45502. (51) Ellison, C. J.; Ruszkowski, R. L.; Fredin, N. J.; Torkelson, J. M. Dramatic Reduction of the Effect of Nanoconfinement on the Glass Transition of Polymer Films via Addition of Small-Molecule Diluent. Phys. Rev. Lett. 2004, 92, 095702. (52) Fakhraai, Z.; Forrest, J. A. Probing Slow Dynamics in Supported Thin Polymer Films. Phys. Rev. Lett. 2005, 95, 25701. (53) Hawker, C. J.; Farrington, P. J.; Mackay, M. E.; Wooley, K. L.; Fréchet, J. M. J. Molecular Ball Bearings: The Unusual Melt Viscosity Behavior of Dendritic Macromolecules. J. Am. Chem. Soc. 1995, 117, 4409−4410. (54) Fryer, D. S.; Peters, R. D.; Kim, E. J.; Tomaszewski, J. E.; de Pablo, J. J.; Nealey, P. F.; White, C. C.; Wu, W. Dependence of the Glass Transition Temperature of Polymer Films on Interfacial Energy and Thickness. Macromolecules 2001, 34, 5627−5634.

(55) Christie, D.; Zhang, C.; Fu, J.; Koel, B.; Priestley, R. D. Glass Transition Temperature of Colloidal Polystyrene Dispersed in Various Liquids. J. Polym. Sci., Part B: Polym. Phys. 2016, 54, 1776−1783. (56) Qi, D.; Fakhraai, Z.; Forrest, J. A. Substrate and Chain Size Dependence of Near Surface Dynamics of Glassy Polymers. Phys. Rev. Lett. 2008, 101, 96101. (57) Ellison, C. J.; Kim, S. D.; Hall, D. B.; Torkelson, J. M. Confinement and Processing Effects on Glass Transition Temperature and Physical Aging in Ultrathin Polymer Films: Novel Fluorescence Measurements. Eur. Phys. J. E: Soft Matter Biol. Phys. 2002, 8, 155− 166. (58) Ellison, C. J.; Torkelson, J. M. Sensing the Glass Transition in Thin and Ultrathin Polymer Films via Fluorescence Probes and Labels. J. Polym. Sci., Part B: Polym. Phys. 2002, 40, 2745−2758. (59) Ellison, C. J.; Torkelson, J. M. The Distribution of GlassTransition Temperatures in Nanoscopically Confined Glass Formers. Nat. Mater. 2003, 2, 695−700. (60) Zheng, X.; Sauer, B. B.; Van Alsten, J. G.; Schwarz, S. A.; Rafailovich, M. H.; Sokolov, J.; Rubinstein, M. Reptation Dynamics of a Polymer Melt near an Attractive Solid Interface. Phys. Rev. Lett. 1995, 74, 407−410. (61) Shin, K.; Obukhov, S.; Chen, J.-T.; Huh, J.; Hwang, Y.; Mok, S.; Dobriyal, P.; Thiyagarajan, P.; Russell, T. P. Enhanced Mobility of Confined Polymers. Nat. Mater. 2007, 6, 961−965. (62) Priestley, R. D.; Ellison, C. J.; Broadbelt, L. J.; Torkelson, J. M.; M, T. J. Structural Relaxation of Polymer Glasses at Surfaces, Interfaces, and In Between. Science 2005, 309, 456−459. (63) Zhang, C.; Guo, Y.; Priestley, R. D. Glass Transition Temperature of Polymer Nanoparticles under Soft and Hard Confinement. Macromolecules 2011, 44, 4001−4006. (64) Guo, Y.; Zhang, C.; Lai, C.; Priestley, R. D.; D’Acunzi, M.; Fytas, G. Structural Relaxation of Polymer Nanospheres under Soft and Hard Confinement: Isobaric versus Isochoric Conditions. ACS Nano 2011, 5, 5365−5373. (65) Feynman, R. P. There’s Plenty of Room at the Bottom. Eng. Sci. 1960, 23, 22−36. (66) Hervet, H.; Degennes, P. G. The Dynamics of WettingPrecursor Films in the Wetting of Dry Solids. C. R. Acad. Des Sci., Ser. II 1984, 299, 499−503. (67) de Gennes, P. G. Wetting: Statics and Dynamics. Rev. Mod. Phys. 1985, 57, 827−863. (68) Reiter, G. Dewetting of Thin Polymer Films. Phys. Rev. Lett. 1992, 68, 75−78. (69) Reiter, G. Unstable Thin Polymer Films: Rupture and Dewetting Processes. Langmuir 1993, 9, 1344−1351. (70) Kerle, T.; Yerushalmi-Rozen, R.; Klein, J.; Fetters, L. J. Van Der Waals Stable Thin Liquid Films: Correlated Undulations and Ultimate Dewetting. Eur. Lett. 1998, 44, 484−490. (71) Geoghegan, M.; Krausch, G. Wetting at Polymer Surfaces and Interfaces. Prog. Polym. Sci. 2003, 28, 261−302. (72) Harkins, W. D.; Feldman, A. Films. the Spreading of Liquids and the Spreading Coefficient. J. Am. Chem. Soc. 1922, 44, 2665−2685. (73) Redon, C.; Brochard-Wyart, F.; Rondelez, F. Dynamics of Dewetting. Phys. Rev. Lett. 1991, 66, 715−718. (74) Brochard-Wyart, F.; de Gennes, P. G. Dynamics of Partial Wetting. Adv. Colloid Interface Sci. 1992, 39, 1−11. (75) Seemann, R.; Herminghaus, S.; Jacobs, K. Dewetting Patterns and Molecular Forces: A Reconciliation. Phys. Rev. Lett. 2001, 86, 5534−5537. (76) Seemann, R.; Herminghaus, S.; Jacobs, K. Shape of a Liquid Front upon Dewetting. Phys. Rev. Lett. 2001, 87, 196101. (77) Herminghaus, S.; Seemann, R.; Jacobs, K. Generic Morphologies of Viscoelastic Dewetting Fronts. Phys. Rev. Lett. 2002, 89, 56101. (78) Reiter, G. Dewetting of Highly Elastic Thin Polymer Films. Phys. Rev. Lett. 2001, 87, 186101. (79) Saulnier, F.; Raphael, E.; De Gennes, P. G. Dewetting of Thin Polymer Films near the Glass Transition. Phys. Rev. Lett. 2002, 88, 4. (80) Gent, A. N.; Hamed, G. R. Fundamentals of Adhesion; Springer: Boston, MA, 1990. K

DOI: 10.1021/acs.macromol.7b00418 Macromolecules XXXX, XXX, XXX−XXX

Perspective

Macromolecules (81) Brown, H. R. The Adhesion Between Polymers. Annu. Rev. Mater. Sci. 1991, 21, 463−489. (82) Brown, H. R. Effect of a Diblock Copolymer on the Adhesion between Incompatible Polymers. Macromolecules 1989, 22, 2859− 2860. (83) Shull, K.; Kramer, E. J. Mean-Field Theory of Polymer Interfaces in the Presence of Block Copolymers. Macromolecules 1990, 23, 4769−4779. (84) Brown, H. R.; Char, K.; Deline, V. R.; Green, P. F. Effects of a Diblock Copolymer on Adhesion between Immiscible Polymers. 1. Polystyrene (PS)-PMMA Copolymer between PS and PMMA. Macromolecules 1993, 26, 4155−4163. (85) Dai, K. H.; Norton, L. J.; Kramer, E. J. Equilibrium Segment Density Distribution of a Diblock Copolymer Segregated to a Polymer/polymer Interface. Macromolecules 1994, 27, 1949−1956. (86) Kulasekere, R.; Kaiser, H.; Ankner, J. F.; et al. Homopolymer Interfaces Reinforced with Random Copolymers. Macromolecules 1996, 29, 5493−5496. (87) Milner, S. T.; Fredrickson, G. H. Reconsidering Random Copolymers at Interfaces. Macromolecules 1995, 28, 7953−7956. (88) Kim, H.; Mohapatra, H.; Phillips, S. T. Rapid, On-Command Debonding of Stimuli-Responsive Cross-Linked Adhesives by Continuous, Sequential Quinone Methide Elimination Reactions. Angew. Chem., Int. Ed. 2015, 54, 13063−13067. (89) Huang, Y.; Kramer, E. J.; Heeger, A. J.; Bazan, G. C. Bulk Heterojunction Solar Cells: Morphology and Performance Relationships. Chem. Rev. 2014, 114, 7006−7043. (90) Lu, L.; Zheng, T.; Wu, Q.; Schneider, A. M.; Zhao, D.; Yu, L. Recent Advances in Bulk Heterojunction Polymer Solar Cells. Chem. Rev. 2015, 115, 12666−12731. (91) Dou, L.; Liu, Y.; Hong, Z.; Li, G.; Yang, Y. Low-Bandgap NearIR Conjugated Polymers/Molecules for Organic Electronics. Chem. Rev. 2015, 115, 12633−12665. (92) Wang, C.; Dong, H.; Hu, W.; Liu, Y.; Zhu, D. Semiconducting π-Conjugated Systems in Field-Effect Transistors: A Material Odyssey of Organic Electronics. Chem. Rev. 2012, 112, 2208−2267. (93) Zhao, X.; Zhan, X. Electron Transporting Semiconducting Polymers in Organic Electronics. Chem. Soc. Rev. 2011, 40, 3728− 3743. (94) Ma, H.; Yip, H.-L.; Huang, F.; Jen, A. K. Y. Interface Engineering for Organic Electronics. Adv. Funct. Mater. 2010, 20, 1371−1388. (95) Hoven, C.; Yang, R.; Garcia, A.; Heeger, A. J.; Nguyen, T. Q.; Bazan, G. C. Ion Motion in Conjugated Polyelectrolyte Electron Transporting Layers. J. Am. Chem. Soc. 2007, 129, 10976−10977. (96) Di, C.; Liu, Y.; Yu, G.; Zhu, D.; C. Di, Y.; Liu, G.; Yu, D. Z.; Di, C.; Liu, Y.; Yu, G.; Zhu, D. Interface Engineering: An Effective Approach toward High-Performance Organic Field-Effect Transistors. Acc. Chem. Res. 2009, 42, 1573−1583. (97) He, Z.; Wu, H.; Cao, Y. Recent Advances in Polymer Solar Cells: Realization of High Device Performance by Incorporating Water/alcohol-Soluble Conjugated Polymers as Electrode Buffer Layer. Adv. Mater. 2014, 26, 1006−1024. (98) Dou, L.; Yang, Y. M.; You, J.; Hong, Z.; Chang, W.-H.; Li, G.; Yang, Y. Solution-Processed Hybrid Perovskite Photodetectors with High Detectivity. Nat. Commun. 2014, 5, 5404. (99) Zhou, Y.; Fuentes-Hernandez, C.; Shim, J.; Meyer, J.; Giordano, A. J.; Li, H.; Winget, P.; Papadopoulos, T.; Cheun, H.; Kim, J.; Fenoll, M.; Dindar, A.; Haske, W.; Najafabadi, E.; Khan, T. M.; Sojoudi, H.; Barlow, S.; Graham, S.; Brédas, J.; Marder, S. R.; Kahn, A.; Kippelen, B.; Bredas, J.-L.; Marder, S. R.; Kahn, A.; Kippelen, B.; Brédas, J.; Marder, S. R.; Kahn, A.; Kippelen, B. A Universal Method to Produce Low-Work Function Electrodes for Organic Electronics. Science 2012, 336, 327−332. (100) Huang, F.; Wu, H.; Cao, Y. Water/alcohol Soluble Conjugated Polymers as Highly Efficient Electron Transporting/injection Layer in Optoelectronic Devices. Chem. Soc. Rev. 2010, 39, 2500−2521.

(101) Yip, H.-L.; Jen, A. K.-Y. Recent Advances in SolutionProcessed Interfacial Materials for Efficient and Stable Polymer Solar Cells. Energy Environ. Sci. 2012, 5, 5994. (102) Chueh, C.; Li, C.; Jen, A. K. Environmental Science Recent Progress and Perspective in Solution- Processed Interfacial Materials for Efficient and Stable Polymer and Organometal Perovskite Solar Cells. Energy Environ. Sci. 2015, 8, 1160−1189. (103) He, Z.; Xiao, B.; Liu, F.; Wu, H.; Yang, Y.; Xiao, S.; Wang, C.; Russell, T. P.; Cao, Y. Single-Junction Polymer Solar Cells with High Efficiency and Photovoltage. Nat. Photonics 2015, 9, 174−179. (104) Jiang, H.; Taranekar, P.; Reynolds, J. R.; Schanze, K. S. Conjugated Polyelectrolytes: Synthesis, Photophysics, and Applications. Angew. Chem., Int. Ed. 2009, 48, 4300−4316. (105) Duarte, A.; Pu, K. Y.; Liu, B.; Bazan, G. C. Recent Advances in Conjugated Polyelectrolytes for Emerging Optoelectronic Applications. Chem. Mater. 2011, 23, 501−515. (106) Hu, Z. C.; Zhang, K.; Huang, F.; Cao, Y. Water/alcohol Soluble Conjugated Polymers for the Interface Engineering of Highly Efficient Polymer Light-Emitting Diodes and Polymer Solar Cells. Chem. Commun. 2015, 51, 5572−5585. (107) Tordera, D.; Kuik, M.; Rengert, Z. D.; Bandiello, E.; Bolink, H. J.; Bazan, G. C.; Nguyen, T. Q. Operational Mechanism of Conjugated Polyelectrolytes. J. Am. Chem. Soc. 2014, 136, 8500−8503. (108) Fang, J.; Wallikewitz, B. H.; Gao, F.; Tu, G.; Müller, C.; Pace, G.; Friend, R. H.; Huck, W. T. S. Conjugated Zwitterionic Polyelectrolyte as the Charge Injection Layer for High-Performance Polymer Light-Emitting Diodes. J. Am. Chem. Soc. 2011, 133, 683− 685. (109) Coulon, G.; Russell, T. P.; Deline, V. R.; Green, P. F. SurfaceInduced Orientation of Symmetric, Diblock Copolymers: A Secondary Ion Mass-Spectrometry Study. Macromolecules 1989, 22, 2581−2589. (110) Russell, T. P.; Coulon, G.; Deline, V. R.; Miller, D. C. Characteristics of the Surface-Induced Orientation for Symmetric Diblock PS/PMMA Copolymers. Macromolecules 1989, 22, 4600− 4606. (111) Fasolka, M. J.; Mayes, A. M. Block Copolymer Thin Films: Physics and Applications. Annu. Rev. Mater. Res. 2001, 31, 323−355. (112) Coulon, G.; Daillant, J.; Collin, B.; Benattar, J. J.; Gallot, Y. Time Evolution of the Free Surface of Ultrathin Copolymer Films. Macromolecules 1993, 26, 1582−1589. (113) Anastasiadis, S. H.; Russell, T. P.; Satija, S. K.; Majkrzak, C. F. Neutron Reflectivity Studies of the Surface-Induced Ordering of Diblock Copolymer Films. Phys. Rev. Lett. 1989, 62, 1852−1855. (114) Stein, G. E.; Kramer, E. J.; Li, X.; Wang, J. Layering Transitions in Thin Films of Spherical-Domain Block Copolymers. Macromolecules 2007, 40, 2453−2460. (115) Collin, B.; Chatenay, D.; Coulon, G.; Ausserre, D.; Gallot, Y. Ordering of Copolymer Thin Films as Revealed by Atomic Force Microscopy. Macromolecules 1992, 25, 1621−1622. (116) Menelle, A.; Russell, T. P.; Anastasiadis, S. H.; Satija, S. K.; Majkrzak, C. F. Ordering of Thin Diblock Copolymer Films. Phys. Rev. Lett. 1992, 68, 67−70. (117) Mansky, P.; Liu, Y.; Huang, E.; Russell, T. P.; Hawker, C. Controlling Polymer-Surface Interactions with Random Copolymer Brushes. Science 1997, 275, 1458−1460. (118) Hawker, C. J.; Elce, E.; Dao, J.; Volksen, W.; Russell, T. P.; Barclay, G. G. Well-Defined Random Copolymers by a “Living” FreeRadical Polymerization Process. Macromolecules 1996, 29, 2686−2688. (119) Huang, E.; Rockford, L.; Russell, T. P.; Hawker, C. J. Nanodomain Control in Copolymer Thin Films. Nature 1998, 395, 757−758. (120) Huang, E.; Russell, T. P.; Harrison, C.; Chaikin, P. M.; Register, R. a.; Hawker, C. J.; Mays, J. Using Surface Active Random Copolymers To Control the Domain Orientation in Diblock Copolymer Thin Films. Macromolecules 1998, 31, 7641−7650. (121) Ryu, D. Y.; Wang, J. Y.; Lavery, K. A.; Drockenmuller, E.; Satija, S. K.; Hawker, C. J.; Russell, T. P. Surface Modification with Cross-Linked Random Copolymers: Minimum Effective Thickness. Macromolecules 2007, 40, 4296−4300. L

DOI: 10.1021/acs.macromol.7b00418 Macromolecules XXXX, XXX, XXX−XXX

Perspective

Macromolecules (122) Lambooy, P.; Salem, J. R.; Russell, T. P. A Method to Confine Thin Solid Organic Films between Flat Rigid Walls. Thin Solid Films 1994, 252, 75−77. (123) Kellogg, G. J.; Walton, D. G.; Mayes, A. M.; Lambooy, P.; Russell, T. P.; Gallagher, P. D.; Satija, S. K. Observed Surface Energy Effects in Confined Diblock Copolymers. Phys. Rev. Lett. 1996, 76, 2503−2506. (124) Koneripalli, N.; Singh, N.; Levicky, R.; Bates, F. S.; Gallagher, P. D.; Satija, S. K. Confined Block Copolymer Thin Films. Macromolecules 1995, 28, 2897−2904. (125) Bates, C. M.; Seshimo, T.; Maher, M. J.; Durand, W. J.; Cushen, J. D.; Dean, L. M.; Blachut, G.; Ellison, C. J.; Willson, C. G. Polarity-Switching Top Coats Enable Orientation of Sub-10-Nm Block Copolymer Domains. Science 2012, 338, 775−779. (126) Segalman, R. A.; Yokoyama, H.; Kramer, E. J. Graphoepitaxy of Spherical Domain Block Copolymer Films. Adv. Mater. 2001, 13, 1152−1155. (127) Bita, I.; Yang, J. K. W.; Jung, Y. S.; Ross, C. a.; Thomas, E. L.; Berggren, K. K. Graphoepitaxy of Self-Assembled Block. Science 2008, 321, 939−943. (128) Xiao, S.; Yang, X.; Edwards, E. W.; La, Y.-H.; Nealey, P. F. Graphoepitaxy of Cylinder-Forming Block Copolymers for Use as Templates to Pattern Magnetic Metal Dot Arrays. Nanotechnology 2005, 16, S324−S329. (129) Kim, S. O.; Solak, H. H.; Stoykovich, M. P.; Ferrier, N. J.; De Pablo, J. J.; Nealey, P. F. Epitaxial Self-Assembly of Block Copolymers on Lithographically Defined Nanopatterned Substrates. Nature 2003, 424, 411−414. (130) Davis, R. L.; Michal, B. T.; Chaikin, P. M.; Register, R. A. Progression of Alignment in Thin Films of Cylinder-Forming Block Copolymers upon Shearing. Macromolecules 2015, 48, 5339−5347. (131) Stoykovich, M. P.; Nealey, P. F. Block Copolymers and Conventional Lithography. Mater. Today 2006, 9, 20−29. (132) La, Y. H.; Stoykovich, M. P.; Park, S. M.; Nealey, P. F. Directed Assembly of Cylinder-Forming Block Copolymers into Patterned Structures to Fabricate Arrays of Spherical Domains and Nanoparticles. Chem. Mater. 2007, 19, 4538−4544. (133) Cheng, J. K.; Rettner, C. T.; Sanders, D. P.; Kim, H. C.; Hinsberg, W. D. Dense Self-Assembly on Sparse Chemical Patterns: Rectifying and Multiplying Lithographic Patterns Using Block Copolymers. Adv. Mater. 2008, 20, 3155−3158. (134) Liu, C. C.; Ramírez-Hernández, A.; Han, E.; Craig, G. S. W.; Tada, Y.; Yoshida, H.; Kang, H.; Ji, S.; Gopalan, P.; De Pablo, J. J.; Nealey, P. F. Chemical Patterns for Directed Self-Assembly of Lamellae-Forming Block Copolymers with Density Multiplication of Features. Macromolecules 2013, 46, 1415−1424. (135) Sivaniah, E.; Hayashi, Y.; Iino, M.; Hashimoto, T.; Fukunaga, K. Observation of Perpendicular Orientation in Symmetric Diblock Copolymer Thin Films on Rough Substrates. Macromolecules 2003, 36, 5894−5896. (136) Park, S.; Lee, D. H.; Xu, J.; Kim, B.; Hong, S. W.; Jeong, U.; Xu, T.; Russell, T. P. Macroscopic 10-Terabit-per-Square-Inch Arrays from Block Copolymers with Lateral Order. Science 2009, 323, 1030− 1033. (137) Hong, S. W.; Huh, J.; Gu, X.; Lee, D. H.; Jo, W. H.; Park, S.; Xu, T.; Russell, T. P. Unidirectionally Aligned Line Patterns Driven by Entropic Effects on Faceted Surfaces. Proc. Natl. Acad. Sci. U. S. A. 2012, 109, 1402−1406. (138) Dobriyal, P.; Xiang, H.; Kazuyuki, M.; Chen, J. T.; Jinnai, H.; Russell, T. P. Cylindrically Confined Diblock Copolymers. Macromolecules 2009, 42, 9082−9088. (139) Shin, K.; Xiang, H.; In, M. S.; Kim, T.; McCarthy, T. J.; Russell, T. P. Curving and Frustrating Flatland. Science 2004, 306, 76−76. (140) Ma, M.; Thomas, E. L.; Rutledge, G. C.; Yu, B.; Li, B.; Jin, Q.; Ding, D.; Shi, A. C. Gyroid-Forming Diblock Copolymers Confined in Cylindrical Geometry: A Case of Extreme Makeover for Domain Morphology. Macromolecules 2010, 43, 3061−3071. (141) Bai, W.; Hannon, A. F.; Gotrik, K. W.; Choi, H. K.; Aissou, K.; Liontos, G.; Ntetsikas, K.; Alexander-Katz, A.; Avgeropoulos, A.; Ross,

C. A. Thin Film Morphologies of Bulk-Gyroid Polystyrene- Block -Polydimethylsiloxane under Solvent Vapor Annealing. Macromolecules 2014, 47, 6000−6008. (142) Higuchi, T.; Tajima, A.; Motoyoshi, K.; Yabu, H.; Shimomura, M. Frustrated Phases of Block Copolymers in Nanoparticles. Angew. Chem. 2008, 120, 8164−8166. (143) Higuchi, T.; Motoyoshi, K.; Sugimori, H.; Jinnai, H.; Yabu, H.; Shimomura, M. Three-Dimensional Observation of Confined PhaseSeparated Structures in Block Copolymer Nanoparticles. Soft Matter 2012, 8, 3791. (144) Yabu, H.; Higuchi, T.; Jinnai, H. Frustrated Phases: Polymeric Self-Assemblies in a 3D Confinement. Soft Matter 2014, 10, 2919. (145) Detcheverry, F. A.; Nealey, P. F.; De Pablo, J. J. Directed Assembly of a Cylinder-Forming Diblock Copolymer: Topographic and Chemical Patterns. Macromolecules 2010, 43, 6495−6504. (146) Li, X.; Liu, Y.; Wan, L.; Li, Z.; Suh, H.; Ren, J.; Ocola, L. E.; Hu, W.; Ji, S.; Nealey, P. F. Effect of Stereochemistry on Directed SelfAssembly of Poly(styrene-B-Lactide) Films on Chemical Patterns. ACS Macro Lett. 2016, 5, 396−401. (147) Bai, W.; Gadelrab, K.; Alexander-Katz, A.; Ross, C. A. Perpendicular Block Copolymer Microdomains in High Aspect Ratio Templates. Nano Lett. 2015, 15, 6901−6908. (148) Liu, G.; Kang, H.; Craig, G. S. W.; Detcheverry, F.; de Pablo, J. J.; Nealey, P. F. Cross-Sectional Imaging of Block Copolymer Thin Films on Chemically Pattemed Surfaces. J. Photopolym. Sci. Technol. 2010, 23, 149−154. (149) Detcheverry, F. A.; Liu, G.; Nealey, P. F.; De Pablo, J. J. Interpolation in the Directed Assembly of Block Copolymers on Nanopatterned Substrates: Simulation and Experiments. Macromolecules 2010, 43, 3446−3454. (150) Ye, X.; Edwards, B. J.; Khomami, B. Elucidating the Formation of Block Copolymer Nanostructures on Patterned Surfaces: A SelfConsistent Field Theory Study. Macromolecules 2010, 43, 9594−9597. (151) Leibler, L. Theory of Microphase Separation in Block Copolymers. Macromolecules 1980, 13, 1602−1617. (152) Matsen, M. W. The Standard Gaussian Model for Block Copolymer Melts. J. Phys.: Condens. Matter 2002, 14, R21−R47. (153) Matsen, M. W.; Schick, M. Stable and Unstable Phases of a Diblock Copolymer Melt. Phys. Rev. Lett. 1994, 72, 2660−2663. (154) Factor, B. J.; Russell, T. P.; Toney, M. F. Grazing Incidence XRay Scattering Studies of Thin Films of an Aromatic Polyimide. Macromolecules 1993, 26, 2847−2859. (155) Wang, C.; Lee, D. H.; Hexemer, A.; Kim, M. I.; Zhao, W.; Hasegawa, H.; Ade, H.; Russell, T. P. Defining the Nanostructured Morphology of Triblock Copolymers Using Resonant Soft X-Ray Scattering. Nano Lett. 2011, 11, 3906−3911. (156) Liu, F.; Gu, Y.; Shen, X.; Ferdous, S.; Wang, H. W.; Russell, T. P. Characterization of the Morphology of Solution-Processed Bulk Heterojunction Organic Photovoltaics. Prog. Polym. Sci. 2013, 38, 1990−2052. (157) Gu, Y.; Wang, C.; Russell, T. P. Multi-Length-Scale Morphologies in PCPDTBT/PCBM Bulk-Heterojunction Solar Cells. Adv. Energy Mater. 2012, 2, 683−690. (158) Russell, T. P. X-Ray and Neutron Reflectivity for the Investigation of Polymers. Mater. Sci. Rep. 1990, 5, 171−271. (159) Penfold, J.; Thomas, R. K. The Application of the Specular Reflection of Neutrons to the Study of Surfaces and Interfaces. J. Phys.: Condens. Matter 1990, 2, 1369−1412. (160) Page, Z. A.; Liu, Y.; Duzhko, V. V.; Thomas, P.; Emrick, T.; Russell, T. P.; Emrick, T. Fulleropyrrolidine Interlayers: Tailoring Electrodes to Raise Organic Solar Cell Efficiency. Science 2014, 346, 441. (161) Choi, H. K.; Chang, J.-B.; Hannon, A. F.; Yang, J. K. W.; Berggren, K. K.; Alexander-Katz, A.; Ross, C. A. Nanoscale Spirals by Directed Self-Assembly. Nano Futures 2017, 1, 015001.

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DOI: 10.1021/acs.macromol.7b00418 Macromolecules XXXX, XXX, XXX−XXX