Article pubs.acs.org/Macromolecules
Morphological Consequences of Frustration in ABC Triblock Polymers Madalyn R. Radlauer,† Christophe Sinturel,§ Yusuke Asai,∥ Akash Arora,‡ Frank S. Bates,‡ Kevin D. Dorfman,‡ and Marc A. Hillmyer*,† †
Department of Chemistry and ‡Department of Chemical Engineering and Materials Science, University of Minnesota, Minneapolis, Minnesota 55455, United States § ICMN, UMR 7374 CNRS, Université d’Orléans, 1 B rue de la Férollerie, 45071 Orléans, Cedex 2, France ∥ Department of Applied Chemistry, Graduate School of Engineering, Nagoya University, Furo-cho, Chikusa-ku, Nagoya 464-8603, Japan S Supporting Information *
ABSTRACT: Three poly(styrene)-block-poly(isoprene)-block-poly(lactide) (PS-b-PI-b-PLA, SIL) triblock terpolymers were synthesized and characterized in the bulk and as thin films. The pronounced incompatibility of the covalently connected PI and PLA led to significant frustration and the tendency to minimize their intermaterial dividing surface area. This resulted in the formation of a core−shell cylinder morphology with exaggerated nonconstant mean curvature from triblock polymers with equal block volume fractions rather than the more typical lamellar morphology. The effect of frustration was magnified in thin films by both confinement and interfacial interactions such that the PI domains became discontinuous. Selfconsistent field theory (SCFT) calculations emphasize that the marked difference in the PS/PI and PI/PLA interaction parameters promotes the formation of nonlamellar morphologies. However, SCFT predicts that lamellar morphology is more stable than the observed cylindrical morphology, demonstrating a limitation that arises from the underlying assumptions.
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morphologies analogous to the diblocks,20 core−shell versions,21 and superstructures such as the knitting pattern22 or cylinders-at-lamellae.23,24 The superstructure morphologies found in ABC triblocks often adopt A/C interfaces, even though these two blocks are not covalently connected.23 The favorability of creating such interfaces can stem from high incompatibility of the covalently connected blocks relative to the outer blocks.25,26 In the case where χACN ≪ χABN or χBCN, the system is considered to be frustrated, and the likelihood of A/C contacts is increased.16,31 This is due to the forced covalent connection of the block pairs with the highest relative incompatibility. The morphologies of several pairs of frustrated and nonfrustrated polymers have been compared including poly(styrene)-block-poly(isoprene)block-poly(2-vinylpyridine) (PS-b-PI-b-P2VP, SIP)27,28 and PIb-PS-b-P2VP (ISP),20 poly(styrene)-block-poly(butadiene)block-poly(2-vinylpyridine) (PS-b-PB-b-P2VP, SBP)29 and PBb-PS-b-P2VP (BSP),29 and poly(styrene)-block-poly(isoprene)block-poly(ethylene oxide) (PS-b-PI-b-PEO, SIO)25 and PI-bPS-b-PEO (ISO).30 On the whole, the frustrated systems (SIP, SBP, and SIO) have much more complicated morphology maps than the corresponding nonfrustrated systems, and only with
INTRODUCTION Block polymers are fascinating self-assembling materials.1−3 Over the past two decades, great expansion of the potential applications of these materials both in the bulk and as thin films has begun to revolutionize thinking around, for example, membrane separations,4−7 nanopatterning,8−11 and microelectronics.12−14 As the field expands, a comprehensive understanding of the factors impacting the morphology of these polymer systems becomes more relevant to guiding future research and associated technologies. For an AB diblock polymer, the principal predictive parameters for microphase separation are the volume fraction of A (fA; f B = 1 − fA), the degree of polymerization (N) based on a common segment reference volume, and the Flory− Huggins interaction parameter (χAB) that quantifies the incompatibility of the two blocks.1,15 These three parameters can be varied to access the principal equilibrium structures for diblock polymers: lamellae, gyroid, cylinders, and spheres. The conformational asymmetry between the two blocks, which depends on the relative segment volumes and segment lengths, can also influence morphological behavior.16−19 For ABC triblock polymers, the relevant parameter set is much larger: fA, f B, (f C = 1 − fA − f B), N, χAB, χBC, χAC, and two conformational asymmetry parameters (between each pair of adjacent blocks).15 Study of this large parameter space has led to the discovery of over 20 stable structures in the bulk, including the © XXXX American Chemical Society
Received: September 27, 2016 Revised: December 2, 2016
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DOI: 10.1021/acs.macromol.6b02112 Macromolecules XXXX, XXX, XXX−XXX
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core−shell cylinders with exaggerated nonconstant mean curvature (non-CMC) between the PS and PI at equal volume fractions of the three blocks, a composition where 3-fold lamellae are generally expected.27 This likely stems from the relatively large incompatibility between the covalently connected PI and P2VP blocks that drives the formation of curved interfaces with smaller interfacial area as compared to the corresponding layered phases. If SIL polymers could achieve hexagonally packed core−shell cylinders over a wide range of compositions, such a system would allow for tunable pore sizes and extents of functionalization by changing the Mn of PI and PLA in the triblock polymer. In this work we targeted two SIL triblock polymers with equal volume fractions of the three blocks and an additional SIL triblock polymer with a lower volume fraction of PLA. After synthesis and molecular characterization, the bulk morphologies of the polymers were determined using TEM and SAXS experiments. SCFT calculations were used to investigate the stability of the observed morphologies. The polymers were then spin-coated onto native silicon oxide, HMDS-modified silicon (HMDS = hexamethyldisilazane), and carbon-coated mica and studied with AFM. The films were also lifted from the carboncoated mica by floating in water54 and characterized by TEM experiments after selective staining of the PI domains with OsO 4 . In both the bulk and thin film, nonlamellar morphologies were observed.
the frustrated polymers are superstructure morphologies with A/C interfaces observed. Theoretical calculations and simulations of ABC triblock polymers in the bulk have been used to determine the thermodynamically stable morphologies and have predicted or corroborated experimentally observed morphologies.15,16,31−35 In particular, self-consistent field theory (SCFT) is the most widely used theoretical framework to study the phase behavior of block polymers.36,37 SCFT has been remarkably successful in aiding understanding as well as discovering new, complex morphologies observed experimentally in block polymers.38−41 SCFT is a mean field theory that is strictly valid in the limit of infinitely long chains but can predict the relative stability of competing ordered phases with reasonable accuracy.36 Understanding and predicting morphologies of triblock polymers becomes even more complicated when dealing with thin films, as confinement and surface interactions come into play.42−44 Both the absolute thickness of the films and their relative value with respect to the domain spacing of the ordered material can greatly impact the structures adopted.45 Matching of surface energy of the substrate or free surface with one or more of the blocks can also change microstructure orientation.42,43 Furthermore, the obtained morphology can be highly dependent on the processing and annealing conditions.46,47 The thin film morphologies of ABC triblock polymers observed using atomic force or scanning force microscopy (AFM or SFM) can be the thin film analogues of their bulk structures or, due to these additional considerations, achieve structures unattainable in the bulk.6,46,48−50 These complex thin film morphologies can be applied as templates or masks to allow for more interesting and complex patterns for applications in nanopatterning.49 Combining the expectation of interesting morphologies due to frustration with the ability to etch two of the blocks, Guo and co-workers synthesized a poly(styrene)-block-poly(isoprene)-block-poly(lactide) (PS-b-PI-b-PLA, SIL) triblock polymer ( f S = 0.60; f I = 0.11; f L = 0.29; number-average molar mass, Mn = 65 kg mol−1; herein abbreviated as S0.60I0.11L0.29-65, where the subscripts are the volume fraction of each block and the number after the dash is the overall Mn in kg mol−1).51 The triblock polymer was spin-coated onto hydrophobically modified silicon substrates. After thermal annealing of S0.60I0.11L0.29-65, core−shell cylinders with the PLA block on the interior were reported for film thicknesses of 65, 100, and 170 nm. Chemical etching of the PLA produced ordered nanoporous films with PI on the pore walls. Functionalization of the PI on the pore walls was then explored in the corresponding bulk system by Bailey and co-workers using a series of SIL triblock polymers with a fixed ratio of f S/f I = 7:3 and Mn values between 15 and 56 kg mol−1.52 For these samples, small-angle X-ray scattering (SAXS) and transmission electron microscopy (TEM) confirmed primarily core−shell hexagonally packed morphologies, analogous to what was seen in thin films. Thin films of S0.62I0.13L0.25-65 displaying perpendicular core−shell cylinders were subsequently utilized by Kubo and co-workers for nanopatterning applications.53 The frustration of this SIL system, the facile etching of the PLA block, and the ability to functionalize or etch the PI block make SIL triblock terpolymers intriguing targets for various nanomaterial applications. Moreover, the possibility of accessing diverse morphologies in this system has not been systematically explored. In the related (and frustrated) SIP triblock polymers, Gido et al. observed hexagonally packed
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METHODS
Materials. All air- and water-sensitive materials were handled on the Schlenk line or in a nitrogen atmosphere glovebox. The materials for the anionic polymerization were thoroughly dried and degassed before the reaction according to standard procedures.25 rac-Lactide was received from Ortec and recrystallized from toluene three times before drying in vacuo and bringing into the glovebox for use. 1,8Diazabicyclo[5.4.0]undec-7-ene (DBU) was purchased from SigmaAldrich and used as received. Solvents used for solvent casting, spincoating, and solvent vapor annealing were acquired using solvent purification columns.55 Polymer Synthesis. The SI diblock polymer with a terminal hydroxyl group was synthesized by sequential anionic polymerization, as previously described.25 Briefly, polymerization of styrene was initiated by sec-butyllithium, after which isoprene was added and the diblock was terminated by end-capping with a single ethylene oxide unit followed by termination with acidified methanol (Scheme 1). Ring-opening polymerization of rac-lactide with DBU as the catalyst provided the desired SIL triblock terpolymers (Scheme 1).56 SI diblock polymer (0.18 mmol, 1 equiv) was added to a Schlenk flask equipped with a stir bar. Dry toluene was added to dissolve the polymer, and the solution was pumped to dryness on the Schlenk line overnight (to remove H2O). The flask was brought into the glovebox, and the polymer was dissolved in dry dichloromethane (DCM, 20 mL). DBU (0.19 mmol, 1.08 equiv) was added by microsyringe, and then rac-lactide (1.02 equiv relative to the desired conversion) was added as a solid using up to 2 mL of DCM to ensure full transfer. After 30 min, the reaction was removed from the glovebox and quenched by addition of benzoic acid (1.8 mmol, 10 equiv). The triblock polymer was purified via dropwise addition to stirring MeOH (600 mL) to precipitate the polymer, which was collected by simple filtration. The precipitation step was repeated after dissolving the polymer in fresh DCM, and the white solid was transferred to a tared glass jar and dried in a vacuum oven at 40 °C overnight to yield the desired material in high yield (>95%). 1H NMR (CDCl3, 500 MHz): 7.24−6.24 (5H per PS unit), 5.34−4.94 (1H per 1,4 PI unit and 1H per PLA unit), 4.85− 4.60 (2H per 3,4 PI unit), 4.17−3.97 (2H per chain, CH2O), 3.63− 3.53 (1H per chain, OH), 2.3−1.2 ppm (aliphatic protons) (Figure S1). B
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Bulk Sample Preparation. Samples for TEM and SAXS were initially prepared by solvent casting from THF ∼ 80 mg of polymer with 0.5 wt % BHT in a Teflon beaker (1.5 cm in diameter) over 1 week, resulting in a thickness of approximately 0.5 mm. This was followed by thermal annealing for 24 h at ∼120 °C in a vacuum oven at roughly 50 mTorr. After these processing techniques, SEC, SAXS, and TEM revealed evidence of degradation of the sample such that shoulders appeared on peaks and different features were observed by TEM (Figures S6−S8). Later, samples for SAXS were prepared by solvent casting from THF ∼ 80 mg of polymer with 0.5 wt % BHT in a Teflon beaker (1.5 cm in diameter) over 1 week in an inert atmosphere. This was followed by thermal annealing for 16 h at ∼120 °C in a vacuum oven at roughly 50 mTorr. These changes to processing eliminated the extra features in the SAXS patterns and were therefore used for all subsequent processing (Figure S7). After removal from the Teflon beaker, the films were stored in individual sealed plastic bags in the freezer. Small-Angle X-ray Scattering (SAXS). SAXS data were collected at the Advanced Photon Source (APS) at Argonne National Laboratory in beamline Sector 5-ID-D. The beamline is maintained by the Dow−Northwestern−Dupont Collaborative Access Team (DND-CAT). In general, the largest peak at low q was chosen as q* and the domain spacing, DSAXS = 2π/q*. For the profiles with hexagonal symmetry the cylinder-to-cylinder distance (Dc−c) was calculated as Dc−c = 2DSAXS/√3 (Figure S9). Transmission Electron Microscopy (TEM). TEM images were acquired using a FEI Tecnai G2 Spirit BioTWIN with a LaB6 gun and an accelerating voltage of 120 kV in bright-field mode. The bulk samples were embedded in an epoxy resin using an Electron Microscopy Sciences EMBed-812 Kit, microtomed, and collected on copper grids for analysis with bulk staining in OsO4 at 50 °C before embedding and vapor staining with OsO4 and I2 at 50 °C for 1 h. By the gauge on the ultramicrotome and the silver to gold color of the sections, the samples were roughly 60−100 nm thick. To determine the cylinder-to-cylinder distances from the TEM images of the bulk samples, the spacings of 4−6 domains were measured and then divided to produce a value for a single domain spacing. This measurement was repeated for all three directions of the hexagonal pattern to account for the asymmetry caused by microtoming, and the average value is reported. For TEM of the thin films, the films were scored and floated in water from the carbon-coated mica surface. Then the film sections were collected on copper grids and stained with OsO4 vapor under various conditions (ambient temperature or 50 °C and from 5 min to 1 h) to observe the best possible resolution of the different domains. Based on thin films spun with the same conditions onto silicon wafers (vide infra), the thickness of these samples was between 25 and 35 nm. Thin Film Preparation and Analysis. Silicon wafers were used as purchased (“native Si”) or modified with HMDS (“HMDS-modified Si”) as follows. The wafer was cleaned by submerging in piranha solution (3:1 H2SO4/H2O2) for 15 min, then rinsed with DI H2O, blown dry with argon, and further dried in an oven at 50 °C for 30 min. The wafer was then submerged in a solution of 10 mL of HMDS (as purchased from Sigma-Aldrich) and 50 mL of dry toluene (from solvent purification columns). After 16−18 h, the wafer was rinsed with dry toluene, and the solvent was removed by blowing with argon. If the wafer was left for a longer period of time (24 h), it was found to be too hydrophobic and the ordering of the films was decreased. Freshly cleaved mica was coated with a thin layer of carbon (“carboncoated mica”) according to literature procedures.57 Before spin-coating, the substrates were cut into 2 cm × 2 cm or 1 cm × 1 cm squares and cleaned with toluene and N2 gas. Spin-coating of the substrates was achieved by preparing 0.5 or 1 wt % solutions of the polymers in toluene. Spin-coating at ambient conditions was performed at 2000 or 3000 rpm for 60 s. Ellipsometry on a J.A. Wollam Co., Inc., V-VASE measured the film thicknesses of the thin films on the Si wafers to be between 15 and 45 nm. Most of the films were around 30 nm thick. Because of the birefringence of mica, film thickness on the mica substrate was not measured but was presumed to be similar to the films prepared at the same spinning speed (rpm) on the Si wafers. Tapping mode atomic force microscopy (AFM) was
Scheme 1. Synthesis of PS-b-PI-b-PLA (SIL) Triblock Polymers
Molecular Characterization. All 1H NMR spectra were acquired on a Bruker Avance III HD with SampleXpress equipped with a 5 mm Prodigy TCI cryoprobe with z-axis gradients. To ensure consistency of integrals between spectra, 32 scans and a 20 s delay time were used for every spectrum. The residual protio peak for CDCl3 was set to 7.26 ppm as a reference. The integrations of the 1H NMR spectra were used to determine the Mn for each of the blocks, which was then used to calculate the volume fractions. The integrations were initially determined using end-group analysis in the spectra of the diblock polymers: the integral of the CH2OH peak with a range of 3.73−3.48 ppm was set to 2. For the triblock polymers, the integral of the PS peak was set to match that of its diblock precursor. This led to good agreement with the integrals of the 3,4 PI peak and the end-group peaks, and the integral of the peak for the 1,4 PI and PLA could be measured so that the Mn of PLA could be determined. Size exclusion chromatography (SEC) analysis was performed at 25 °C on an Agilent 1260 Infinity liquid chromatograph system equipped with three Waters Styragel columns in series, a Wyatt DAWN Heleos II 18-angle light scattering detector, and a Wyatt OPTILAB T-rEX refractive index detector and using tetrahydrofuran (THF) as the eluent with a flow rate of 1 mL min−1 (Figure S2). Molar mass dispersity (Đ) was calculated from the light scattering data using Astra software. Thermal Characterization. Thermogravimetric analysis (TGA) traces were collected on a TA Instruments Q500 under house nitrogen with a heating ramp rate of 10 °C/min (Figure S3). Differential scanning calorimetry (DSC) traces were collected on a TA Instruments Discovery DSC using a temperature ramp rate of 10 °C/min (Figure S4). The samples were heated to 120 °C, cooled to −85 °C, and then heated to 250 °C. Glass transition temperatures (Tg) were measured by analysis of the second heating ramp. There was an exothermic transition present for all three triblock polymers starting around 180 °C that was assigned as cross-linking. This transition shifted to higher temperature or was absent with the addition of 0.5 wt % butylated hydroxytoluene (BHT) as a radical inhibitor. Dynamic mechanical analysis (DMA) was performed on a press-molded sample of S0.33I0.34L0.33-49 with 1 wt % BHT using 25 mm parallel plates on a TA Instruments ARES-G2 rotational rheometer (Figure S5). The DMA temperature sweep upon heating was recorded with the following parameters: rate = 5 °C/min: ω = 3 or 1 rad/s, strain kept at less than 10%. No order−disorder transition was observed over the experimental temperature range (up to 185 °C). C
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Macromolecules Table 1. Molecular Characterization of Triblock Terpolymers in This Study polymer
Mn,PSa (kg mol−1)
Mn,PIa (kg mol−1)
Mn,PLAa (kg mol−1)
Mn,totala (kg mol−1)
Đb
f PSa
f PIa
f PLAa
S0.33I0.34L0.33-49 S0.40I0.41L0.19-39 S0.32I0.31L0.37-16
15.9 15.9 4.9
14.2 14.2 4.1
18.9 9.2 6.7
49.0 39.3 15.7
1.01 1.01 1.02
0.33 0.40 0.32
0.34 0.41 0.31
0.33 0.19 0.37
Dc−cTEMc Dc−cSAXSd 50 36 23
62 48 29
Dc−cSCFTe 61 46 21
a Determined from 1H NMR spectroscopy by end-group analysis of the SI diblock polymer. bDetermined from SEC in THF (light scattering). cThe average of at least five cylinder-to-cylinder distances in the TEM images. dCylinder-to-cylinder distance determined from SAXS. DSAXS = 2π/q* and Dc−cSAXS = 2DSAXS/√3. eCylinder-to-cylinder distance calculated by SCFT.
be monodisperse, and both the densities and the statistical segment lengths were assumed to be temperature independent. This assumption of temperature independence means that changes in temperature, such as the difference between the temperature at which we annealed the bulk samples (120 °C) and the temperature used for the calculations (150 °C), will have a negligible effect on the outcome of the calculations. The following effective binary interaction parameters were obtained from the literature65−67 and transformed appropriately for vref = 0.118 nm3 (where T is the temperature in Kelvin):
performed on a Bruker Nanoscope V Multimode 8 scanning probe microscope. Thin Film Annealing. All thermal annealing (TA) was performed in a vacuum oven at roughly 50 mTorr. The oven was equilibrated to the desired temperature as measured by an alcohol thermometer resting on the platform within the vacuum oven. Then the samples, arranged on a glass dish, were placed on the platform in the oven and put under vacuum. After the designated time, the vacuum was vented, and the samples were removed from the oven and placed on a cool glass dish or watch glass for 5−10 min before being stored for analysis. Solvent vapor annealing (SVA) was performed in a small covered glass jar at ambient temperature (∼22 °C) with a large Teflon screwcap used as a stand on which the sample could be held (Figure S10, left). The jar was filled to the bottom edge of the cap with the desired solvent, and the sample was placed on the top of the cap. The jar was sealed and held closed with a weight for the desired amount of time, after which the jar was opened and the sample was removed and stored for analysis. For these experiments, rapid drying was assumed to occur as soon as the jar was opened. Thin Film Etching. To hydrolyze the PLA, the sample was placed on the spin-coating apparatus, and any dust was removed by blowing N2 (gas) over the sample. A drop of 0.5 M NaOH solution in 40/60 (v/v) methanol/water was transferred using a pipet onto the center of the sample and left there for 1 min. Putting the drop in the center of the film helped prevent potential delamination of the film from the substrate (Figure S10, right). The sample was spun at 4000 rpm and rinsed with 40/60 (v/v) methanol/water to remove any residual NaOH. The sample was then dried with N2 (gas) and stored for analysis. Theory. SCFT for block polymers has been discussed previously.37,58,59 In brief, the SCFT framework is based on the standard Gaussian model, in which the polymer chains are treated as infinitely thin threads and their configurations are modeled using random-walk statistics.58 The local interactions driving demixing of dissimilar segments are modeled using Flory−Huggins interaction parameters χij, while short-ranged hard-core repulsions are enforced by an incompressibility constraint.58 Within this description, the SCFT formalism results in a set of coupled nonlinear equations that are solved numerically.59,60 Each self-consistent solution obtained in this way corresponds to a stable or metastable state that can be associated with a mesophase; a spatially periodic solution represents an ordered structure, while a homogeneous solution corresponds to the disordered phase. From such a solution, one can calculate relevant structural and thermodynamic quantities, such as free energy, domain spacing, and density profiles of different monomer types in the mesophase. The SCFT parameter space for the SIL triblock terpolymers involves five parameters: χPS−PIN, χPI−PLAN, χPS−PLAN, f PS, and f PI. The degree of polymerization (number of Flory−Huggins lattice sites), N, was determined using the experimental values for the volume fractions and molar masses and published values for the densities of each block (T = 140 °C): ρPS = 0.969 g/cm3,61 ρPI = 0.830 g/cm3,61 and ρPLA = 1.154 g/cm3.62 Literature values of statistical segment lengths, bPS = 0.54 nm,63 bPI = 0.60 nm,63 and bPLA = 0.70 nm,64 adjusted to the common reference volume vref = 0.118 nm3, were also used. From the statistical segment lengths, fixed values for the conformation asymmetry parameters, εI = bI/bS = 1.11 and εL = bL/bS = 1.30, were determined. For all calculations, the polymers were assumed to
χPS − PI =
58 − 0.07 T
(1)
χPS − PLA =
57 − 0.06 T
(2)
χPI − PLA =
250 − 0.41 T
(3)
At T = 150 °C and for moderate value of N = 500, the above equations yield χPS−PIN ≈ 33, χPS−PLAN ≈ 37, and χPI−PLAN ≈ 87, demonstrating that the system is frustrated. In the SCFT calculations, a grid size of 128 × 128 was used for the core−shell cylinders, while a grid size of 128 was used for the lamellae. The calculations were performed in a symmetry-constrained mode using the open-source PSCF code.59
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RESULTS AND DISCUSSION SIL triblock terpolymers were synthesized using anionic polymerization (PS and PI blocks) and ring-opening transesterification polymerization (PLA block) according to previously reported procedures.25,56 The products were characterized by 1H NMR spectroscopy, SEC, TGA, DSC, and DMA experiments (see the Supporting Information for data, Figures S1−S5). The Mn of each block as well as the overall Mn was determined from the integration of the 1H NMR spectra (Table 1). The volume fractions of each block were then calculated using the homopolymer densities at 140 °C.61,62 These data indicated that two triblock polymers of different overall size with roughly equal volume fractions of each block and a third triblock polymer from the larger diblock precursor with a lower fraction of PLA were successfully prepared (Table 1). Herein, these polymers will be referred to as S0.33I0.34L0.33-49, S0.32I0.31L0.37-16, and S0.40I0.41L0.19-39, respectively. SEC analysis for all polymers gave low molar mass dispersity values Đ < 1.1 (Table 1). The thermal characterization of the polymers produced the expected results with Td (5% mass loss) above 300 °C, Tg,PI ≈ −60 °C, Tg,PLA ≈ 53 °C, and Tg,PS ≈ 85 °C for all samples. Bulk Specimens. We prepared bulk samples of the three block polymers by solvent casting and thermal annealing. SAXS profiles indicated that all three samples displayed hexagonal symmetry (Figure 1), and the domain spacing (D) increased with overall Mn. For S0.33I0.34L0.33-49 q* = 0.116 nm−1 and DSAXS = 54 nm; for S0.40I0.41L0.19-39, q* = 0.150 nm−1 and DSAXS D
DOI: 10.1021/acs.macromol.6b02112 Macromolecules XXXX, XXX, XXX−XXX
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Figure 1. 1D SAXS patterns for S0.33I0.34L0.33-49 (a), S0.40I0.41L0.19-39 (b), and S0.32I0.31L0.37-16 (c) obtained at room temperature with q* and expected higher order reflections for hexagonal symmetry indicated. Samples were solvent cast for 1 week at ambient temperature under an inert atmosphere and thermally annealed for 16 h at 120 °C with 0.5 wt % BHT used as a stabilizer.
= 42 nm; and for S0.32I0.31L0.37-16, q* = 0.246 nm−1 and DSAXS = 26 nm. TEM images of the bulk samples were acquired after embedding in epoxy resin, staining with OsO4, and microtoming (Figure 2). OsO4 was chosen because it heavily stains the PI, lightly stains PS, and essentially leaves PLA unstained, providing excellent contrast between the three domains.24,68 The majority of the TEM images were consistent with core− shell cylinders displaying considerable non-CMC (see Figure S11 for the transverse view), though the presence of particlelike scatterers cannot be ruled out by either the TEM or SAXS data for S0.32I0.31L0.37-16. From these images, the cylinder-tocylinder spacing (Dc−cTEM) was estimated at 50 nm for S0.33I0.34L0.33-49, 36 nm for S0.40I0.41L0.19-39, and 23 nm for S0.32I0.31L0.37-16. Cylinder-to-cylinder distance was also calculated from the SAXS data (Dc−cSAXS) giving 62, 48, and 29 nm for S0.33I0.34L0.33-49, S0.40I0.41L0.19-39, and S0.32I0.31L0.37-16, respectively, in reasonable agreement with the TEM data. We also performed SCFT calculations for the three polymer samples listed in Table 1 (vide infra) and calculated the domain spacing. The domain spacings estimated from SCFT calculations are 61, 46, and 21 nm for S0.33I0.34L0.33-49, S0.40I0.41L0.19-39, and S0.32I0.31L0.37-16, respectively. For the two higher molar mass polymers, the domain spacings estimated by SCFT are in good agreement with the SAXS experiments; however, SCFT underestimates the domain spacing at low molar mass. Because SCFT is strictly valid for infinitely long chains, the assumption underlying the theory is not as rigorous as the molar mass of the polymer decreases.69 To explain the formation of hexagonal symmetry in a SIP triblock polymer where symmetric composition would typically favor a lamellar morphology, Stadler et al. posited that the curvature of domains allows for a minimization of the intermaterial dividing surface (IMDS) area between the blocks with the greatest incompatibility while concurrently increasing the contacts between the blocks of lower incompatibility.24 This explanation is consistent with our observations of a nearly CMC cylindrical PLA core within the exaggerated non-CMC
Figure 2. TEM images of S0.33I0.34L0.33-49 (a), S0.40I0.41L0.19-39 (b), and S0.32I0.31L0.37-16 (c) bulk samples after staining with OsO4. The PI domains are black, the PS domains are gray, and the PLA domains are white. The films were solvent cast over a week from THF and then annealed at 120 °C for 24 h with 0.5 wt % BHT used as a stabilizer. The scale bars represent 50 nm.
cylindrical shell of PI, thereby minimizing the IMDS area between the PLA and PI blocks. As highlighted by Gido et al.,27 these non-CMC surfaces arise due to packing frustration, minimizing the stretching energy of the chains at the cost of E
DOI: 10.1021/acs.macromol.6b02112 Macromolecules XXXX, XXX, XXX−XXX
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Macromolecules enthalpically unfavorable non-CMC surfaces.38 As SIL triblock polymers with χPS−PI = 0.069,65 χPI−PLA = 0.181,67 and χPS−PLA = 0.07566 (at 150 °C) have similar incompatibilities to the SIP system, core(PLA)−shell(PI) cylinders in a PS matrix may be anticipated to minimize the IMDS area and maximize circularity between PLA and PI while expanding the IMDS area, at the cost of circularity, between PS and PI. Gido et al. examined the source of the non-CMC nature of the cylinders in their SIP triblock terpolymer and explained that the straight edges and corners at the divide between the PI and PS surfaces are derived from the requirement that the PS fill the outer edge of the hexagonal unit uniformly despite the relatively low volume fraction of PS.27 This phenomenon, termed “packing frustration”, was shown by Matsen and Bates to theoretically influence the formation of non-CMC morphologies.38 For the SIP system, a PS block that makes up a volume fraction of less than 0.4 was predicted to display significant nonCMC of the IMDS.27 In our work, all three SIL polymers contain a PS volume fraction equal to or less than 0.4 such that amplified non-CMC could be expected for the cylindrical morphology. In contrast, for SIL polymers synthesized by Bailey and co-workers where the PS block made up a larger portion of the total polymer volume (f PS between 0.48 and 0.57), core−shell cylinder with nearly CMC were observed.52 Figure 3 shows the density profile from SCFT for the hexagonally packed core−shell cylinders (P6mm) calculated for S0.33I0.34L0.33-49 at T = 150 °C using the interaction parameters reported in eqs 1−3. It is clear from the figure that the PLA
forms virtually circular core cylinders (green) shielded by PI (red), while PS (blue) fills the matrix. Moreover, the calculations capture the experimentally observed faceting of the PS/PI interface (Figure 2), producing a surface with considerable non-CMC. As shown in Figure 3, the SCFT results anticipate the amplified non-CMC of the cylindrical morphology; however, our further calculations actually predict lamellae to be the stable phase at the chosen temperature of T = 150 °C for S0.33I0.34L0.33-49. This latter result is consistent with previously reported calculations wherein an ABC triblock polymer with equal volume fractions of the three blocks is expected to form a lamellar structure to minimize the stretching of the blocks by producing flat interfaces. These calculations are at odds with the observed morphologies (in previous work with SIP27 and herein with SIL) with χBC ≫ χAB, and the BC interface is curved to minimize the excessively unfavorable contacts between highly incompatible B and C blocks, producing cylindrical morphologies.24 Thus, it is important to examine the effects of relative differences in the interaction parameters within the triblock polymer and aim to adjust our calculations to account for these effects. Lyatskaya and Birshtein developed a strong segregation theory to examine the role of interaction parameters, particularly the ratio, χAB/χBC, on the relative stability of lamellae and cylinders in ABC triblocks.70 They predicted that for ABC triblock polymers with equal volume fractions of all blocks the lamellae-to-cylinder transition occurs at χAB/χBC ≈ 0.05. The interaction parameters reported in eqs 1−3 do not predict such a high relative incompatibility at any experimentally relevant temperatures. Even at T = 80 °C, which is below the processing temperature used for our polymers, χAB/ χBC ≈ 0.33, which is significantly higher than the ratio predicted by the theory to produce a core−shell structure. It is possible that the calculated value of χAB/χBC ≈ 0.05 is an artifact of the assumptions underlying their theory. Indeed, the effect of immiscibility between the A and C blocks, χAC, was not taken into account in their work. To investigate this issue and go beyond the approximations in the strong segregation limit, we used full mean-field theory, SCFT, to study the stability of lamellae and core−shell cylinders formed in the model ABC triblock polymers with equal volume fractions of all blocks. The SCFT calculations allow us to examine the effect of immiscibility between the A and C blocks, χAC. In these calculations, χN was used in place of χ, but N acts as a scaling factor such that the ratio χABN/χBCN is directly comparable to χAB/χBC. Figure 4 plots the free energies for the lamellar and core−shell cylindrical morphologies as a function of the ratio χABN/χBCN for different values of χACN and χBCN. For all values of χACN and χBCN considered, the lamellae-to-cylinder transition occurs in the range 0.01 < χABN/ χBCN < 0.1, which is quite small compared to the ratio, χABN/ χBCN = 0.39, calculated using the interaction parameters in eqs 1 and 3 at T = 150 °C. In fact, the range of transition values predicted by these SCFT calculations encompasses the one predicted by Lyatskaya and Birshtein.70 To compare with the results presented in Figure 3, additional SCFT calculations were carried out to produce a density profile for the core−shell cylinder morphology of an SIL polymer after adjusting the χNAB/χNBC ratio to be below the calculated lamellae-to-cylinder transition value (Figure S12, χNAB/χNBC = 0.05, χNAC = 60, and χNBC = 100). As before, the polymer was modeled with equal volume fractions of each block. In this
Figure 3. (a) Top view of the hexagonally packed core−shell cylinders (P6mm) calculated using SCFT for the polymer S0.33I0.34L0.33-49 at T = 150 °C. (b) Density profile along the line [110] (solid line in (a)) within the primitive unit cell shown by dotted lines in (a) with PS in blue, PI in red, and PLA in green. The degree of polymerization calculated using the parameters provided in the theory subsection is N = 702. The free energy of the structure relative to the disordered phase is ΔF/nkBT = −1.3205, with n being the number of chains in the system. F
DOI: 10.1021/acs.macromol.6b02112 Macromolecules XXXX, XXX, XXX−XXX
Article
Macromolecules
Figure 4. SCFT-computed free energies relative to the free energies of the disordered phase per chain, F − Fdis/nkBT, of lamellae and hexagonally packed core−shell cylinders (P6mm) formed in the model ABC triblock terpolymers with equal volume fractions of all blocks. For the different values of χACN and χBCN, the lamellae-to-cylinder transition occurs in the range 0.01 < χABN/χBCN < 0.1.
separation. In this work, the dispersity of the SIL triblock polymers is quite low (Đ ≤ 1.02) and thus would not be expected to fully account for the large difference in the SCFT predicted transition value (χABN/χBCN < 0.1) and the minimum expected transition value (χABN/χBCN = 0.39). A full explanation of the discrepancy between the experimental and SCFT studies may also include consideration of conformational asymmetry. Indeed, previous studies on diblock polymers demonstrated that conformation asymmetry not only influenced the order−order transition boundaries74,75 but also stabilized complex phases in block polymers.76 In this work, fixed values calculated from the statistical segment lengths were used as the conformation asymmetry parameters, a procedure that has previously resulted in successful prediction of experimental morphologies,39,64 though an in-depth study of conformational asymmetry in triblock polymers (and even more complicated systems) may reveal important factors for future tuning of SCFT studies. The discrepancy may also arise from the uncertainty in the interaction parameters reported in the literature. In particular, the most dominant interaction parameter used here, χPI−PLA, was obtained by fitting the order−disorder transition values for symmetric and nearly symmetric diblock polymers to the meanfield predictions.67 Because the interaction parameters obtained from different methods vary significantly, there can be
profile, there is a significant concentration of A (PS) at the B/C (PI/PLA) interface, helping to reduce the enthalpically costly B/C (PI/PLA) contacts. Also, the faceting of the shell formed by B blocks is diminished in this structure. The SCFT prediction of lamellae as the stable phase for S0.33I0.34L0.33-49, rather than the experimentally observed core− shell cylinders, leads to the critical question: why do the prediction and experiment differ? One possible source of this discrepancy is that deviations from a perfectly monodisperse system are neglected in the SCFT calculations. Previous theoretical and experimental studiesmostly focused on diblock polymershave shown that molar mass dispersity can exert a pronounced effect on the morphological map. Specifically, an increase in the polydispersity of the A block in an AB diblock polymer can shift the relationship between the expected morphology and fA, decrease the minimum χN required to achieve an ordered structure, and change domain curvature.71,72 2D real-space SCFT calculations on ABC triblock terpolymers also demonstrated that increased dispersity of an end block or midblock could significantly alter the morphological map (cases where Đ ≥ 1.5 for one of the blocks were examined).73 The effects of molar mass dispersity on morphology are attributed to the ability of a disperse sample to optimally arrange short and long chains within a domain, thereby decreasing the entropic penalty of microphase G
DOI: 10.1021/acs.macromol.6b02112 Macromolecules XXXX, XXX, XXX−XXX
Article
Macromolecules
Figure 5. AFM phase images obtained in tapping mode of S0.33I0.34L0.33-49 as-spun (a) and from solvent vapor annealing with THF for 5 min (b) and 10 min (c) on native Si. The sample thickness measured by ellipsometry was 28 nm. The scale bars represent 100 nm.
In the bulk, all three SIL polymers adopted a core(PLA)− shell(PI) cylinder morphology with exaggerated nonconstant mean curvature (non-CMC) at the PI/PS interface. Development of core−shell cylinders at such high volume fractions of the nonmatrix blocks presents interesting possibilities for the use of these polymers in nanopatterning because the diameter of the etchable domain and the distance between the domains can be readily tuned. Indeed, PLA diameters of ∼23 nm for S0.33I0.34L0.33-49 and ∼8 nm for both S0.40I0.41L0.19-39 and S0.32I0.31L0.37-16 were determined from the TEM images of the bulk materials. Even though the latter two triblock polymers have PLA domains of roughly the same size, the center-tocenter distance between the PLA domains more than doubled in going from the larger S0.40I0.41L0.19-39 to the smaller S0.32I0.31L0.37-16. Appraisal of thin films of these triblock polymers was therefore pursued to determine the morphology and domain size tunability on various substrates. Thin Film Specimens. Dilute polymer solutions (1 wt %) were spin-coated onto native Si oxide and HMDS-modified Si substrates. For each polymer, the morphology and the orientation of the domains did not vary significantly between the two substrates, even after thermal or solvent vapor annealing techniques. Because of dewetting and delamination of S0.32I0.31L0.37-16 thin films under the tested annealing conditions, the larger polymers became the focus of this work (see Figure S13 for the AFM images of S0.32I0.31L0.37-16). In AFM tapping mode, the soft PI will have the most negative phase and appears as the darkest domain in the phase images while both PS and PLA appear as lighter areas.51 The as-spun polymers revealed microphase separation of the domains, though unlike S0.62I0.13L0.25-65 reported by Kubo et al., the triblock polymers reported herein did not immediately develop core−shell cylinders oriented perpendicular to the substrate surface upon spin coating.53 As-spun AFM phase images of S0.33I0.34L0.33-49 showed a mix of cylindrical and wormlike features with dark, round, PI domains throughout the matrix on both native and HMDS-modified Si substrates (Figure 5a and Figure S14a). The as-spun AFM phase image of S0.40I0.41L0.19-39 on native Si showed circular bright domains and dark domains throughout the matrix (Figure S15a). Figure 5 depicts phase images of S0.33I0.34L0.33-49 after SVA with THF for 5 or 10 min on native Si. The phase images exhibit both perpendicularly oriented core−shell cylinders and wormlike structures wherein discontinuous PI (dark) domains border bright core domains and a somewhat less bright matrix surrounds these core−shell structures. Over longer SVA times,
uncertainties upward of 100% at experimentally relevant temperatures.77 Yet this could not account for the order of magnitude difference between the ratio of χNAB/χNBC in the experimental system (∼0.39) versus the ratios predicted to yield a stable core−shell cylinder morphology by either SCFT or strong segregation theory (
