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Perspective pubs.acs.org/Macromolecules

50th Anniversary Perspective: Block PolymersPure Potential Christopher M. Bates*,†,‡ and Frank S. Bates*,§ †

Materials Department and ‡Department of Chemical Engineering, University of California, Santa Barbara, Santa Barbara, California 93106, United States § Department of Chemical Engineering and Materials Science, University of Minnesota, Minneapolis, Minnesota 55455, United States ABSTRACT: Block polymers have undergone extraordinary evolution since their inception more than 60 years ago, maturing from simple surfactants to an expansive class of macromolecules encoded with exquisite attributes. Contemporary synthetic accessibility coupled with facile characterization and rigorous theoretical advances have conspired to continuously generate fundamental insights and enabling concepts that target applications spanning chemistry, biology, physics, and engineering. Here, we parse the vast literature to examine the forefront of the field and identify exciting themes and challenging opportunities that portend a bracing future trajectory. This Perspective celebrates the visionary role played by Macromolecules in advancing our understanding of this remarkable class of materials.

I. PREAMBLE The title “block polymers” conjures up a host of competing images: mundane plastics and exotic high-tech devices; intuitively simple nanostructured morphologies and dauntingly complex theories; ideal polymerization mechanisms and maddeningly demanding synthetic requirements. Over the past 50 years, Macromolecules has played a central role in the development of this subject, with numerous seminal contributions that have advanced our ability to prepare, characterize, theoretically model, and technologically exploit this class of materials in myriad ways. During its first 5 years in print (1968−1972) this journal published 18 articles and communications dealing with block polymers. Since then more than 5000 papers have appeared in Macromolecules, about half during the past decade, evidencing enormous growth of the field. This Perspective peers into the future and explores selected opportunities likely to emerge in the coming years that will afford both intellectually challenging fundamental advances and exciting technological applications associated with block polymers. We begin with a brief summary of what constitutes a block polymer and take a glimpse at the extraordinary wealth of structural features that can be built into such giant molecules. Figure 1 illustrates a few of the many ways to connect chemically distinct sequences of repeat units into macromolecules that contain two or more polymeric subunits, referred to as blocks, which can be configured into linear, branched, cyclic, and hybrid molecular architectures. An explosion in the discovery and implementation of innovative polymerization methods capable of generating previously unattainable levels of man-made architectural complexity, with nearly any combination of block chemistries, represents perhaps the most enabling development in the field over the past two decades. We do not address here the synthetic © XXXX American Chemical Society

Figure 1. Illustration highlighting select examples of currently accessible molecular designs.

challenges that lie ahead. Instead, we make the assumption, supported by recent history, that any molecular architecture and combination of polymers motivated by potentially compelling structures and properties will be synthetically feasible. This includes the preparation of linear multiblock polymers containing arbitrary block sequences, star- and graftblock polymers, bottlebrush block polymers that contain one or more types of densely packed side chains emanating from a common backbone, and hierarchical configurations that combine two or more of these elements (e.g., bottlebrush multiblocks as illustrated in Figure 1). The pace of developments in recent years has been simply breathtaking, driven by numerous advances in experimental techniques enabling the synthesis and characterization of a remarkable range of materials, complemented by synergistic developments in polymer theory and computation. Following a brief review in section II of the key developments in the past leading up to the current state of knowledge and technology, Received: October 30, 2016 Revised: December 7, 2016

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DOI: 10.1021/acs.macromol.6b02355 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules based on the random phase approximation (RPA)12 that becomes exact approaching the critical point in the limit of infinite molecular weight (degree of polymerization N → ∞). (One of the authors was inspired by this landmark article to explore the fundamental consequences of macromolecular self-assembly, setting him on a journey that appears to offer no mechanism for escape.) This publication provided an analytical expression for the disordered state structure function accessed by scattering experiments and laid the foundation for subsequent incorporation of fluctuation effects by Fredrickson and Helfand in 1987.13 Semenov then made a critical link in 1985 to the strong segregation limit,14 establishing the now familiar trade-off between chain stretching and interfacial area that leads to the well-known scaling laws for domain dimensions as a function of χ and N (see section IV). Commercial development of automated rheological tools beginning in the 1970s, application of neutron and synchrotron X-ray scattering to block polymers in the 1980s, and discovery of controlled free-radical polymerization in the 1990s fueled the explosive growth of knowledge and expanding range of technological applications evident today. While there is no distinct point in time when the focus on linear AB and ABA type block copolymers began to broaden, a defining moment occurred with the introduction of star block polymers, formed by coupling living SI (or IS) diblocks using chlorosilane linking agents. Curiously, this molecular architecture produced the first network morphology,15 although it took a decade to identify the “wagon-wheel” structure as a cocontinuous topology and nearly two decades passed before the morphology was firmly established as the gyroid phase.16,17 Another important development was the introduction of more than two distinct block chemistries led by the Stadler,18 Matsushita,19 and co-workers in the 1990s, coupled with detailed structure analysis and expanded theoretical treatments. Today dozens of morphologies have been documented with flexible ABC triblock terpolymers, several of which are shown in Figure 2. The introduction of more blocks (e.g., tetrablock

we focus on several emerging themes. Molecular precision and perfection, discussed in section III, are overarching considerations in designing block polymers, made possible by the aforementioned advances in synthetic chemistry. Future progress in this field will require grappling with the daunting range of molecular options available, demanding more predictive capabilities as considered in section IV. Although most theoretical and experimental studies of block polymers assume equilibrium behavior, recent discoveries with simple diblocks and more complicated linear multiblocks underscore the prevalence and importance of nonequilibrium states. This presents challenges and opportunities associated with processing as considered in section V. We finish with a few thoughts regarding properties and products in section VI and summarize with a prospective in section VII.

II. PAST Block copolymers first appeared in the literature with little fanfare in the early 1950s in the form of large surfactant molecules denoted “Pluronics”, designed by the Wyandotte Chemical Corporation for laundering textiles.1,2 Subsequently marketed under various names including “Polaxomers”, this class of amphiphilic compounds, formed from poly(ethylene oxide) (PEO) and poly(propylene oxide) (PPO) into PEO− PPO diblocks and PEO−PPO−PEO triblocks, still find many important uses, including as key ingredients in cosmetic and pharmaceutical formulations, cell media stabilizers, in a host of medical applications, and as templating agents for the production of mesoporous materials.3 Pluronics are consequently among the most studied polymer surfactants. The discovery of living anionic polymerization by Swarc in 1956 provided a facile mechanism for economically producing block polymers (including Pluronics) on a grand scale.4 By the mid-1960s poly(styrene) (PS or S), poly(isoprene) (PI or I) and poly(butadiene) (PB or B) based thermoplastic elastomers (e.g., SIS, SBS) were available commercially, a business that that today represents one of the largest sectors of the block polymer industry, with applications ranging from adhesives to footwear to asphalt modifiers. Access to controlled polymerization spurred scientists and engineers to probe the consequences of coupling thermodynamically incompatible blocks together. Seminal work by Kawai, Hashimoto, coworkers, and others established the basic principles that govern the relationships between composition, molecular weight, morphology, and the associated length scales of bulk materials.5 Many of the experimental tools that we take for granted today, such as small-angle scattering (SAXS and SANS), dynamic mechanical spectroscopy (DMS), electron microscopy (TEM and SEM), differential scanning calorimetry (DSC), and various other techniques, were adapted by academic and industrial researchers for the purpose of characterizing the structure and properties that define these fascinating materials. Realization of the powerful opportunities presented by this emerging technology motivated new theoretical approaches. Preliminary concepts initially advanced by Meier6 in the late 1960s were followed by the pioneering work of Helfand7,8 in the 1970s based on the narrow interface approximation, which built on the conceptual framework of self-consistent field theory (SCFT) developed by Edwards.9 In the early 1990s, Vavasour and Whitmour10 relaxed the narrow interface constraint and Matsen and Schick provided the first general numerical method to solve the exact field equations without approximations.11 In 1980, Leibler published a paper in

Figure 2. Simulated morphological complexity accessible with ABC triblock terpolymers (χABN = χBCN = 35 and χACN = 15). Dotted lines are phase boundaries that are not determined exactly. Data and illustrations are courtesy of An-Chang Shi, adapted with permission from ref 20. B


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requirements only out of necessity within the confines of (in)organic methodology, still demanding careful regulation of molecular purity with otherwise potentially disastrous consequences, e.g., the thalidomide debacle. Polymer scientists, at times considered the ugly ducklings of both the chemistry and physics communities,26 must inevitably wrangle with asynchronous chain initiation and termination during conventional polymerizations, which impart significant distributions in molecular size and composition characteristic of all synthetic macromolecular products. Historically, this unavoidably hinders perfect or even moderately precise control over polymer molecular weight, dispersity, sequence, architecture, and functionalization, with consequences spanning material prediction (section IV), processing (section V), properties, and products (section VI). Yet, nature also ubiquitously leverages mixtures of inexact fibrous proteins to construct important structural scaffolds within living organisms. What balance then can or should the chemist strike when emulating evolutionary accomplishments? Here, we address pertinent topics showcasing contemporary molecular design freedom by surveying select examples from the recent polymer literature that accentuate emerging concepts involving precision. We frame the discussion by arguing that the appropriate level of precision, context dependent, is “just enough, and no more”ça suffit! Two traditionally disparate scientific communities have approached issues of macromolecular design and precision with largely orthogonal strategies.27 Synthetic biologists steeped in the traditions of nature have devised beautiful recombinant DNA techniques that facilitate the expression of nearly any artificial protein containing natural or noncanonical amino acids (Figure 3).28 These biomaterials, often joined together with

terpolymers) and block types (A, B, C, D, ...) convoluted with the architectural freedom illustrated in Figure 1 simultaneously offers unbounded opportunities for dividing space into intricately tailored nanostructures. However, as discussed in sections III, IV, and V, our enthusiasm for increasingly elaborate molecular designs is tempered by the bewildering complexity associated with predicting and actually achieving specific morphologies. Thus far, we have focused on flexible block polymers. During the past few decades many other block types have been combined into diblock and multiblock polymers including semiflexible, helical, and stiff chains, densely grafted bottlebrush polymers, and globular proteins as illustrated in Figure 1. Understanding how these hybrid compounds self-assemble in the bulk state, governed by the strong driving force to fill space at essentially uniform density, presents a host of complicating factors. The relatively simple trade-off between chain stretching and interfacial area, which dictates the spatial arrangement of flexible block polymer melts, no longer suffices to explain the states of order and disorder encountered with these mixed block moities. Moreover, these nonclassical blocks often are accompanied by complex inter- and intramolecular interactions that challenge the assumption whereby segment−segment interactions can be captured by the conceptually simple Flory−Huggins χ parameter approach. This problem is further exacerbated when charged species are introduced. Yet these materials find increasingly important technological applications, for example, in light-emitting diodes (LEDs),21 organic photovoltaic devices,22 and ion conducting membranes.23 These chemically diverse block polymers also may have large effective χ parameters that necessitate very low molecular weights in order to access the order−disorder transition. This raises several questions. What constitutes a block polymer? In particular, where does the field of block polymers transition to the realm of surfactants and liquid crystals? Application of block polymers to patterning thin films for use in the manufacturing of semiconductor devices and computer memory is fast approaching this crossover regime,24,25 which returns us to the beginning of this section and the origins of the field. There is no simple definition of what is, or is not, a block polymer, and we will not attempt to place restrictions on the label here. Not surprisingly, the self-assembly behaviors on either side of this hypothetical divide (e.g., surfactants versus thermotropic liquid crystals versus diblock copolymers) can be strikingly similar. Our purpose with this Perspective is to identify interesting and exciting topics likely to emerge as important and enabling areas for basic research and the development of technology. We have selected issues that are germane to all block polymers, regardless of molecular architecture or block types, but make no attempt to be comprehensive. Our remarks should thus not be construed as a review of the field.

Figure 3. Schematic of recombinant DNA technology used to produce perfectly monodisperse and sequence-controlled proteins. Reproduced with permission from from ref 34. Copyright 2001 The Royal Society of Chemistry.

III. PRECISION AND PERFECTION The words “precision” and “perfection” arouse remarkably varied interpretations to different audiences. Biological machinery has on the one hand evolved to produce exquisite higher-order protein structures made possible only with impeccable control over constituent amino acid sequence, undoubtedly embodying “perfection”. Correspondingly, life often tolerates virtually no variance in protein connectivity, with even single site mutations capable of generating debilitating diseases. Small molecule chemists relax these

blocky motifs, have spawned significant achievements across biology, medicine, and materials science,29 as have complementary solid-phase syntheses capable of producing welldefined peptide30 and peptoid31 oligomers. For instance, hydrophilic proteins flanked by hydrophobic residues and optionally decorated with adhesion-promoting sequences can form reversible32 shear-thinning33 hydrogels suitable for cell encapsulation, injection, and drug delivery. Engineered peptides, peptoids, and proteins are of course inherently C


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chemistry today. A sister publication describes identical parent PDMS monodisperse oligomers functionalized on each chain end with a single moiety, blends of which surprisingly selfassemble like block copolymers with periodicity scaling characteristic of strongly segregated AB diblock melts (L0 ∼ N0.67).40 This result is rather astonishing since end blocks with N = 1 hardly qualify under any definition as a “polymer”. A complementary synthetic strategy pioneered by Hawker and colleagues also generates monodisperse materials (reportedly Đ = 1.0) through meticulous column chromatographic purification of relatively narrow (but polydisperse) precursors produced via controlled polymerization (Figure 4B).41 (Note the distribution of the as-synthesized polymer, Đ ≈ 1.2, is probably considered “controlled” by most standards yet consists of many different molecular weights, even at a low average ⟨N⟩ = 8.) A major advantage of such a methodology is the potential for relatively large scale (multigram) yields. While their publications to date only describe homopolymer purification, the technique will undoubtedly evolve to block polymer separations. Currently, both iterative step growth and column chromatography are synthetically taxing, time-consuming, and limited to low molecular weights, but the nascent development of truly monodisperse synthetic macromolecules represents a turning point in polymer science. Clever utilization of automated synthetic routines might partially alleviate slow fabrication times and facilitate further percolation of discrete polymers,42 but accessing high molecular weight analogues remains a serious outstanding challenge. The unrelenting drive toward monodisperse perfection is pragmatically balanced by the possibility of useful imperfection. Moderate yet defined molecular weight distributions (roughly 1.4 < Đ < 2) retaining excellent chain end functional group fidelity achievable with controlled polymerization techniques introduce the prospect of synthesing block polymers with carefully prescribed dispersity. Such materials not only selfassemble but do so in potentially useful ways markedly different than monodisperse analogues. Systematic experimental studies on block polymers intentionally containing at least one disperse block, including AB diblocks,43 ABC triblocks,44 and ABA triblocks,45 have discerned notable effects on microphase stability and phase portraits, domain spacing, and translational order. These observations, theoretically rationalized in the framework of self-consistent field theory (see section IV), highlight the possibility of deliberately injecting molecular weight dispersity as a self-assembly design tool, a sort of precise imprecision. Significant domain dilation driven by dispersity effects can transform otherwise typical block polymers into photonic materials. A notable example with potential commercial viability is Dow’s poly(olefins) derivatives, which display vibrant colors across the visible spectrum.46 We next turn to precision in intramolecular design. A primary example is polymer architecture, which exerts considerable influence over microphase behavior and accompanying mechanical properties.47 Effectively infinite combinations of molecular architecture, connectivity, and (a)symmetry can be envisioned. Here, select examples drawn from the expansive literature associated with this topic are highlighted to emphasize the connection between microstructure and macroscopic properties. Multiblock thermoplastic elastomers and plastics are a prototypical archetype. Perhaps the least precise yet simplest to produce multiblock polymers are polyurethanes, prepared by step-growth mechanisms that couple together many sequences of hard and soft blocks into linear multiblock

restricted to a limited class of monomers that can form amide bond linkages, but the impressive power of biology cannot be understated. While protein block polymers may never penetrate commodity markets reliant on cutthroat economics, the prospects of artificial proteins could prove game-changing for polymer science as a whole in the decades to come. Significant strides in macromolecular precision and perfection have also been realized by the organic chemistry community, historically the bedrock of block polymer synthesis. The aforementioned explosion of living and controlled polymerization methodologies is beginning to provide the toolkit necessary to bridge the gap between natural and synthetic capabilities. Such unprecedented levels of human control often come at an extraordinary cost in complexity, with obvious attendant monetary implications. The remainder of this section considers block polymer precision themes framed in the synthetic polymer perspective and is primarily intended as a complement to our recent discussion35 highlighting sequence complexity effects. The familiar issue of molecular weight dispersity has seen remarkable developments associated with precision. Low dispersities produced by living and controlled polymerizations (1 ≤ Đ < 1.2) are significantly improved relative to uncontrolled analogues but a far cry from monodisperse (Đ = 1), especially in the high molecular weight limit since the corresponding standard deviation σ = Mn(Đ − 1)1/2 scales accordingly. (Here, we acknowledge the perversion36 of the terms “monodisperse” and “polydisperse” but remain resigned to their use due to overwhelming prevalence within the polymer community.) Recent advances in polymer chemistry are just beginning to provide the methodology necessary to truly probe the Đ → 1 limit (Figure 4). Meijer and co-workers

Figure 4. Matrix-assisted laser desorption/ionization (MALDI) data highlighting complementary strategies that can generate effectively monodisperse synthetic polymers. (A) Iterative coupling to produce monodisperse block polymers. Adapted with permission from ref 37. (B) Large-scale column chromatographic separation of disperse (Đ ≈ 1.2) homopolymers originating from controlled radical polymerization. Adapted with permission from ref 41.

have reported the first example of a low molecular weight and practically monodisperse (Đ = 1.000 01) AB diblock (A = poly(dimethylsiloxane), B = poly(lactide)) they coin a “discrete oligomer” (Figure 4A).37 Their iterative step-growth coupling strategy, partially rooted in both spirit and chemistry from prior dendrimer literature,38,39 is an extraordinary synthetic tour de force that underscores the amazing flexibility of synthetic D


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generally minimizes entanglements (the presence of which would cause the zero shear viscosity η0 ∼ N3.4), Rouse-like dynamics (η0 ∼ Nγ, γ ∼ 6 (1))67 coupled with extremely high molecular weights (in many cases exceeding millions of daltons) still produce incredibly slow dynamics that hamper assembly. Solvent-casting has usually been employed as a workaround,68 which can lead to useful self-assembly, albeit probably nonequilibrium. One might attempt to ascertain true equilibrium phase behavior in these situations by using a combination of solvent-casting and thermal annealing to demonstrate a path-independent final structure, although we view such conclusions as tenuous at best. A more detailed discussion on equilibrium is presented in section V. Branched architectures can be cleverly exploited to tailor molecular space-filling. For instance, the interplay between block side-chain length and backbone stretching has been implicated to rationalize domain spacing scaling69 and morphological70 trends in brush materials of varying disposition. Installing grafts which themselves are block polymers can additionally generate radial composition profiles (core−shell nanotubes71,72) of interest for templating applications.73 One could envision extrapolating these concepts to construct complex three-dimensional molecular structures by precisely tailoring local rigidity, packing, and composition using a design toolkit including backbone and side-chain lengths and stiffness, grafting density, branch connectivity and position, and block distribution. The coalescence of such intricate sequence and architectural effects represents a relatively unexplored direction in the context of block polymers, although promising reports with architecturally elaborate homopolymers suggest further evolution should be forthcoming.74,75 Advances in synthetic chemistry also translate into essentially unlimited opportunities to precisely ornament macromolecules with useful functionality. One example intimately coupled to macroscopic performance is the installation of special “triggers” that upon appropriate provocation instigate polymer degradation. Often a single carefully designed end group, combined with an appropriate monomer, is sufficient to amplify the effect of one stimuli-induced chemical reaction and fully depolymerize the parent polymer. End-capped poly(phthalaldehyde) is a classic example of such “self-immolative” behavior, driven thermodynamically upon acid-catalyzed deprotection by T > TC (where TC is the ceiling temperature). Conceptually analogous strategies have been extensively investigated with various homopolymer scaffolds, for instance in the field of drug delivery, but adaptation to block polymers is largely an outstanding challenge. As noted by Lane and co-workers,76 most examples of self-immolative polymers are to date produced via step-growth polymerization, and the production of low-dispersity block polymers is consequently difficult. This may not be intrinsically detrimental (as outlined above) but is without a doubt synthetically constraining. Furthermore, imparting clean orthogonality between block polymer processing conditions and stimuli-induced annihilation can be problematic. Explicit incorporation of structure-directing moieities (e.g., hydrogen bonding, charged) can further manipulate selfassembly. However, challenges associated with designing and imposing precise molecular interactions are underscored by issues treating even classical polymer−polymer mixing with the pervasive Flory−Huggins parameter χ. While a simplistic pairwise description of enthalpy arising from like and dislike segment interactions leads to a straightforward theoretical

polymers. Imprecision in block length distribution notwithstanding, these materials form well-defined albeit disordered morphologies with relatively narrow distributions of nanodomain sizes.48,49 Transformation of the neat monomers from low viscosity liquids to a microphase-separated solid offers numerous versatile processing routes to products including foams, adhesives, and coatings (see section IV). Kratons (ABA triblock polymers) exhibit improved mechanical properties relative to homologous AB diblocks, an advantage derived entirely from the triblock architecture, which permits bridging between discrete glassy domains through a continuous rubbery matrix. (The phase diagrams for AB and ABA block polymers are very similar.) Extension of this design principle to more complicated architectures utilizing regular multigrafted glassy side chains protruding from a rubbery backbone has yielded additional improvements in some mechanical properties like elongation at break.50 Even more intricate doubly asymmetric (in both connectivity and composition) miktoarm star polymers are theoretically anticipated to further improve thermoplastic elastomer mechanical properties through substantial perturbations to the diblock phase diagram,51,52 an exciting development perhaps rejuvenating a classic and well-studied architecture.53,54 Additional contemporary interest in miktoarm star block polymers derives from a reduced (χN)ODT relative to homologous diblocks,55 which could provide value for lithographic patterning applications in the low N limit.56 The cyclic block polymer architecture similarly generates reductions in domain spacing due to a contraction of pervaded volume.57 Advantageous properties engendered by these and other complicated architectures, while often compelling, generally come at considerable cost in synthetic complexity, a recurring theme in the balance between precision and practicality. A closely related architecture variably referred to as brush, comb, graft, and bottlebrush has recently become available with the precision necessary to generate densely attached side chains (e.g., 100%). These brush polymers are often easier to controllably synthesize than other branched polymers (in large part thanks to Grubbs’ empyrean catalyst) and can be thought of as elongated star polymers in the limit of large Nbackbone/Nside chain. The synthetic feasibility of successively stringing together homopolymer grafts to produce AB,58 ABA,59−61 and ABC62 brush block polymers has already been established with no obvious limitations preventing further sequence complexity. Steric encumbrance introduced by densely grafted side chains is typically argued to promote backbone extension relative to Gaussian coils and impart wormlike stiffness with a corresponding increase in persistence length (lP). The domain spacing scaling d ∼ Nα exhibits a concomitant increase from α = 0.66 toward α = 1, although the exact magnitude depends on specific molecular characteristics. Recent theoretical insight from Matsen63 has revealed a notable decorrelation of chain orientation as a function of distance from the block−block interface, implying moderate enhancements in chain stiffness without actually achieving the limit of truly rigidrod macromolecules. Nevertheless, the domain spacings accessed by brush block polymers, on the order of hundreds of nanometers in bulk64 and thin films,65 are perhaps useful for photonic applications66 in conceptual analogy to the aforementioned dispersity effect. A particularly vexing unsolved issue with high molecular weight brush block polymers is thermally inducing sufficient molecular mobility to achieve thermodynamic equilibrium. Even though the brush architecture E


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Macromolecules definition of χ, experiments actually measure an effective interaction parameter χeff that encompasses enthalpic and entropic effects conflated with ill-defined and poorly understood nonidealities.77 Furthermore, the value of χeff determined for identical chemical constituents with different molecular connectivity (e.g., blends vs blocks) is not identical,78 a rather unsatisfactory state of affairs given the underlying theoretical framework. Thus, even universally describing simple A−B polymer interactions is not a trivial endeavor, and accurately predicting χeff for a given pair of polymers is a process consequently guided by chemical intuition rather than quantitative physical principles. Temporarily setting issues with prediction aside (delayed further until section IV), the salient point as related to precision is that the interplay between χ (magnitude, sign), molecular design (chemistry, sequence, architecture), and any ancillary functionality (including its distribution) is poorly understood except with the most rudimentary combinations. For example, the sequence of blocks and associated χ values can drastically differentiate the phase behavior of ABC and ACB triblock polymers due to frustration effects.79 Continued extrapolation to higher levels of sequence complexity beyond roughly three blocks quickly becomes a tremendously complex proposition.35 Notwithstanding theoretical issues tackling χ, experiments indicate that the judicious decoration of block polymers can markedly influence physical behavior. Hawker and colleagues have shown that tuning intermolecular interactions with AB/ B′C blends where B and B′ contain low levels of complementary hydrogen-bonding motifs generates squarepacked arrays of cylinders.80 (One normally expects a hexagonal arrangement with P6mm plane group symmetry.) Additionally, modulating microscopic interactions with positive, negative, or zwitterionic charge can yield drastic and often unexpected results as captured in recent theoeretical work by Sing, Zwanikken, and Olvera de la Cruz.81 Park and Balsara have demonstrated that even light sulfonation of PS−PEP (PEP = poly(ethylenepropylene)) induces microphase transitions at roughly constant volume fraction.82 Probably even more shocking is the discovery that installation and deprotonation of a single end group on poly(styrene-blockethylene oxide) cleanly imparts order and transforms lamellae into hexagonally packed cylinders (Figure 5).83 These

perturbations to otherwise standard AB diblock assembly. One conclusion is irrefutable: the mean-field notion of χ is simply incapable of capturing the actual influence exerted by miniscule concentrations of structure-directing moieties. With all of the strategies available to carefully tweak molecular design, what techniques can the polymer chemist invoke to predictably prescribe mesoscale assembly? Painfully few. We are to date frustratingly constrained by only a handful of typical morphologies characteristic of nearly every block polymer (and liquid crystal and surfactant). While selfassembled structural complexity certainly scales with the number of distinct chemical constituents in a block polymer,88 the sum of known unique unit cells accessible via self-assembly is small relative to the 230 available nonmagnetic space groups (and even smaller if magnetic block polymers become popular!). Further compounding the problem is our effective inability to predictably determine a priori the influence of subtle variations in the vast parameter space available for molecular design on mesoscale structure (section IV). Envision a scenario where computer simulations unambiguously forecast optimum transport properties89 (perhaps ionic, electronic, mass, etc.) in a membrane with low but specific symmetry (e.g., some triclinic derivative) containing 10 nm pores oriented along a specific crystallographic direction. The problem is, theory and simulations struggle to make such a bold prediction, and even if one were available, synthesizing an arbitrary selfassembling material is today basically impossible (not to mention the serious challenges associated with alignment, section V). How can this glaring deficiency be surmounted? We believe one solution lies in precision. Carefully regulating hierarchical structure will likely require the convolution of sequence control and deliberately arranged inter- and intramolecular interactions in analogy to proteins. As emphasized throughout this section, chemistry provides a preponderance of tools necessary to construct complex polymer sequences in a variety of architectures and selectively adorn them with additional functional moieties. If polymer scientists were somehow able to package this plethora of strategies into a comprehensible set of design principles useful for contriving higher-order structure, they could perhaps approach the acumen regularly exploited when coaxing DNA into intricate origami patterns. Such an endeavor will undoubtedly require close collaboration with theory, a point now addressed in more detail.

IV. PREDICTION Progress in block polymer science and engineering has relied on advances in the statistical mechanical treatment of selfassembly for the past half century. Tracing the evolution of our understanding of flexible linear AB diblock and ABC triblock polymers provides insight into the synergistic nature of the relationships between theory and experiment. By the mid-1980s the phase behavior of diblocks seemed to be virtually complete. Helfand and Leibler had quantitatively accounted for the wellestablished spherical (BCC), cylindrical (hexagonal), and lamellar (striped) phases and criteria for order and disorder as a function of composition fA and the combined parameter χN.90 However, assessing whether other ordered phases might lurk in parameter space required specification of additional candidate symmetries, a limitation that largely remains today. SCFT calculations by Matsen and Schick played a critical role in establishing that the double gyroid morphology (G) exists at equilibrium based on a lower free energy than other network

Figure 5. Unexpected morphological transformations instigated by the manipulation of a single block polymer end group accentuate the importance of precision effects in self-assembly. Adapted with permission from ref 83.

observations are augmented by additional examples of individual charged groups carefully installed at the block− block interface84,85 or at chain ends86 notably modifying TODT and chain dynamics. Such effects have been sporadically investigated and anticipated87 theoretically, but no definitive guidelines have materialized which fully justify the often strong F


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Macromolecules structures such as the double diamond and the plumber’s nightmare.11 Subsequent analysis using the mean-field treatment provided illuminating insights into the roles of network connectivity and interfacial curvature in the free energy competition that favors 3-fold over 4- and 6-fold network functionality.91 Identification of the pentacontinuous core− shell double gyroid in a poly(styrene)-b-poly(isoprene)-bpoly(dimethylsiloxane) (SID) triblock terpolymers relied directly on SCFT calculations for interpretation of TEM micrographs.92 In 2002, another network phase dubbed O70 (Fddd space group symmetry) was found in poly(styrene)-bpoly(isoprene)-b-poly(ethylene oxide) (SIO) triblocks,93 which was also subsequently accounted for with SCFT calculations.94 The power of theory was demonstrated by extension of the ABC result to the AB diblock limit, leading to prediction of the O70 phase in a slender region of phase space where the gyroid window approaches the critical point associated with the meanfield treatment. Within two years this prediction was confirmed by Takenaka and co-workers based on SAXS measurements conducted with PS−PI diblock copolymers.95 Surprisingly, there was more to come with simple flexible diblock copolymers. In 2010, Lee et al. reported the discovery of the Frank−Kasper σ-phase in a poly(isoprene)-b-poly(lactide) (PI−PLA) specimen, a low symmetry ordered structure with 30 nearly spherical particles per unit cell.96 Four years later, An-Chang Shi and co-workers showed that this equilibrium morphology is accounted for by SCFT, revealing that the σ-phase emerges between the classical BCC and cylindrical (Hex) states at sufficiently high conformational asymmetry, (bA/bB)2 ≥ 1, and elevated segregation strength, where the A block forms the pointlike particles and b represents the statistical segment length.97 This important finding was nearly anticipated by Grason and co-workers98,99 in elegant publications that predicted the A15 structure, another Frank− Kasper phase, in structurally asymmetric ABn branched copolymers with n ≥ 2. Clearly, the theoretical machinery was in place to predict the existence of the σ-phase before it was found experimentally. What was lacking was the insight to evaluate this particular phase symmetry. Herein lies the current dilemma in predicting equilibrium phase behavior based on a given molecular design: What structural candidates should be evaluated in the free energy competition? Even for the restricted case of relatively simple nearly spherical particles there are dozens of plausible packing scenarios. Nevertheless, Grason demonstrated the predictive power of SCFT for elucidating how molecular architecture translates into mesoscopic structure. Placing such computational tools directly into the hands of practicing experimentalists should accelerate the discovery of new morphologies.100 Many applications of block polymers require processing the material as a viscoelastic liquid in the disordered state followed by ordering upon cooling, driven either by large enough χN > 0 or crystallization of one or more blocks.101 Fluctuation effects become important in the vicinity of the order−disorder transition (ODT), and the associated consequences on phase behavior are amplified as the molecular weight is reduced. Recent work by Morse and co-workers102 has shown that block polymer melts appear to conform to a corresponding states principle, where an additional parameter, N̅ = Nv2b6, first introduced by Fredrickson and Helfand,13 accounts for deviations from the SCFT treatment (Figure 6); mean-field behavior is recovered in the limit N̅ → ∞. These authors have obtained quantitative agreement between several different

Figure 6. (A) Corresponding states principle accounts for universal (χN)ODT versus invariant degree of polymerization N̅ at low molecular weights and symmetric volume fractions where strong deviations from mean-field behavior dominate the phase behavior. Points represent molecular simulations; dotted, solid, and dashed curves are ROL theory, one-loop fluctuation theory, and self-consistent mean-field theory, respectively. Reproduced with permission from ref 102. Copyright 2014 the American Physical Society. (B) Illustration spanning the homogeneous disordered state at χN ≪ (χN)ODT, fluctuating disordered state at χN < (χN)ODT, weakly segregated lamellae at χN > (χN)ODT, and strongly segregated lamellae at χN ≫ (χN)ODT.

molecular simulation models and a renormalized one-loop (ROL) theory that incorporates an effective interaction parameter (χe) that is a nonlinear function of the segment− segment interaction strength α within the context of Flory− Huggins theory χ ≅ zα/kT. Comparison of the predictions of the ROL theory and simulation results with experimental data obtained from symmetric, low molecular weight, lamellarforming poly(isoprene)-b-poly(lactide) (PI−PLA) diblock copolymers, including a direct assessment of the heat of transition at the ODT, provides encouraging, although not definitive evidence that fluctuation effects can be accounted for in all low molecular weight flexible diblock copolymers (N̅ < 500).103 Subsequent dispersity corrections to this treatment by Beardsley and Matsen produced virtually exact agreement.104 We look forward to the extension of this work to asymmetric and ultimately multiblock polymers. Much progress has been made over the past decade that moves block polymer theory far beyond the now classical SCFT. Fredrickson’s seminal book “The Equilibrium Theory of Inhomogeneous Polymers” provides an outstanding introduction to developments over the past decade that promise to revolutionize the modeling and simulation of mesoscopic soft materials.105 Ideally, predictive approaches that obviate the need to initially specify an ordered state symmetry yet evolve to local or global minima in the overall free energy surface for particular block architectures and sets of molecular parameters are most desirable. Equally if not more powerful would be cracking the so-called “inverse design” problem. Given a target morphology and the available molecular design tools as explicated above, how can the polymer chemist distill the effectively infinite parameter space (including block sequence, composition, dispersities, χij, statistical segment lengths, etc.) into a synthetically tractable formulation to achieve the desired result? Recent simulation breakthroughs employing “swarm G


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(DDQC) phases.115 Limited chain stretching at such low molecular weights, where the contour length approaches the persistence length, may be better modeled using rod−coil concepts including wormlike chains, even though both blocks are normally considered to be flexible polymers. Current theory and most simulations rely on the classical Flory−Huggins χ parameter to account for pairwise segment− segment interactions. Unfortunately, as alluded to in section III, results obtained for χ(T) using different experimental approaches generally are not quantitatively self-consistent. For example, expressions obtained by SANS from single phase blends of homopolymers do not necessarily account for the molecular weight required to place the ODT for the corresponding diblock at a precisely specified temperature.78 Although the gross magnitude of χ can be qualitatively estimated for most pairs of polymers using existing experimental data or solubility parameters, the variability of literature values, even for extensively studied systems like PS− PI, PS−PEO, and PI−PEO, presents a real problem in predictive design. This problem is magnified with multiblock polymers containing three or more types of repeat units, for example PS−PI−PEO, where subtle variations in the relative magnitude of the interaction parameters (e.g., χSI vs χSO vs χIO) produce dramatic changes in the predicted ordered phases as a function of composition and temperature. Introduction of charges species to the blocks, a topic currently being developed with important technological implications (see sections III and VI), presents additional levels of complexity. We are not sanguine about the prospects for rectifying this dilemma in the near future. There is no reason to expect that block polymer phase behavior can be accurately represented by collections of individual interaction parameters that do not depend on composition and can be used interchangeably without regard for the detailed nature of the interactions at the segment scale. This does not generally work with low molecular weight liquids. For example, the van der Waals equation of state (EOS) captures the qualitative character of liquids and gases, including (mean-field) critical behavior based on a single interaction term.116 However, representing even relatively simple low molecular weight molecules on a unified corresponding states diagram requires at least one additional parameter (e.g., acentric factor) that accounts for non-idealities such as differences in molecular geometry. Industrially relevant equations of state (e.g., Peng−Robinson) contain many more empirical parameters. Capturing even modestly complex vapor−liquid, liquid−liquid, and liquid−solid equilibrium phase behavior in binary mixtures, such as azeotropes and eutectics, requires enthalpy of mixing models with more than one parameter (e.g., a two-constant as opposed to a oneconstant Margules expression116). Flory−Huggins type single parameter (χ) regular solution theories are just inadequate. Moreover, relatively subtle differences in the χ parameters in ABC triblock copolymers can produce massive changes in phase behavior. Multiblock polymer morphology is considerably more delicate and susceptible to variations in block interactions, which scale as χ ∼ N−1 at the ODT (see section V), than the phase behavior of low molecular weight binary or ternary liquid mixtures. Nevertheless, molecular simulations will become increasingly important for understanding block polymer self-assembly, and polymer−polymer phase behavior in general, as computational algorithms become more sophisticated and more efficient, powered by inexorable gains in computer speed and memory

intelligence optimization” indicate realizing this lofty goal could be on the horizon (Figure 7). A further review of the

Figure 7. Swarm multiblock polymer inverse optimization illustrating convergence on a targeted line and spot morphology by varying a blend (composition, Ni) of two triblocks and a homopolymer. Reproduced with permission from ref 108.

remarkable advances in the past few years that employ field theoretic simulations (FTS), which incorporate fluctuations into SCFT, is beyond the scope of this Perspective. We refer the reader to the literature, which portends enabling developments in the near future.106,107 Adapting SCFT to other block motifs is not simple. Stiff and semiflexible blocks present challenges that cannot be addressed using the Gaussian chain assumptions associated with flexible macromolecules.109−111 When the block persistence length approaches or exceeds the segregated domain size, orientational interactions between blocks must be accounted for, introducing additional parameters and computational difficulties. Considerable progress has been made with rod−coil diblock copolymer melts, but extension to other molecular architectures awaits future developments. A recent contribution by Matsen and co-workers highlights the difficulties of modeling bottlebrush block polymers, even for simple 1-dimensional lamellae.63 Dense arrays of flexible side chains produce radial stretching that couples to the backbone configuration directly influencing the overall periodic length of the morphology. Predictive calculations that identify the consequences of varying the density, length, and chemical makeup of the chains emanating from the backbone of individual blocks and anticipating the collective effects of connecting such bottlebrushes into precisely prepared hierarchical block architectures would provide a terrific design tool. Some progress has already been made in delineating bottlebrush homopolymer behavior using polymer physics scaling arguments.112 Recent trends in the application of block polymers to the patterning of thin films has driven the development of high χ and low N materials, literally encroaching on the crossover regime to surfactants and liquid crystals. de Pablo and colleagues113,114 have demonstrated the utility of computational approaches to guiding the design and processing of block copolymers when subjected to alignment on chemically patterned surfaces and the control of topological constraints (see section V). Polymers with N < 50 are likely to deviate from mean-field theory beyond what is captured by the corresponding states principle for flexible blocks. For example, asymmetric (f L < 0.3) diblocks containing short (N ≈ 10) poly(lactide) (L) blocks exhibit fascinating low symmetry ordered structures including the Frank−Kasper σ and dodecagonal quasicrystal H


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properties121 and asymmetric swelling with solvents selective toward the matrix B blocks.122 Flow fields are especially efficient at driving isotropically ordered (i.e., polycrystalline) block polymers into states of long-range order. Mathis, Hadziioannou, and Skoulios first demonstrated that oscillatory shearing is remarkably effective at generating monodomain materials,123 and this seminal work spawned numerous investigations that probed the roles of molecular architecture, microdomain geometry, frequency, strain amplitude, and proximity to order−disorder and order−order transitions in generating unprecedented levels of long-ranged order and microdomain perfection in bulk specimens. Particularly illuminating in situ experiments, enabled by mounting shearing devices in neutron and X-ray beams, led to the discovery that the orientation of lamellae, perpendicular and parallel (and even transverse) to the plane of shear, could be controlled by the processing history and various molecular variables including “viscoelastic contrast”.124 More recently, Segalman,125 Osuji,126 and colleagues have shown that magnetic fields can be exploited for the purpose of orienting block polymers based on diamagnetic anisotropy between domains. Although the coupling between the orientational order parameter and magnetic field is much weaker than with hydrodynamic flow, the block polymers can be tuned to respond locally through covalent127 or noncovalent128 incorporation of liquid crystalline moieties with anisotropic field susceptibility. Even poly(styrene)-b-poly(4-vinylpyridine) (PS−P4VP) diblock copolymers can be aligned in a sufficiently strong magnetic field.126 At the opposite extreme, symmetric disordered block polymers near the ODT exist in a fluctuating state (see Figure 6B) that closely resembles a bicontinuous microemulsion.129 Similar morphologies are prevalent in segmented multiblock polymers such as poly(urethane)s where the molecular architecture, including dispersity in block length and overall molecular weight, presents insurmountable barriers to the formation of long-range ordered structures. We do not fully understand the thermodynamic and kinetic obstacles that mediate transformation of defective, nonequilibrium morphologies into lower free energy states.130,131 Microphase-separated block polymers (and soft materials in general) are governed by complex free energy surfaces that depend on many variables. Line defects, such as dislocations and disclinations in cylindrical and lamellar structures, may be locally stable but globally metastable and difficult or impossible to anneal away.132 Moreover, the mechanisms responsible for transitioning between states of order and disorder are not well understood. Domain fusion and fission, the operative mechanisms for converting a spherical to cylindrical morphology and vice versa, do not require interdomain diffusion, much like what occurs when a water−oil suspension coarsens into macroscopic phases by droplet coalescence. On the other hand, transitions between disorder and various ordered states with point particles (nominally spheres) are entirely dictated by chain exchange mechanisms that require diffusion of molecules between domains. Recent work with sphere forming diblock polymers and tetrablock terpolymers has revealed the formation of periodic crystals and aperiodic quasicrystals where symmetry selection is mediated by chain exchange at rates that are critically dependent on the extent of supercooling below TODT, i.e., the segregation strength.115,133,134 Surprisingly, rapidly cooling the disordered melt more than about 30 degrees below TODT leads to a kinetically trapped state of liquidlike particle packing (denoted LLP) that may be best characterized as a soft

capacity. We are now on the cusp of realizing simulated phase behavior of block polymers with nearly atomistic detail. Recent progress by Siepmann et al.117 in modeling mixtures of olefin oligomers using Monte Carlo simulations offers some hope that χ parameters in the future may be predictable based on detailed molecular structures. We close this section on a practical note, summarizing a useful approach to establishing χ(T) experimentally for the purpose of engineering diblock copolymers, and possibly multiblocks, with specific ODT temperatures.63 Order and disorder, hence TODT, can be precisely established in a diblock copolymer using small-angle scattering measurements or the linear dynamic elastic modulus G′(T) determined in the low frequency regime. With two or more sufficiently spaced TODT values, obtained from diblocks prepared with the same composition but different molecular weights, χ(T) can be modeled using an assumed form for (χN)ODT. In practice, we usually accept the mean-field result, (χN)ODT = 10.5 for fA = 1/ 2. Plotting 1/TODT versus the degree of polymerization N (calculated with a specified but arbitrary statistical segment reference volume v) and assuming χ = AT−1 + B leads to values of the coefficients A and B. Obviously this approach ignores fluctuation effects, i.e., N̅ . (For some purists this approach is heretical; a theory colleague recently used the term “garbage”. Ideally, a proper fluctuation corrected model should be employed (e.g., the ROL theory102), which requires values for bA and bB.) Nevertheless, with some exceptions,118 the limited range in N that produces experimentally accessible TODT values (intermediate to the glass or melting transition and degradation temperatures, e.g., 100−250 °C for PS−PI) usually results in a linear plot and an expression for χ(T) that is very useful for synthesizing diblock copolymers with predictable ODT temperatures.119 In principle, this procedure can be repeated as a function of composition. We look forward to the development of a more complete corresponding states theory, including the effects of composition, which could be implemented with this experimental method.

V. PROCESSING Understanding the structure and properties of materials generally begins with equilibrium concepts. In practice, the desired functions of such materials almost inevitably involve nonequilibrium states. Perhaps the most familiar macromolecular examples are drawn from biology, where the lifeenabling behaviors of proteins rely on hierarchical structures that often do not reflect the lowest free energy state under physiological conditions. Innumerable less dramatic but technologically essential nonequilibrium structures define the field of materials science and engineering, highlighted by processes that generate useful metallic alloys such as steel, and nanostructures semiconductor devices that include integrated circuits and computer memory. Inevitably, optimized performance is intricately coupled to tailored levels of structural order, ranging from disordered nonequilibrium glasses to nearly perfect single crystals. Soft materials present unique challenges in this regard and block polymers offer unparalleled opportunities to understand and exploit structural manipulation through the control of morphology via processing. Keller and co-workers first described the production of macroscopically aligned SBS triblock copolymers through melt extrusion.120 SAXS patterns revealed formation of a “single crystal” of hexagonally packed poly(styrene) cylinders, and subsequent work demonstrated highly anisotropic mechanical I


Perspective

Macromolecules glass. Ordering is thwarted due to the extinction of block polymer chain exchange speculated to accompany strong segregation of the blocks at T ≪ TODT. Remarkable similarities between the complex phases discovered in diblocks and multiblocks, such as the Frank−Kasper σ-phase and a dodecagonal quasicrystal (DDQC), and the phase behavior of many metals and alloys, including vitrification upon rapid cooling and various diffusionless (e.g., Martensitic) transitions, offer intriguing opportunities to probe the fundamental nature of symmetry breaking in both hard and soft materials. Introduction of more than two chemical block types (e.g., ABC, ABAC, ABCD, etc.) produces unavoidable asymmetries in the associated interaction parameters (χij), which can complicate accessing equilibrium morphologies. For example, the sequence of solubility parameters for poly(isoprene), poly(styrene) (S), and poly(ethylene oxide) (O), δI = 8.1, δS = 9.1, and δO = 10.1 (in units of [cal/cm3]1/2), leads to χIS ≈ χSO ≈ 1/4χIO. Hence, an SIO triblock terpolymer will order in two stages when cooled from a homogeneous state at high temperature: first, the O block segregates from the mixed SI domain to reduce I−O contacts followed by segregation of S and I at sufficiently low temperature.135 Depending on the overall composition (e.g., f I, f S, and f O = 1 − f I − f S), noncovalently bonded S−O interfaces may offer the lowest free energy morphology. Equilibrating a metastable morphology formed during cooling or solvent casting requires breaking and re-forming interfaces and collectively moving individual blocks through unfavorable states, which is greatly exacerbated by multiblock and nonlinear molecular architectures that require coordinated rearrangements of blocks. As an additional speculative scenario, one might imagine inducing long-range hexagonal core−shell cylinder order by imposing large-amplitude oscillatory flow in an ABC triblock terpolymer that exhibits an equilibrium network morphology. Cessation of flow would then allow an order−order transition through interfacial rearrangements, guided by epitaxial relationships, resulting in an anisotropic material with different domains that could be individually continuous in 1, 2, or 3 dimensions. Realizing such advanced materials will require close collaboration between synthetic chemists, experts in structural and property characterization, processing, and theory and simulation. Sorting out the correct combinations of molecular architecture, chemical building blocks, and processing conditions from the bewildering array of structural and processing variables seems like a daunting design challenge. Yet the rewards could be rich. Imagine mechanically robust, perhaps melt blown, films that selectively transport and separate gases or ions with high flux and efficiency. Connecting nonlinear rheological factors, especially extensional flows,136 with the thermodynamic behavior of self-assembled bulk morphologies through experimentation and theoretical modeling is a prerequisite for innovation in this area. The future has already arrived in the limit of two dimensions. Block polymer thin films offer an astonishing ability to controllably pattern surfaces at length scales rapidly approaching 1 nm.137−139 Surface chemistry140−142 or topological143,144 guiding fields can be used to control in-plane alignment and accomplish “directed self-assembly” of lamellae (Figure 8A),145 cylinders,146 spheres,147 and in principle other morphologies. Application to the fabrication of semiconductor and magnetic devices through clever pattern transfer148−152 and replication strategies153,154 is critically dependent on minimizing the level of defects, including struts that bridge domains, dislocations,

Figure 8. Processing block polymers into highly (A) perfect158 or (B) irregular159 patterns can generate remarkably useful properties. Both scale bars represent 100 nm. (A) Thin film lamellae of PS−PMMA aligned on a chemical prepattern. Adapted with permission from ref 158. (B) An ionic liquid-containing polymer electrolyte membrane with cocontinuous but globally disordered domains exhibits both high ionic conductivity and storage modulus. Adapted with permission from ref 159.

and fluctuating interfaces. Currently, state-of-the-art processing utilizing the most mature materials has achieved defect densities (referred to as “defectivity” in the field) below 30 per cm2, a level of perfection unattainable just a few years ago.155 Robustly measuring and even reliably quantifying such miniscule levels of imperfection, arising from a complicated convergence of material attributes and processing steps, with sufficient statistics to extrapolate toward a commercial reality remains a formidable challenge.156 As mentioned above, an incomplete understanding of the mechanisms and pathways at play during molecular reorganization creates uncertainty in both the true lower limits of defectivity and those realistically achievable in practice. Yet, modern block polymer design, synthesis, and processing strategies, qualified by advanced morphological characterization tools (e.g., AFM, tomographic TEM, GISAXS, metrology, etc.) and guided by self-consistent mean-field theory and computer simulation, offer hope that this technology may impact industry and society in the coming years.157 Expansion of these methods to 3-dimensional structures seems feasible. Imagine combining holographic patterns at the micrometer scale through illumination with light, which locally disorders or reorders regions of a bulk block polymer,160 in combination with dynamically imposed shear or magnetic fields that result in domain alignment. Revisiting again the aspiration contemplated in section III, the fanciful concept mingling light, order, and disorder and external fields could lead to hierarchical and dimensionally-specified control over electronic, ionic, and molecular transport as well as optical and mechanical properties. Engineering ordered phase grain boundaries, much like what is done with hard materials, will play a key role in realizing such material objectives. Blending block polymers and homopolymers with precisely controlled compositions, molecular architectures, molecular weights, and dispersities (see section III) will enable the crafting of almost any local interfacial structure, provided the appropriate predictive tools are available (see section IV). Glimpses of what is possible have been provided in two dimensions through the addition of judicious amounts of homopolymer to diblock copolymers in order to accommodate sharp bends with perpendicular lamellae in thin films.161 On the other hand, processing strategies that exploit nonequilibrium states offer terrific opportunities to tailor the structure and properties of final products. A practical example is polymerization-induced phase separation, which can produce cocontinuous but irregular three-dimensional structures on J


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yet versatile multiblock elastomers and plastics, will contribute to grow in the coming years and constitute the greatest volume and revenue share of the block polymer market for the foreseeable future. However, such commodity applications based largely on mechanical properties and high volume processing capture only a small portion of the innovations occurring in the field. We believe block polymers will play a key role in the development of sustainable polymers derived from natural sources such as sugar and cellulose.172 Today a very small fraction ( pKa,PAA induces deprotonation of PAA, with electrostatic repulsion driving chain elongation and pore blockage. (B) Dry membrane. (C) pH = 6.88. Reproduced with permission from ref 199. Copyright 2014 Elsevier. Figure 11. (A, B) Scanning electron micrographs of a Taxus Liberté stent comprising a stainless steel strut coated with drug-eluting PS− PIB−PS triblock polymer: (A) representative scanning electron micrograph (SEM) of an expanded stent with the white arrow demarcating the SIBS polymer; (B) magnified SEM with the black arrow indicating the underlying stainless steel strut. Reproduced with permission from ref 202. Copyright 2010 Elsevier. (C, D) Atomic force micrographs (scale bars 200 nm) of a polymer-coated stent surface: (C) PS−PIB−PS only; (D) PS−PIB−PS loaded with 8.8 wt % Paclitaxel. Reproduced with permission from ref 203. Copyright 2005 Elsevier.

Here, high cost will not necessarily deter commercial viability, as the active material could represent a trivial fraction of the overall mass of a separation device. Imagine a 100 nm film of exquisitely fabricated block polymer supported on a macroporous hollow fiber, emulating the configuration of traditional hollow fiber separation units. The total amount of block polymer required per square meter of active surface area is around 0.1 g! We envision enormous opportunities in this field, perhaps augmented by the strategic placement of catalytically active nanoparticles within the film for applications in reactive separations. Clearly, the combination of imaginative science and engineering, fueled by creative chemistry, makes these value-added possibilities extraordinarily attractive for development. And there are fascinating templating technologies that exploit block polymers in structuring nanoporous inorganic materials such as zeolites.200 While the block polymer guides formation of the transient mesostructure, its presence disappears with sintering, a bit like the remnant Cheshire Cat’s smile in Alice in Wonderland. As we work our way up the value chain, a few biomedical materials applications should be highlighted. Today many vascular stents are covered with drug eluting poly(styrene)-bpoly(isobutylene)-b-poly(styrene) (PS−PIB−PS) triblock copolymer (Figure 11).201 This passive application exploits the unique combination of drug solubility in PIB and the mechanical and adhesive properties of the material, somewhat like a plasticized pressure-sensitive adhesive. Innumerable other biomedical applications rely on either long-term stability or controlled degradation coupled with demanding linear and nonlinear elastic and plastic mechanical properties. Medical implants are subject to hostile biological environments that challenge materials design. Nanoscale morphologies may expose chemical reaction pathways that are not encountered in bulk materials. Recent experience with poly(ether urethane)-b-poly(dimethylsiloxane) (PEU−PDMS) multiblock polymers highlights some of the challenges facing the industry. Combining PEU and PDMS blocks provides attractive mechanical properties for use in coating pacemaker leads and several commercial products are marketed based on these materials. However, when configured with 3−4 nm size microphase-separated PDMS and PEU domains, an unanticipated hydrolysis reaction occurs,204 which over time degrades the molecular weight and simultaneously the mechanical performance;205 PEU and PDMS alone experience much slower hydrolysis.206 We suspect that modest concentrations

(ca. 2 wt %) of water associated with the PEU blocks expose PDMS at the domain interfaces to a chemical purlieu never encountered in bulk PDMS, resulting in slow cleavage of the soft blocks under physiological conditions, ultimately compromising performance. Medical devices intended for long-term implantation require new block chemistries capable of resisting oxidative and hydrolytic degradation over at least 10 years. As emphasized repeatedly, such materials must be engineered to simultaneously account for structural, processing, and ultimate mechanical performance. Most currently marketed block polymer-based products are characterized by a relatively narrow range of domain dimensions, focused between roughly 5 and 50 nm. Recent advances have pushed these limits to smaller and larger scales, and we expect these trends to continue, driven by technological demands, enabled by precise synthesis (see section III) and guided by theory and simulation (section IV). Aforementioned progress in two-dimensional patterning is likely to be extended down to about 1 nm, beyond which the concepts associated with block polymers become moot. At the other extreme, bottlebrush and hierarchically assembled block polymers have reached into the optical regime, generating regular periodic structures capable of functioning as one-dimensional photonic crystals.207 Precisely controlling self-assembled morphology at the micrometer scale and actually reproducibly processing such materials is challenging. Flexible linear blocks require impractically high molecular weights to achieve domain periods much greater than 100 nm. Strong segregation mean-field theory shows that a monodisperse PS−PI diblock requires Mn ≈ 3.2 × 105 g/mol to reach a lamellar period of 0.1 μm, which is unworkable for several reasons. For example, such a diblock would be nearly impossible to process (placing TODT at 200 °C requires Mn ≈ 2 × 104 g/mol), and higher molecular weights become synthetically challenging. M


Perspective

Macromolecules Tailored branching, including bottlebrushes, offers molecular design strategies that can extend the spatial range of individual macromolecules, while avoiding intractable molecular dynamics by limiting entanglements. Yet, even these advantages are tempered at ultrahigh molecular weights (section III) like those necessary to push Bragg reflection significantly further into the infrared. Precise placement of chemically dissimilar blocks endowed with suitable functionality could guide the formation of three-dimensional order at the micrometer scale. We expect that the strategic placement of charged groups will be especially useful in manipulating the strength and range of interactions between blocks. Implementation of such an approach will require predictive computer simulations to guide the synthetic chemist.

Biographies

VII. PROSPECTIVE

Christopher M. Bates earned a B.S. degree in Chemistry at the University of WisconsinMadison in 2007 and received a Ph.D. from The University of Texas at Austin in 2013 under the guidance of C. Grant Willson. After a postdoc with Robert H. Grubbs at the California Institute of Technology, Christopher recently moved to the University of California, Santa Barbara as an Assistant Professor in the Materials and Chemical Engineering Departments.

The field of block polymers has witnessed extraordinary developments since its inception more than 6 decades ago. We optimistically identify emerging opportunities spanning the gamut of topics that will propel this discipline in the coming decades, including synthesis, structural and dynamic characterization, theory and simulation, and translation to products. As should be clear from this Perspective, we view these specialties as increasingly interdependent. Blurring the artificial boundaries separating chemistry, engineering, materials science, and physics will be increasingly necessary to maintain innovative momentum. These essential constituents must continue to evolve from isolated specializations into well-integrated collaborative research endeavors. We are encouraged to see block polymers percolating through the scientific community well beyond specifically trained practitioners. Modern simplicity in executing controlled polymerizations, the convenience of synchrotron radiation facilities, commercially available rheological characterization tools, astounding developments in electron microscopy, and tractable theory are just a few of the factors that make block polymer discovery accessible to specialists and nonspecialists alike. Overcoming the challenges identified in this Perspective portends additional decades of exciting advances that will be made available to the scientific and engineering communities through the continued journalistic leadership of Macromolecules.

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Frank S. Bates is a Regents Professor and a member of the Chemical Engineering and Materials Science department at the University of Minnesota. He received a B.S. in Mathematics from SUNY Albany and M.S. and Sc.D. degrees in Chemical Engineering from MIT. Between 1982 and 1989 Bates was a member of the technical staff at AT&T Bell Laboratories and then joined the University of Minnesota, where he served as department head from 1999 to 2014. He is a member of the US National Academy of Engineering and the American Academy of Arts and Sciences.

AUTHOR INFORMATION

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Corresponding Authors

*E-mail [email protected] (C.M.B.). *E-mail [email protected] (F.S.B.).

ACKNOWLEDGMENTS This work was supported by (FSB) the National Science Foundation through the University of Minnesota MRSEC under Award DMR-1420013. C.M.B. thanks UCSB for funding and the Materials Research Laboratory, a Materials Research Science and Engineering Center (MRSEC) under support from the National Science Foundation (DMR-1121053). We gratefully acknowledge Brian Long and Peter Allen for rendering the cover artwork that accompanies this manuscript and the reviewers for excellent suggested improvements.

ORCID

Frank S. Bates: 0000-0003-3977-1278 Notes

The authors declare the following competing financial interest(s): Frank S. Bates has royalty interests in, and serves as Chief Financial Officer and on the Board of Directors of Valerian Materials, a company involved in the commercialization of β-methyl-δ-valerolactone. The University of Minnesota also has equity and royalty interests in Valerian Materials. These interests have been reviewed and managed by the University of Minnesota in accordance with its Conflict of Interest policies.

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REFERENCES

(1) Mankowich, A. M. Micellar Molecular Weights of Selected Surface Active Agents. J. Phys. Chem. 1954, 58, 1027−1030.

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