Aperiodic “Bricks and Mortar” Mesophase: a New Equilibrium State of

Article pubs.acs.org/Macromolecules

Aperiodic “Bricks and Mortar” Mesophase: a New Equilibrium State of Soft Matter and Application as a Stiff Thermoplastic Elastomer Weichao Shi,† Andrew L. Hamilton,† Kris T. Delaney,† Glenn H. Fredrickson,*,†,‡,§ Edward J. Kramer,†,‡,§,○ Christos Ntaras,# Apostolos Avgeropoulos,*,# Nathaniel A. Lynd,⊥,¶ Quentin Demassieux,∥ and Costantino Creton∥ †

Materials Research Laboratory, ‡Department of Chemical Engineering, and §Department of Materials, University of California at Santa Barbara, California 93106, United States # Department of Materials Science and Engineering, University of Ioannina, University Campus, Ioannina, Greece 45110 ⊥ McKetta Department of Chemical Engineering, University of Texas at Austin, Austin, Texas 78712, United States ¶ Materials Sciences Division, Lawrence Berkeley National Laboratory, Berkeley, California 94720, United States ∥ Laboratory of Soft Matter Science and Engineering, ESPCI Paristech-CNRS-UPMC, 10 rue Vauquelin, 75005 Paris, France S Supporting Information *

ABSTRACT: A new thermodynamically stable, aperiodic “bricksand-mortar” (B&M) cellular mesophase structure is reported in PS1b-(PI-b-PS2)3 miktoarm copolymer and PS homopolymer blends [PS1, long polystyrene; PI, poly(isoprene); PS2, short polystyrene], where PS comprises discrete hard “bricks” and PI the continuous soft “mortar”. The mesophase is unique in its extreme domain volume fractions, its lack of positional order, and quasi-long-range orientational order. On the basis of this unusual mesophase structure, a series of PS-based thermoplastic elastomers are realized, combining rigidity from an exceptionally high content of discrete glassy PS domains (up to 82 wt %) and high extensibility with recoverable elasticity from a low content of continuous rubbery PI (down to 18 wt %). The new elastomers show sharp yielding behavior while maintaining good elasticity at large strains. Tensile-SAXS experiments reveal that voiding plays an important role for the mechanical behavior and voids can open/close reversibly with/without loading. Plastic deformation only results in a slight loss of recoverable elasticity.

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INTRODUCTION The self-assembly of block copolymers (BCPs) has been of considerable interest over the past decades, including neat systems,1 blends with homopolymer,2 and more general BCP alloys.3 While many different types of structural motifs are possible, rare are thermodynamically stable nanoscale cellular morphologies, where a high volume fraction discrete phase is well dispersed in a low volume fraction continuous matrix. Such materials could find a wide range of applications in food and colloid science (such as high internal phase emulsions), photolithography, catalyst supports, separation membranes, energy and electronic devices, and thermoplastic elastomers (TPEs).4−14 TPEs are an important class of materials, which exhibit elastic mechanical properties at room temperature and can be easily processed at elevated temperatures. Traditional polystyrenebased TPEs contain a low volume fraction of discrete glassy PS domains (typically less than 0.3) embedded in a continuous rubbery matrix, rendering the material elastic, but soft.12,13 From a mechanics perspective, if glassy PS domains could be maintained as discrete objects to high volume fraction, while sustaining a continuous low volume fraction rubbery domain, © XXXX American Chemical Society

then we would expect that such a nanocellular material would possess the unusual combination of high modulus from the discrete glassy phase and recoverable elasticity from the soft continuous matrix, representing a new class of TPEs.12−14 Unfortunately, such a nonconventional state of self-assembly has never been achieved by equilibrium processing of linear BCP architectures.15−18 In the conventional case of a linear AB diblock or symmetric ABA triblock copolymer, it is rarely possible to realize a cellular mesophase such as the hexagonal phase (HEX) with “hard” A cylinders beyond an A volume fraction, fA, exceeding 0.3; for larger fA, the A domains become continuous in at least two dimensions, destroying the possibility of recoverable elasticity. When such a linear block copolymer is blended with a homopolymer of type A, the self-assembly can either be dominated by microphase separation when the homopolymers are short and well confined within block copolymer domains, or by macrophase separation in the case of sufficiently long Received: June 4, 2015 Revised: July 14, 2015

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RESULTS AND DISCUSSION The neat miktoarm star copolymer with f PS = 0.37 shows hexagonally packed PS cylinders in a PI matrix (PS HEX).13 In comparison, linear PS-b-PI diblock or symmetric PS-b-PI-b-PS triblock copolymers at the same composition show lamellar structures (LAM).28 This significant phase boundary shift was induced by the specific molecular architecture. The asymmetric junction connecting PS1 to three PI blocks prefers a concave curvature toward PS domains. In addition, as predicted by polymer brush theory, bidisperse PS brushes afford more mixing and conformational entropy to relax chain stretching, reducing the entropy penalty of concave PS domains.29,30 Thus, the three short PS2 blocks are helpful in stabilizing the discrete PS HEX phase toward high fPS. The observed phase behavior of the pure miktoarm copolymer is consistent with the predictions of self-consistent field theory (SCFT) for the A1−(B−A2)3 architecture.12 The binary blends of miktoarm star block copolymer and low molecular weight hPS (5.0 and 9.9 kg/mol respectively, see Figure 2 and Supporting Information, part B) show a regular order−order transition (OOT) from PS HEX to asymmetric LAM, and finally to disordered micelles as the hPS fraction is increased. Because of the small molecular weight of the hPS, the homopolymer can penetrate the PS brushes comprising the block copolymer mesophase to maximize translational entropy, producing a swelling of the PS domains and driving a flattening, and ultimately an inversion, of the interfacial curvature. This condition is known as “wet brush” behavior in characterizing the response of the PS1 and PS2 blocks to hPS molecules. Such HEX-LAM OOTs and LAM to disordered micelles phase transitions are consistent with the phase diagram of a conventional AB diblock/A homopolymer blend at low homopolymer molecular weight.19 When the molecular weight of the hPS is very large, comparable to or larger than that of the PS1 block in the miktoarm copolymer, such as the 79.5 kg/mol case shown in Figure 2, macrophase separation occurs over a broad composition range, which is deleterious for material properties. In this case, the long hPS molecules cannot penetrate the copolymer PS brushes (known as “dry brush” conditions) and the brush layers acquire an effective “autophobic” attraction, driving out the intervening hPS into its own macroscopic phase.20 A unique feature of the present miktoarm and hPS blends is the appearance of a new “bricks-and-mortar” (B&M) cellular mesophase when the hPS molecular weight falls between those of the PS1 and PS2 blocks. The morphologies are shown in Figure 2a for blends of the miktoarm star block copolymer (f PS = 0.37) and hPS (44.6 kg/mol) at different hPS weight fractions. In this case, the hPS molecules do not swell the short PS2 brushes but can partly penetrate into the long PS1 brushes. The gradient distribution of the long and short PS blocks at the interface evidently softens the interfaces to bending. Meanwhile, the partial penetration of the hPS into PS1 blocks gives more entropic freedom (translational and conformational) than the dry brush case, leading to a reduced tendency for macrophase separation. Most remarkably, the hPS molecules swell the PS domains significantly while retaining their discrete nature. This majority of hard PS domains (the “bricks”) intercalated in the minority of continuous rubbery PI matrix (the “mortar”) makes the B&M structure unique among equilibrium copolymer blend morphologies. The equilibrium

homopolymers, where a BCP-rich microphase coexists with a homopolymer-rich disordered phase.2,3,19,20 Dispersity of diblock and homopolymer components can assist swelling discrete domains while suppressing macrophase separation. Nonetheless, the amount of A homopolymer that can be accommodated is rather limited.21 Another effective strategy is to adopt asymmetric star block architectures, such as ABn (n > 1).22−26 Because of the overabundance of B blocks near junction points localized at A−B interfaces, the interfaces are driven to curve toward the A domains with slight penalty of conformational entropy, in order to provide a convex surface on the B side that can relieve the more significant crowding of the nB blocks per chain. This shift of the A−B interfacial curvature manifests a shift in order−order boundaries, stabilizing discrete A domains to higher fA than in linear AB block copolymers.24 However, such ABn molecules cannot provide recoverable elasticity since both ends of the soft B blocks are not anchored in hard A blocks. In this study, we present a blending strategy combining advantages of a nonlinear architecture and a bidisperse polymer brush effect to create a first-of-its-kind equilibrium mesophase with an exceptionally high volume fraction of glassy domains dispersed in a soft rubbery matrix, the composition exhibiting both high modulus and elastic recovery. Our experimental realization is a PS1 -(PI-b-PS 2 ) 3 and hPS (polystyrene homopolymer) binary blend system. The miktoarm block architecture PS1-(PI-b-PS2)3 has one long polystyrene block (PS1, number-average molecular weight, MW ∼ 80 kg/mol) and three PI-b-PS2 diblock branches, where each short PS2 block has a MW of ∼10 kg/mol. The length of the mediating PI block is varied to adjust fPS. In this study, we focus on a particular miktoarm star copolymer with f PS = 0.37. A series of hPS with varying MW (ranging from 5.0 to 79.5 kg/mol) were used, spanning the MW range of the bidisperse PS1 and PS2 blocks. The details of the samples and preparation methods can be found in the Supporting Information (part A).12,13,27 A schematic illustration of the molecular structure is given in Figure 1, parts a and b. A new stable “bricks-and-mortar”

Figure 1. (a) Schematic illustrations of the miktoarm architecture PS1(PI-b-PS2)3; (b) cartoon of bricks and mortar and the inset showing the molecular arrangement at the interface; (c) TEM image of the featured “bricks-and-mortar” structure, where the scale bar corresponds to 100 nm.

(B&M) mesophase (Figure 1c) was discovered in these blends, unique in that it lacks long-range positional order, but possesses quasi-long-range orientational order. The high achievable fractions of hard polystyrenic “bricks” are well dispersed within a rubbery “mortar” matrix, offering unique mechanical property combinations with high modulus, yet high recoverable elasticity. The boundary of the B&M structure observed here extends to 82 wt % of discrete PS domains in an 18 wt % PI matrix, which overlaps with traditional toughened PS-based thermoplastics (such as high impact polystyrene, HIPS), yet the B&M composition remains remarkably elastic rather than plastic at high strains. B

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Figure 3. Phase diagram (a) and domain spacing (b) for the miktoarm star copolymer (f PS = 0.37) and hPS blends. The orange curve corresponds to the domain spacing of complete unbinding of lamellar structures at all compositions, which separates the diagram into an upper B&M region and lower OOT region. The macrophase separation region is not depicted since the domain size is generally much larger (>1 μm).

confirmed that the orientational order is not correlated with the microtome procedure or specific orientation used to produce the TEM images, but is intrinsic to the material (and verified by SAXS; see Figure 4). The lack of positional order can be further supported by SAXS data (Figure 2b). The neat miktoarm stars show sharp peaks with featured position ratios (1:31/2:41/2:71/2:121/2), confirming ordered hexagonally packed cylinders in space. The blend with short hPS (9.9 kg/mol, 50 wt %) exhibits LAM structure with up to fifth ordered peak in the SAXS curve. In contrast, the SAXS profile does not represent any featured peaks for the blend with typical B&M structure (blend with hPS, 44.6 kg/mol, 50 wt %). The corresponding SAXS curves (Supporting Information, part D) for the B&M structures generally do not exhibit higher scattering peaks, evidently confirming the lack of positional order. As shown in Figure 3b, the domain spacing was calculated accordingly (D = 2π/q*, where q* is the primary peak position). The domain spacing increases significantly with increasing amount of hPS, and also shows strong dependence on the molecular weight of hPS.27 Compared with traditional microphase self-assembly on the ∼10 nm scale and macrophase segregation over microns, the B&M mesophase features have a wide size distribution, from several tens (Figure 3b) to several hundred nanometers (TEM images in Figure 2). The ubiquity of the B&M phase was demonstrated by reproducing it in another set of miktoarm and hPS blends, where fPS = 0.48 in the miktoarm copolymer (Supporting Information, part E). The B&M mesophase reflects a delicate balance between tendencies for microphase separation and

Figure 2. (a) TEM images for the binary blends of miktoarm copolymer (f PS = 0.37) and hPS with varying hPS weight fractions and hPS molecular weights. The scale bars are 200 nm in all images. The full TEM diagram for all compositions can be found in the Supporting Information (part B). (b) SAXS profiles of neat block copolymer and blends with 50 wt % hPS. The neat miktoarm block copolymer shows a HEX structure; the blend with 50% hPS (9.9 kg/mol) exhibits a LAM structure; the blend with 50 wt % hPS (44.6 kg/mol) lacks any positional order. The full SAXS profiles for all compositions can be found in the Supporting Information (part D).

nature of the B&M structures was confirmed by thermal annealing for 2 days and 10 days at 150 °C (Supporting Information, part C). The SAXS curves for 2-day and 10-day annealing nearly overlap by appropriate shifting vertically along the intensity axis. Consequently, we believe that the new B&M structures are not kinetically trapped metastable states. The above results are summarized into a phase diagram given in Figure 3a. From TEM, we conclude that the B&M nanostructure is fairly uniform on a large scale, but has significant local variations in size, shape, and connectivity of domains. Evidently, there is no long-range positional order of the structure, much like a polymeric microemulsion, although some images and regions suggest quasi-long-range orientational order. We have C

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orders are lost. Moreover, the microemulsion phase is bicontinuous in contrast to the monocontinuous character and extremely asymmetric composition of the B&M phase reported here. We believe that the B&M phase is an unusual type of fluctuation-stabilized mesophase,31−33 differing in its statistical and topological properties from other structured yet disordered polymer phases including micellar phases and microemulsions.34 Additional support is provided by selfconsistent field theory (SCFT) simulations (Supporting Information, part F). Since SCFT neglects thermal fluctuations, it predicts only conventional OOTs and incorrect macroscopic instability boundary, which covers the actual B&M, macrophase separation and micellar phases in experiments. More advanced computational techniques, such as field-theoretic simulations,35 are necessary to capture thermal fluctuation effects directly, although the task is especially challenging for the B&M phase due to the broad range of length scales that characterize the phase. Among a large number of possible applications of the B&M mesophase, here we discuss its potential (in the specific PS−PI system) as a remarkably stiff yet strong thermoplastic elastomer. A unique feature of the PS1−(PI−PS2)3 architecture is that the junctions between PI blocks and the short PS2 tails supply multiple anchoring points in the glassy PS domains, compared with fewer anchoring points in linear-chain BCP counterparts.13 From the TEM images, the B&M mesophase structure is observed at up to 50 wt % hPS (70 wt % PS in total). Discrete PS-rich domains are stabilized to even higher 70 wt % hPS (82 wt % PS in total), although domain sizes are polydisperse and approaching the macrophase length scale ∼500 nm (Supporting Information, part G). Since the rubbery PI phase apparently still comprises the continuous matrix, we would expect that nearly reversible elasticity is well maintained up to high strain, while the high PS content contributes stiffness and strength. This expectation was validated through mechanical tensile testing. Monotonic tensile tests indicated that the Young’s modulus of miktoarm block copolymer alloys increases with increasing hPS fraction (Supporting Information, part H). Moreover, the step-cycle tests indicated that high elastic recovery above 90% was obtained in blends with hPS content up to 70 wt % (82 wt % PS in total). The modulus of the limiting composition is 99.2 (±4.0) MPa, 2 orders of magnitude larger than traditional linear polystyrene-based thermoplastic elastomers (∼1 MPa, f PS less than 30 wt %).13 The remarkable elastic recovery above 90% up to break is direct evidence of the continuous nature of the rubbery phase. Notably, the stiffness and recoverable elasticity of this B&M structure is without precedent in classic TPEs (such as polyurethanes or PS−PI−PS) or polymer alloys that have neither been cross-linked nor loaded with inorganic fillers. One striking feature is that the strong TPEs exhibit sharp yielding phenomena in the envelope of the step cycle and monotonic tensile tests (Figure 4a, at 0.2 strain and 5.5 MPa stress). In glassy and semicrystalline polymers, yielding is usually a sign of the onset of significant plastic deformation that cannot be fully recovered when the load is removed.36 Sharp yielding point marks the plastic instability under stretching. However, for these strong and hard TPEs, the recovery is still higher than 90% even after the sharp yielding point. We therefore believe the yielding behavior cannot be fully attributed to conventional plastic deformation, and some other type of recoverable deformation need to be investigated.

Figure 4. Mechanical testing results for PS1-(PI−PS2)3 (f PS = 0.37) and hPS (50 wt %, 44.6 kg/mol) blend: (a) Monotonic and step-cycle tensile tests; (b) 2-D SAXS patterns for monotonic tests at different strains with the arrow indicating the tensile direction; (c) void fraction variation during the step-cycle test with strain increasing from cycle to cycle.

macrophase separation created by the unusual mixed wet/dry brush conditions and strong conformational asymmetry of miktoarm star copolymers with intermediate MW hPS. The long-range positional order in the B&M state was destroyed by thermal fluctuations into large (but not macroscopic) discrete domains of nonuniform shape, but with orientational bias. The orientational order seems to arise from a layering of highly defective asymmetric lamellae in which the glassy domains are interrupted by numerous random perforations by the PI elastomer. Nonetheless, the B&M phase is not preceded by a LAM phase at lower hPS content, so it is difficult to view it as emerging continuously from LAM by a proliferation of defects. The fluctuation-driven nature of the B&M phase has striking similarity with the previously reported and well-understood bicontinuous polymeric microemulsion phase,31−34 which also appears in a narrow channel between microphase separation and macrophase separation. In the microemulsion case, the fluctuation stabilized phase emerges by thermal unbinding of highly swollen lamellae, but both translational and orientational D

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continuous sheet of rubbery matrix surrounding the cellular hard PS domains. In contrast, voids usually cannot fully recover in conventional (including rubber-modified) thermoplastics because of the hydrostatic stress sustained by the deformed continuous glassy phase.38−40 We infer that the small unrecoverable part (less than 10%) of the deformation observed for the B&M phase is associated with minor plastic deformation during each cycle. The high recoverable elasticity of TPEs relies on the existence of a continuous rubbery phase, while discrete glassy domains contribute a reinforcement of the modulus. Combining the TEM, SAXS, and mechanical data, we thus have compelling evidence that the PS glassy domains of the B&M phase remain discrete up to 82 wt % PS in total. Notably, this composition range strongly overlaps the regime of traditional rubber-toughened thermoplastics, where the low fractions of rubbery component comprise the discrete domains in glassy matrix, such as high impact PS (HIPS). It is interesting to compare the B&M morphology with the well-known “salami” morphology in HIPS. HIPS is a complex, reactively formed blend of PS homopolymer, a graft copolymer of butadiene and styrene (PBgS) and some residual PB homopolymer.36,38 The content of PB rubber is usually less than 20 wt % and it is almost entirely contained within micronsized inclusions rich in the graft copolymer that have macrophase separated from the continuous PS matrix. Cross sections of the rubber-rich inclusions have a “salami” microphase appearance with locally continuous PB domains surrounding discrete PS domains. Unlike the B&M structure, however, the rubbery PB domains in HIPS are only continuous within the salami inclusions, not throughout the entire composition. Thus, HIPS is plastic rather than elastic and has no significant recoverable deformation beyond the yielding point. Cavitation of the PB domains produces voids that act as localized centers to relieve internal stress or merge to form crazes, which significantly improves the impact strength of HIPS.41 In contrast, the B&M mesophase is uniformly cellular in its morphology throughout the entire material. The continuous rubber phase affords the recoverable elasticity of a TPE, while the unusually high content of discrete PS confers significant modulus.

Accompanied by yielding, the specimens showed whitening phenomena, but did not exhibit “necking”. To gain insight into the deformation behavior, we carried out tensile-SAXS experiment to probe the mechanism. The initial anisotropic scattering pattern reflected the quasi-long-range orientation order in the sample before loading. Significant intensity growth was observed as loading was increased to the vicinity of the yielding point. The 2-D scattering pattern under strain showed additional anisotropic features (Figure 4b), including especially strong intensity in the vertical stretching direction and weaker amplification in the horizontal direction (Supporting Information: part I). These features are consistent with void formation under tensile stretching; similar behavior is observed with traditional elastomers (such as filled natural rubbers), thermoplastics and thermosets.37−42 Using a well-established three-phase model,37 we can quantitatively calculate the volume fraction of voids during tensile stretching (Figure 5). The neat miktoarm block

Figure 5. Volume fraction of voids during monotonic tensile stretching for miktoarm-hPS blends at different hPS fractions. The neat miktoarm block copolymer did not show void formation during stretching.

copolymer did not show void formation until break. The blend with 30 wt % hPS showed void formation when the strain was above 0.5, whereas the blend with 50 wt % hPS showed significant void development when the strain was above 0.13. Moreover, voids formed nearly instantaneously with no perceptible threshold upon loading in the blend with 70 wt % hPS. Compared with the stress−strain curves (Figure 4a and Supporting Information: part G), the onset of the yielding (deviation from linear mechanical behavior) was slightly ahead of voiding for the blends with 30 and 50 wt % hPS, while voiding and onset of yielding are coincident in the material with 70 wt % hPS content. We infer that voiding plays an important role in the yielding behavior, and the volume (or weight) fraction of the glassy phase dictates the onset of void formation. Significant void development is responsible for the appearance of a sharp yielding point and results in the plastic instability under stretching. To gain further insight into the excellent elastic properties, the scattering variation was monitored during the step-cycle tensile tests (Figure 4c). We observed that voids do not form before the yield point and the void fraction increases with strain in the subsequent cycles. One remarkable feature is the full recovery/healing of voids when the load was removed in each cycle. This suggests that the voids are concentrated in the thin

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CONCLUSION In conclusion, a carefully designed PS1-(PI-b-PS2)3 miktoarm star block copolymer was observed to form a remarkable equilibrium, nanocellular emulsion when blended with hPS of intermediate molecular weight, combining the advantages of a nonlinear architecture and a bidisperse polymer brush effect. This new nanostructured emulsion, the “bricks-and-mortar” mesophase, was identified in compositions that delicately balance tendencies for microphase and macrophase separation and is believed to be stabilized by thermal fluctuations. The B&M phase lacks long-range positional order but exhibits quasi-long-range orientational order. In the specific PS/PI realization explored here, the unusually high content of discrete glassy PS domains imparts high stiffness and the low content of continuous rubbery PI matrix sustains recoverable elasticity, which leads to mechanical behavior consistent with a hardtough-strong thermoplastic elastomer. Beyond the specific context, our work reveals a powerful protocol for producing equilibrium, melt processable, meso-cellular phases with ultrahigh volume fractions of discrete A domains, embedded in a continuous B matrix. Such compositions could benefit E

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applications where the B component must be maintained as continuous, e.g., for chemical resistance or electrical or ionic conductivity, while it is desirable to have a majority of a second component A, either for reasons of cost or to balance other properties.

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The authors declare no competing financial interest. ○ Edward J. Kramer passed away on Dec. 27th, 2014.

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ACKNOWLEDGMENTS This research was supported by the Institute for Collaborative Biotechnologies through Grant W911NF-09-0001 from the U.S. Army Research Office. The content of the information does not necessarily reflect the position or the policy of the Government, and no official endorsement should be inferred. Extensive use was made of the MRL Shared Experimental Facilities supported by the MRSEC Program of the NSF under Award No. DMR 1121053; a member of the NSF-funded Materials Research Facilities Network. This work was performed at the DuPont-Northwestern-Dow Collaborative Access Team (DND-CAT) located at Sector 5 of the Advanced Photon Source (APS). DND-CAT is supported by E.I. DuPont de Nemours & Co., The Dow Chemical Company and Northwestern University. Use of the APS, an Office of Science User Facility operated for the U.S. Department of Energy (DOE) Office of Science by Argonne National Laboratory, was supported by the U.S. DOE under Contract No. DE-AC0206CH11357.

EXPERIMENTAL SECTION

Sample Preparation. The miktoarm star block copolymer and PS homopolymer were mixed with different ratios in toluene. The total weight fraction of polymers was 2% in the solution. Also, 0.5 wt % of antioxidant BHT, with respect to the total polymers, was added to protect the PI blocks from oxidation. The mixture was stirred at room temperature in a fume hood and toluene was allowed to evaporate slowly. The dried samples were then annealed at 150 °C for 48 h (or 10 days) in a high vacuum chamber (10−8 mbar). The prepared samples were then ready for further characterization. Transmission Electron Microscopy (TEM). Ultrathin (∼100 nm) sections were cut by a cryo-ultramicrotome at −90 °C. The thin slices were collected by copper grids (CF-300 from Electron Microscopy Sciences) and then stained in osmium tetroxide vapors for 2−4 h to enhance the contrast for TEM. Mechanical Tensile Testing. The specimens were in dog-bone shape by 7 mm gauge length and a cross-section of 2 mm wide and ∼0.5 mm thick. Monotonic and step-cycle tensile tests were carried out with five specimens on a house-made tensile tester. The load cell had capacity of 40 N. The bottom clamp was stationary while the upper one can be stretched. The crosshead speed of was kept at 5 mm/min for all the specimens, which produced an initial strain rate of 0.012 s−1. Synchrotron Small-Angle X-ray Scattering (SAXS). A part of the SAXS experiments were carried out using the Advanced Light Source beamline 7.3.3 at Lawrence Berkeley National Laboratory. The X-ray wavelength of the beam is 0.124 nm. Real-Time Tensile−SAXS Experiment. The synchrotron smallangle X-ray scattering (SAXS) experiments were carried out using the Advanced Photon Source beamline 5-ID at Argonne National Laboratory. The wavelength of the source was 0.124 nm. The sample to detector distance was set to 7496 mm so that the lowest wave vector was 0.0155 nm−1, which is equivalent to 400 nm. The beamline was adapted to an Instron 850 tensile tester. The load cell capacity was 200 N. Both upper and bottom clamps were movable to create symmetric tensile stretching. The strain rate was kept at 0.012 s−1 for all tests.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.5b01210. Part A: Samples used and molecular characterization. Part B: Full TEM diagram for miktoarm copolymer ( f PS = 0.37) and hPS. Part C: 10-day annealing experiment. Part D: SAXS curves. Part E: Bricks-and-mortar structure in miktoarm (f PS = 0.48) and hPS blends. Part F: Phase diagram of miktoarm and hPS by SCFT. Part G: Monotonic and step-cycle tensile testing results. Part H: Monotonic and modulus at different hPS fractions. Part I: Scattering curves for the tensile-SAXS experiment at different strains (PDF)

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REFERENCES

(1) (a) Bates, F. S.; Fredrickson, G. H. Annu. Rev. Phys. Chem. 1990, 41, 525. (b) Bates, F. S.; Fredrickson, G. H. Phys. Today 1999, 52, 32. (c) Bates, F. S.; Hillmyer, M. A.; Lodge, T. P.; Bates, C. M.; Delaney, K. T.; Fredrickson, G. H. Science 2012, 336, 434. (2) (a) Winey, K. I.; Thomas, E. L.; Fetters, L. J. Macromolecules 1991, 24, 6182. (b) Winey, K. I.; Thomas, E. L.; Fetters, L. J. Macromolecules 1992, 25, 2645. (c) Winey, K. I.; Gobran, D. A.; Xu, Z.; Fetters, L. J.; Thomas, E. L. Macromolecules 1994, 27, 2392. (3) (a) Court, F.; Yamaguchi, D.; Hashimoto, T. Macromolecules 2006, 39, 2596. (b) Chen, F.; Kondo, Y.; Hashimoto, T. Macromolecules 2007, 40, 3714. (c) Yamaguchi, D.; Takenaka, M.; Hasegawa, H.; Hashimoto, T. Macromolecules 2001, 34, 1707. (d) Morita, H.; Kawakatsu, T.; Doi, M.; Yamaguchi, D.; Takenaka, M.; Hashimoto, T. Macromolecules 2002, 35, 7473. (4) Silverstein, M. S. Prog. Polym. Sci. 2014, 39, 199. (5) (a) Mezzenga, R.; Ruokolainen, J.; Fredrickson, G. H.; Kramer, E. J.; Moses, D.; Heeger, A. J.; Ikkala, O. Science 2003, 299, 1872. (b) Mezzenga, R.; Fredrickson, G. H.; Kramer, E. J. Macromolecules 2003, 36, 4457. (c) Mezzenga, R.; Ruokolainen, J.; Fredrickson, G. H.; Kramer, E. J. Macromolecules 2003, 36, 4466. (6) Park, M.; Harrison, C.; Chaikin, P. M.; Register, R. A.; Adamson, D. H. Science 1997, 276, 1401. (7) Hawker, C. J.; Russell, T. P. MRS Bull. 2005, 30, 952. (8) Kim, J. K.; Yang, S. Y.; Lee, Y.; Kim, Y. Prog. Polym. Sci. 2010, 35, 1325. (9) Sai, H.; Tan, K. W.; Hur, K.; Asenath-Smith, E.; Hovden, R.; Jiang, Y.; Riccio, M.; Muller, D. A.; Elser, V.; Estroff, L. A.; Gruner, S. M.; Wiesner, U. Science 2013, 341, 530. (10) Yao, H.-B.; Fang, H.-Y.; Wang, X.-H.; Yu, S.-H. Chem. Soc. Rev. 2011, 40, 3764. (11) Holten-Andersen, N.; Fantner, G. E.; Hohlbauch, S.; Waite, J. H.; Zok, F. W. Nat. Mater. 2007, 6, 669. (12) Lynd, N. A.; Oyerokun, F. T.; O’Donoghue, D. L., Jr.; Handlin, D. L.; Fredrickson, G. H. Macromolecules 2010, 43, 3479. (13) Shi, W.; Lynd, N. A.; Montarnal, D.; Luo, Y.; Fredrickson, G. H.; Kramer, E. J.; Ntaras, C.; Avgeropoulos, A.; Hexemer, A. Macromolecules 2014, 47, 2037. (14) Patel, R. M.; Hahn, S. F.; Esneault, C.; Bensason, S. Adv. Mater. 2000, 12, 1813. (15) Cheung, K. C.; Gershenfeld, N. Science 2013, 341, 1219. (16) Ruzette, A.-V.; Leibler, L. Nat. Mater. 2005, 4, 19.

AUTHOR INFORMATION

Corresponding Authors

*(G.H.F.) E-mail: [email protected]. *(A.A.) E-mail: [email protected]. F

DOI: 10.1021/acs.macromol.5b01210 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules (17) Fulvio, P. F.; Mayes, R. T.; Wang, X.; Mahurin, S. M.; Bauer, J. C.; Presser, V.; McDonough, J.; Gogotsi, Y.; Dai, S. Adv. Funct. Mater. 2011, 21, 2208. (18) Begley, M. R.; Philips, N. R.; Compton, B. G.; Wilbrink, D. V.; Ritchie, R. O.; Utz, M. J. Mech. Phys. Solids 2012, 60, 1545. (19) Matsen, M. W. Phys. Rev. Lett. 1995, 74, 4225. (b) Matsen, M. W. Macromolecules 1995, 28, 5765. (20) Leibler, L. Makromol. Chem., Macromol. Symp. 1988, 16, 1. (21) (a) Matsen, M. W. Phys. Rev. Lett. 2007, 99, 148304. (b) Thompson, R. B.; Matsen, M. W. Phys. Rev. Lett. 2000, 85, 670. (22) Dyer, C.; Driva, P.; Sides, S. W.; Sumpter, B. G.; Mays, J. W.; Chen, J.; Kumar, R.; Goswami, M.; Dadmun, M. D. Macromolecules 2013, 46, 2023. (23) Beyer, F. L.; Gido, S. P.; Velis, G.; Hadjichristidis, N.; Tan, N. B. Macromolecules 1999, 32, 6604. (24) Milner, S. T. Macromolecules 1994, 27, 2333. (25) (a) Matsen, M. W.; Schick, M. Macromolecules 1994, 27, 6761. (b) Matsen, M. W. Macromolecules 2012, 45, 2161. (26) Avgeropoulos, A.; Hadjichristidis, N. J. Polym. Sci., Part A: Polym. Chem. 1997, 35, 813. (27) Shi, W.; Hamilton, A. L.; Delaney, K. T.; Fredrickson, G. H.; Kramer, E. J.; Ntaras, C.; Avgeropoulos, A.; Lynd, N. A. J. Am. Chem. Soc. 2015, 137, 6160. (28) Khandpur, A. K.; Förster, S.; Bates, F. S.; Hamley, I. W.; Ryan, A. J.; Bras, W.; Almdal, K.; Mortensen, K. Macromolecules 1995, 28, 8796. (29) (a) Milner, S. T.; Witten, T. A.; Cates, M. E. Macromolecules 1988, 21, 2610. (b) Milner, S. T.; Witten, T. A.; Cates, M. E. Macromolecules 1989, 22, 853. (30) Dan, N.; Tirrell, M. Macromolecules 1993, 26, 6467. (31) Fredrickson, G. H.; Helfand, E. J. Chem. Phys. 1987, 87, 697. (32) Bates, F. S.; Rosedale, J. H.; Fredrickson, G. H. J. Chem. Phys. 1990, 92, 6255. (33) Klushin, L. I.; Skvortsov, A. M. Macromolecules 1992, 25, 3443. (34) (a) Bates, F. S.; Maurer, W.; Lodge, T. P.; Schulz, M. F.; Matsen, M. W.; Almdal, K.; Mortensen, K. Phys. Rev. Lett. 1995, 75, 4429. (b) Bates, F. S.; Maurer, W. W.; Lipic, P. M.; Hillmyer, M. A.; Almdal, K.; Mortensen, K.; Fredrickson, G. H.; Lodge, T. P. Phys. Rev. Lett. 1997, 79, 849. (c) Hillmyer, M. A.; Maurer, W. W.; Lodge, T. P.; Bates, F. S. J. Phys. Chem. B 1999, 103, 4814. (d) Fredrickson, G. H.; Bates, F. S. J. Polym. Sci., Part B: Polym. Phys. 1997, 35, 2775. (35) Fredrickson, G. H. The Equilibrium Theory of Inhomogeneous Polymers; Oxford: New York, 2006. (36) Strobl, G. The Physics of Polymers, 2nd ed., Springer: New York, 2010. (37) (a) Zhang, H.; Scholz, A. K.; de Crevoisier, J.; Vion-Loisel, F.; Besnard, G.; Hexemer, A.; Brown, H. R.; Kramer, E. J.; Creton, C. Macromolecules 2012, 45, 1529. (b) Zhang, H.; Scholz, A. K.; VionLoisel, F.; Merckel, Y.; Brieu, M.; Brown, H. R.; Roux, S.; Kramer, E. J.; Creton, C. Macromolecules 2013, 46, 900. (38) Fischer, M.; Hellmann, G. P. Macromolecules 1996, 29, 2498. (39) Kuboki, T.; Ben Jar, P.-Y.; Takahashi, K.; Shinmura, T. Macromolecules 2002, 35, 3584. (40) Salomons, G. J.; Singh, M. A.; Bardouille, T.; Foran, W. A.; Capel, M. S. Macromolecules 1999, 32, 1264. (41) Bubeck, R. A.; Buckley, D. J., Jr.; Kramer, E. J.; Brown, H. J. J. Mater. Sci. 1991, 26, 6249. (42) Declet-Perez, C.; Francis, L. F.; Bates, F. S. ACS Macro Lett. 2013, 2, 939.

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DOI: 10.1021/acs.macromol.5b01210 Macromolecules XXXX, XXX, XXX−XXX