Bulk and Thin Film Morphological Behavior of Broad Dispersity Poly

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Bulk and Thin Film Morphological Behavior of Broad Dispersity Poly(styrene-b-methyl methacrylate) Diblock Copolymers Joan M. Widin,† Myungwoong Kim,‡ Adam K. Schmitt,† Eungnak Han,‡ Padma Gopalan,†,‡ and Mahesh K. Mahanthappa*,†,‡ †

Department of Chemistry, University of WisconsinMadison, 1101 University Ave., Madison, Wisconsin 53706, United States Department of Materials Science & Engineering, University of WisconsinMadison, 1509 University Ave., Madison, Wisconsin 53706, United States

‡

S Supporting Information *

ABSTRACT: We describe the morphological implications of broad molecular weight dispersity on the bulk and thin film self-assembly behavior of seven model poly(styrene-blockmethyl methacrylate) (SM) diblock copolymers. Derived from sequential nitroxide-mediated polymerizations, these unimodal diblock copolymers are comprised of narrow dispersity S blocks (Đ ≤ 1.14) and broad dispersity M blocks (Đ ∼ 1.7) with total molecular weights Mn,total = 29.2−42.9 kg/mol and M volume fractions f M = 0.35−0.63. Small-angle X-ray scattering (SAXS) and transmission electron microscopy (TEM) analyses demonstrate that these diblock copolymers microphase separate into lamellar and cylindrical morphologies with substantially larger microdomain spacings at lower overall molecular weights as compared to their narrow dispersity analogues. The observed microphase-separated melt stabilization is also accompanied by a substantial shift in the lamellar phase composition window to higher values of f M. In thin films, these polydisperse copolymers form perpendicularly oriented morphologies with modest degrees of lateral order on substrates functionalized with P(S-ran-MMA) neutral polymer brush layers.

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INTRODUCTION Block copolymer self-assembly into nanoscale morphologies with precise periodicities in both bulk and thin film geometries offers opportunities for the development of new commodity thermoplastics, thermoplastic elastomers, and adhesives1,2 as well as value-added nanostructured templates for nextgeneration microelectronic devices and high-density magnetic storage media.3−5 Block copolymer melt-phase morphologies reflect a delicate balance of the enthalpic repulsions between the constituent chemically dissimilar homopolymer segments and the unfavorable entropy of chain stretching. This force balance thermodynamically directs formation of well-known equilibrium morphologies such as spherical, cylindrical, and lamellar phases.6 Monodisperse diblock copolymer phase behavior is parametrized by two variables: the copolymer composition or volume fraction (fA = 1 − f B) and the segregation strength (χABN), where N is the overall degree of polymerization and χAB is the temperature-dependent Flory− Huggins interaction parameter that quantifies the incompatibility between the A and B blocks. 7,8 Geometrically constraining a block copolymer in a thin film geometry introduces additional variables that control microdomain structure, orientation, and lateral order with respect to the underlying substrate.4,9−12 In the context of directed selfassembly of block copolymer templates for microelectronics fabrication, much work has focused on manipulating the film © XXXX American Chemical Society

thickness, the surface energies of each block, and the interfacial interactions of each block with the substrate to generate useful, perpendicularly (vertically) oriented morphologies with high degrees of long-range lateral order.13,14 During the past 25 years, synthetic polymer chemists have developed a variety of controlled radical polymerization (CRP) techniques,15−18 which enable access to designer soft materials bearing unusual combinations of chemical functionalities. For example, these synthetic methodologies facilitate enchainment of monomers containing diverse molecular motifs useful in stimuli-responsive polymers for bioanalytical sensing19,20 and in etch-resistant materials for block copolymer nanolithography.21,22 In contrast to anionic polymerizations that typically produce narrow dispersity block copolymers, CRP methods often introduce significant chain length and composition dispersities into the resulting copolymers. The effects of segmental dispersity on the melt and thin film phase behavior of block copolymers are only partially understood, as molecular weight homogeneity has long been considered a prerequisite for well-defined copolymer self-assembly.23 Access to new, polydisperse block copolymers derived from CRPs and tandem polymerizations has spurred investigations of Received: March 1, 2013 Revised: May 2, 2013

A

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(TIPNO) was synthesized as previously reported.47 Styrene and methyl methacrylate were stirred over Brockman Type I basic alumina for 1 h and gravity filtered to remove inhibitors immediately prior to use. In some cases, methyl methacrylate (MMA) was additionally purified by stirring over CaH2 for 12 h followed by vacuum distillation. Polymer compositions were determined by quantitative 1H NMR spectroscopy in CDCl3 on a Bruker AC+ 300 spectrometer, using a pulse repetition delay of 10 s. All spectra were referenced relative to tetramethylsilane (TMS). Size Exclusion Chromatography (SEC). SEC analyses employed a Viscotek GPCMax System equipped with two Polymer Laboratories Resipore columns (250 mm × 4.6 mm), a differential refractometer, light scattering module (7° and 90° detection), a differential viscometer, and a UV/vis detector. Tetrahydrofuran was used as the mobile phase at 40 °C with a 0.8 mL/min flow rate. The instrument was calibrated using 10 narrow dispersity polystyrene standards with Mn = 0.580−377.4 kg/mol (Polymer Laboratories, Amherst, MA). This calibration curve was converted to an equivalent calibration curve for poly(methyl methacrylate) using reported Mark−Houwink constants.48 Synthesis of 2,2,5-Trimethyl-3-(2-oxyethyl isobutryl)-4-phenyl-3-azahexane (1). 1 was synthesized by a variation on the method of Banik et al.49 Cu powder (873 mg, 13.7 mmol), CuBr (99 mg, 0.681 mmol), and PMDETA (236 mg, 1.36 mmol) were combined in a Schlenk flask and degassed by three freeze−thaw cycles. A combination of TIPNO (1.00 g, 4.54 mmol) and ethyl 2bromoisobutyrate (1.33 g, 6.81 mmol) was degassed by three freeze−thaw cycles in a second Schlenk flask followed by dilution with anhydrous toluene (20 mL). The latter solution was transferred to the Cu/CuBr/PMDETA mixture and stirred vigorously for 18 h at 22 °C. The reaction mixture was exposed to air, diluted with CH2Cl2 (10 mL), and filtered through a plug of silica gel with ethyl acetate as the eluent to remove the metallic impurities. The resulting pale yellow filtrate was concentrated and purified by silica gel column chromatography (10:1 hexanes/ethyl acetate). Residual ethyl 2bromoisobutyrate was removed by coevaporation of 1 with toluene under vacuum. Typical yields: 60−70%. 1H NMR (300 MHz, CDCl3 with added PhNHNH2) δ (ppm): 7.60−7.25 (m, 5H, Ph), 4.17 (q, 2H, J = 7.0 Hz, CH3CH2OCO), 4.15 (q, 2H, J = 7.2 Hz, CH3CH2OCO), 3.69 (d, 1H, J = 11.1 Hz, NCH[CH(CH3)2](Ph)), 3.41 (d, 1H, J = 10.3 Hz, NCH[CH(CH3)2](Ph)), 2.40 (m, 1H, NCH[CH(CH3)2](Ph), 1.83 (m, 1H, NCH[CH(CH3)2](Ph)), 1.62 (s, 6H, OC(CH3)2CO), 1.59 (s, 6H, OC(CH3)2CO), 1.55 (s, 6H, OC(CH3)2CO), 1.48 (s, 6H, OC(CH3)2CO), 1.31 (t, 3H, J = 7.1 Hz, CH3CH2OCO), 1.28 (t, 3H, J = 7.1 Hz, CH3CH2OCO), 1.24 (d, 6H, J = 6.7 Hz, NCH[CH(CH3)2](Ph)), 1.16 (d, 6H, J = 6.2 Hz, NCH[CH(CH3)2](Ph)), 0.93 (s, 9H, (CH3)3N), 0.81 (s, 9H, (CH3)3N), 0.73 (d, 6H, J = 6.5 Hz, NCH[CH(CH3)2](Ph)), 0.40 (d, 6H, J = 6.8 Hz, NCH[CH(CH3)2](Ph)). Representative Synthesis of Poly(styrene) Homopolymer. A mixture of 1 (300.0 mg, 0.933 mmol) and purified styrene (20.0 g, 0.193 mol) was degassed by three freeze−thaw cycles in a Schlenk flask and placed under N2(g). The flask was placed in an oil bath at 125 °C. After 6 h, the polymerization reaction was arrested by rapidly cooling the flask in a 0 °C bath for 15 min. The resulting polymer was precipitated in MeOH and dried in vacuo. SEC: Mn = 15.9 kg/mol, Mw/Mn = 1.12 (against PS standards). Representative Parallel Syntheses of SM Diblock Copolymers. PS macroinitiator (Mn = 15.9 kg/mol, Mw/Mn = 1.12, 1.00 g, 63.0 μmol) was weighed into each of six vials in a glovebox. Purified MMA (3.85 mL, 38.0 mmol) and TIPNO (194 μL of a 50 mg/mL solution in MMA, 44.1 μmol) were added to each vial. The vials were placed in a multiwell stir plate at 125 °C for 0.5−9 h until the desired MMA monomer conversion was achieved. The reactions were stopped by removing the reaction vials from the glovebox and immersing them in a 0 °C bath for at least 15 min. The resulting polymers were precipitated in MeOH and dried in vacuo. Nitroxide End-Group Removal from SM Diblock Copolymers. Each diblock copolymer (200−300 mg) was dissolved in a 2:1 (v/v) solution of PhNHNH2/toluene (3 mL) in a glass pressure tube

the consequences of broad and continuous segmental dispersity on their phase behaviors.24−27 Matsushita and co-workers have described the morphologies adopted by multicomponent blends of narrow dispersity diblock and triblock copolymers,28−30 yet these multimodal blends exhibit notably different behaviors from broad dispersity copolymers with unimodal and continuous molecular weight distributions. AB diblock31−33 and (AB)n multiblock27 copolymers with broad dispersities in both blocks (“symmetrically polydisperse” copolymers) exhibit composition-dependent phase windows similar to their monodisperse analogues.8 Complementary studies of “asymmetrically polydisperse” AB diblocks34−36 and ABA triblocks,25,37,38 in which only one segment is polydisperse, demonstrate shifts in the composition-dependent morphology windows. In symmetrically and asymmetrically polydisperse copolymer melts, microphase separation has been observed at lower than expected overall molecular weights with attendant increases in the order−disorder transition temperature (TODT). The microphase-separated domains are also significantly dilated as compared to their narrow dispersity analogues.25,32,35,37,38 In some cases, segmental dispersity enables formation of stable, nonclassical bicontinuous morphologies.36,39 Theoretical treatments of broad dispersity copolymer melts capture a subset of these experimentally observed phenomena.40−44 The effect of chain length dispersity on the morphological behavior of block copolymer thin films is less well-studied. Only one report has described the effects of broad and continuous copolymer dispersity on high molecular weight poly(methyl methacrylateb-n-butyl acrylate) (PMMA−PnBA) diblock copolymer thin film morphologies and their orientations relative to the substrate.45 Comprised of a broad dispersity PnBA block (Đ = 1.31−1.76) and a narrow dispersity M block, these copolymer thin films self-assemble with high degrees of lateral order reminiscent of strongly segregated, narrow dispersity copolymers. In this paper, we describe the phase behavior of seven weakly segregated poly(styrene-b-methyl methacrylate) (SM) diblock copolymers as bulk materials and in thin films. Prepared by nitroxide-mediated polymerizations (NMP), these copolymers are comprised of a narrow dispersity S block (Đ = Mw/Mn ∼ 1.1) and a broad dispersity M block (Đ ∼ 1.7). Using a combination of small-angle X-ray scattering (SAXS) and transmission electron microscopy (TEM), we determine the bulk morphologies and characteristic microdomain spacings of these polymers. From this foundation, we investigate the thin film self-assembly behavior of these SM copolymers on substrates modified with hydroxyl-terminated P(S-r-MMA) random copolymer brushes. By varying the surface brush composition, we identify neutral surface wetting conditions that drive formation of perpendicularly oriented microdomains, as determined by scanning electron microscopy (SEM). These studies demonstrate the ability to direct the assembly of these polydisperse SM copolymers into well-defined perpendicular lamellar and cylindrical morphologies within the geometric confinement of thin films.

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EXPERIMENTAL SECTION

Materials. All reagents and monomers were obtained from SigmaAldrich Chemical Co. (Milwaukee, WI), and they were used as received unless stated otherwise. All reagent grade solvents were used without purification. CuBr was purified by an established procedure,46 and 1,1,4,7,7-pentamethyldiethylenetriamine (PMDETA) was vacuum distilled prior to use. 2,2,5-Trimethyl-4-phenyl-3-azahexane-3-nitroxide B

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and heated to 118 °C for 8 h. Upon cooling, the solution was diluted with chloroform (∼50 mL) and washed with water and dilute hydrochloric acid (∼2 M) in a separatory funnel. Separation and concentration of the organic fraction yielded an orange solid, likely due to the presence of residual TIPNO. The polymer was further purified by precipitation from dichloromethane (10 mL) into hexanes (500 mL) and subsequent vacuum drying. The isolated polymer retained a faint orange hue, which we tentatively attribute to the presence of a small amount of free TIPNO. No further analyses were conducted to identify the exact origin of this color, and these polymers were used as obtained. Differential Scanning Calorimetry (DSC). DSC analyses were performed using a TA Instruments DSC-100 over a temperature range T = 25−200 °C, using a heating and a cooling ramp rate of 5 °C/min. The thermal history of the sample was erased in the first cycle by heating to 200 °C. All reported thermal analyses derive from the second heating cycle. Only one glass transition temperature Tg = 88− 89 °C was observed in all SM samples. Thermogravimetric Analysis (TGA). Thermogravimetric analyses were performed using a TA Instruments Q500 thermogravimetric analyzer. Polymer sample weight loss was measured over the temperature range T = 25−600 °C using a ramp rate of 10 °C/min. Small-Angle X-ray Scattering (SAXS). Synchrotron small-angle X-ray scattering measurements were performed at the 5-ID-D beamline of the DuPont-Northwestern-DOW Collaborative Access Team Synchrotron Research Center at the Advanced Photon Source (Argonne, IL), using a beam energy of 16 keV (λ = 0.7293 Å−1) and a 3.067 m sample-to-detector distance. Two-dimensional SAXS patterns were recorded on a MAR-CCD detector (133 mm diameter active circular area) with 2048 × 2048 pixel resolution. Temperaturedependent SAXS analyses utilized a Linkam DSC stage, with a 5 min pre-equilibration delay prior to data collection at a given temperature (typical exposure times ∼1 s). Prior to SAXS analysis, each sample was annealed for 5−10 min at 220 °C and slowly cooled to 25 °C. Transmission Electron Microscopy (TEM). Portions of the synchrotron SAXS samples were embedded in medium-soft epoxy (EMBed-812, Electron Microscopy Sciences), prior to sectioning at 22 °C using a diamond knife to produce ∼80−95 nm thick samples. The sections were placed on 400-mesh Cu grids and exposed to the vapor above a 0.5 wt % aqueous RuO4 solution for 36 h to preferentially stain the S block. Energy-filtered TEM micrographs were acquired with a LEO 912 EFTEM using a 120 kV accelerating voltage in bright-field mode. Fabrication of Nonpreferential Substrates and Thin Film Block Copolymer Self-Assembly. Solutions of hydroxyl-terminated P(S-r-MMA) random copolymers9 in toluene (1% w/w) were spincoated at 2000 rpm onto precleaned (piranha solution treated) silicon wafers and annealed under vacuum at 220 °C for 6 h. Annealed wafers were washed with toluene to remove ungrafted chains. On these random copolymer-modified wafers, solutions of SM-46 or SM-51 in toluene (0.7% w/w) were spin-cast at 4000 rpm and annealed under vacuum at 220 °C for 3 h. Block copolymer film thicknesses were measured by ellipsometry (Rudolph Research Auto EL). Top-down scanning electron micrographs of the thin film morphologies were acquired with a LEO-1550 VP field-emission SEM operating with an accelerating voltage of 1 kV.

Scheme 1. Sequential NMP Synthesis of Polydisperse Yet Unimodal SM Diblock Copolymers Initiated by 1

superior control in styrene homopolymerizations as compared to the TEMPO-based analogue.50 Using NMP of styrene at 125 °C initiated by 1, we synthesized three narrow dispersity, TIPNO-terminated S homopolymers with Mn = 15.9−18.1 kg/ mol and Đ = Mw/Mn = 1.10−1.14. Sequential chain extension block copolymerizations of MMA from these S macroinitiators at 125 °C produced complex mixtures of the desired broad dispersity SM diblock copolymers with a significant fraction of unreacted parent S (Figure 1a). We ascribed this result to uneven initiation of the S macroinitiator chain ends, followed by rapid and poorly controlled methacrylate monomer propagation. In order to facilitate uniform chain initiation and to reduce the rate of MMA polymerization, we studied the effects of adding variable amounts of free TIPNO to these chain extension reactions. Adding 70 mol % TIPNO (relative to the S macroinitiator) to these chain extension reactions yielded polydisperse, unimodal SM diblock copolymers with no evidence of unreacted parent S homopolymer (Figure 1b). SEC analyses with UV−vis detection at λ = 254 nm confirmed the uniform incorporation of the S blocks across the entire copolymer molecular weight distribution, indicating the absence of any M homopolymer. Since this polymer system is not susceptible to analytically useful degradation reactions that enable the direct characterization of the dispersities of each segment of the copolymer, we conducted MMA homopolymerizations initiated by 1 at the same [monomer]/[initiator] ratios as in the chain extension reactions. These reactions demonstrate that Đ ∼ 1.7 for the M blocks (see Figure S2), as expected from the well-studied PMMA chain termination by disproportionation.51 In lieu of complex deconvolutions of the SEC data subject to assumptions regarding the hydrodynamic volumes of the constituent homopolymer blocks,34 we assume that the dispersities of the M segments in our block copolymers are Đ ∼ 1.7. Under these optimized polymerization conditions, a series of seven SM diblock copolymers were synthesized with Mn,total = 29.2−42.8 kg/mol with Đ = 1.25−1.57 (from SEC against PS standards) and M block volume fractions f M = 0.35−0.63. Total copolymer molecular weights (Mn,total) and volume-normalized degrees of polymerization (N) were calculated using Mn,SEC for the parent S macroinitiator, the copolymer composition derived

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RESULTS AND DISCUSSION Polydisperse SM Diblock Copolymer Synthesis. Unimodal yet polydisperse SM diblock copolymers were synthesized by sequential nitroxide-mediated block copolymerizations of styrene and MMA (Scheme 1). NMP of MMA is well-known to produce broad dispersity PMMA, whereas the NMP of styrene is well controlled.17 TIPNO-derived alkoxyamine initiator 1 was prepared by atom transfer radical coupling of TIPNO47 with ethyl 2-bromoisobutyrate catalyzed by (PMDETA)CuBr with Cu(0) as the terminal reducing agent in 60−70% yield.49 This TIPNO-based initiator 1 furnishes C

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density differences between PS and PMMA homopolymers, we recalculated χS‑MMA based on previous work by Russell et al. in which they determined χS‑MMA by fitting the correlation hole scattering associated with a disordered, deuterated SM diblock copolymer.53 We find that χS‑MMA(T) = 0.0209 + 2.87/T with respect to the 118 Å reference volume (see Supporting Information for details). From this expression, we calculate that all seven of our samples are weakly segregated with χS‑MMAN = 10.4−14.8 at T = 140 °C. Thermal Properties of Polydisperse SM Diblock Copolymers. The thermal stability of PMMA produced by free radical polymerizations depends upon the identity of the end groups generated by chain termination.54,55 We initially observed some irreproducibility in the melt-phase morphologies and domain spacings associated with polydisperse SM copolymers derived from NMP after thermal annealing at 220 °C. Thus, we investigated the thermal stabilities of these diblocks using thermogravimetric analysis (TGA). Upon heating SM-39 from 25 to 400 °C at a constant rate of 10 °C/min under a flow of nitrogen, we observed two broad thermal decomposition features near T = 215 and 275 °C with complete decomposition by ∼400 °C, as shown in the TGA trace given in Figure 2. NMP is known to poorly control the

Figure 1. (a) Sample SEC trace of a mixture of a broad dispersity SM diblock copolymer and a parent S macroinitiator derived from sequential chain extension block copolymerization of MMA in the absence of added TIPNO. (b) Overlay of SEC traces for unimodal yet polydisperse SM-46 (Mn = 35.3 kg/mol, Đ = 1.32; blue) produced by sequential NMP of MMA at 125 °C initiated from a S macroinitiator (Mn = 18.1 kg/mol, Đ = 1.14; red) in the presence of 70 mol % TIPNO.

from quantitative 1H NMR spectroscopy, and the somewhat arbitrary 118 Å reference volume (see Supporting Information for details). The volume fractions ( f M) of MMA in the copolymers were calculated using the previously reported homopolymer melt densities at 140 °C.52 Table 1 summarizes the detailed molecular characteristics of these materials. For convenience, we adopt the naming convention SM-X, where X is the copolymer volume fraction f M. In order to account for

Figure 2. Thermogravimetric analysis (TGA) of SM-39 diblock copolymer before (dashed blue line) and after refluxing the polymer with PhNHNH2 (solid red line), indicating enhanced thermal stability upon nitroxide end-group cleavage.

polymerization of MMA, since nucleophilic nitroxide radicals catalyze the disproportionation of propagating M chain ends to furnish olefin-terminated (unsaturated) and hydrogen-termi-

Table 1. Molecular Characteristics of Polydisperse SM Diblock Copolymers S block sample SM-35 SM-39 SM-44 SM-46 SM-51 SM-56 SM-63

Mn,SECa

(kg/mol)

18.1 18.1 18.1 18.1 15.9 14.6 14.6

M block Đ

a

1.14 1.14 1.14 1.14 1.12 1.10 1.10

block copolymer

b

Mn,totalb

Mn,NMR (kg/mol) 11.1 12.8 16.1 17.2 18.5 21.0 28.2

(kg/mol)

29.2 30.9 34.2 35.3 34.4 35.6 42.8

Đa

f Mc

Nd

χNe

1.25 1.33 1.34 1.32 1.44 1.44 1.57

0.35 0.39 0.44 0.46 0.51 0.56 0.63

374 395 434 447 432 445 530

10.4 11.0 12.1 12.4 12.0 12.4 14.8

Determined by size exclusion chromatography (SEC) in THF at 40 °C against a polystyrene calibration curve. bCalculated from quantitative 1H NMR spectroscopy in CDCl3, using the Mn,SEC for the S block. cVolume fraction f M determined from 1H NMR using melt homopolymer densities at 140 °C.52 dTotal volume normalized degree of polymerization of the block copolymer calculated using the 118 Å3 reference volume to account for density differences in the constituent homopolymers. eSegregation strength calculated using χSM = 0.02785 at 140 °C (see Supporting Information for details). a

D

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Scheme 2. Removal of Nitroxide End Groups from SM Diblock Copolymers Using Phenylhydrazine

nated (saturated) chain ends along with the free nitroxide.55,56 Since disproportionation competes with the reversible activation−deactivation polymerization process, we reasoned that NMP of MMA yields three distinct chain populations having TIPNO-, olefin-, and H-terminated chain ends. We speculate that the initial thermal degradation wave near 215 °C arises from activating the TIPNO-terminated chains above the ceiling temperature of PMMA (Tc ∼ 220 °C for neat monomer51), which triggers their spontaneous depolymerization. Further heating triggers the decomposition of the remaining diblock copolymer chains over a broad temperature range spanning 275−400 °C. The precise mechanism of this second decomposition process is unknown. In order to prevent these decomposition reactions, we studied thermal activation of the nitroxide-terminated chain ends in the presence of chain transfer agents to cleave the nitroxide moieties and irreversibly cap the chains with H atoms. Given the potential utility of block copolymer nanopatterning in semiconductor and microelectronic device fabrication, we sought to avoid the use of metallic reagents such as (nBu)3SnH.55 We instead investigated the use of PhNHNH2 as a chain transfer agent according to the reaction shown in Scheme 2. Phenylhydrazine is known to readily reduce nitroxide radicals to the corresponding hydroxylamines.47 Treatment of SM diblock copolymers with PhNHNH2 in toluene at 118 °C for 8 h yielded polymers with dramatically enhanced thermal stabilities. TGA analysis of SM-39 after treatment with PhNHNH2 demonstrated that this diblock remains stable up to 290 °C (5 wt % loss in Figure 2). Quantitative 1H NMR analysis of the samples after treatment with PhNHNH2 showed that the chemical composition of the polymer does not change after this nitroxide end-group removal protocol. SEC analyses further confirm that the Mn and Đ are unaffected by this modification reaction. Consequently, all SM diblock copolymer samples were subjected to this postsynthetic modification reaction prior to morphological analysis. Melt-Phase Characterization. The self-assembled morphologies of polydisperse SM diblock copolymer melts were investigated using a combination of synchrotron small-angle Xray scattering (SAXS) and transmission electron microscopy (TEM). Samples were thermally annealed at 220 °C for 5 min to enhance chain mobility and to enable microdomain ordering, followed by cooling to 25 °C for SAXS analyses. SAXS analyses clearly demonstrate that SM-39 and SM-44 form a hexagonally packed cylinders morphology at 140 °C (Figure 3), albeit with only modest long-range order as evidenced by the broad peaks and the limited number of higher order scattering maxima. SM46 exhibits a broad SAXS maximum at q* = 0.0171 Å−1 (d = 36.8 nm at 140 °C) along with broad reflections at √3q*, √4q*, √7q*, and √9q*, possibly characteristic of a hexagonally packed cylinders morphology. However, the presence of additional broad scattering maxima in the azimuthally integrated intensity pattern suggests that this

Figure 3. Azimuthally integrated synchrotron SAXS intensity profiles at 140 °C for polydisperse SM diblock copolymers listed in Table 1. The black triangles indicate the positions of scattering maxima expected for a hexagonally packed cylinders morphology, whereas the diamonds indicate the peak positions for a lamellar morphology. SM46 (open circles) exhibits broad reflections that are ascribed to a disorganized microphase-separated structure between cylinders and lamellae (see text for details).

sample sits in a transition zone between hexagonally packed cylinders and lamellae (open circles in Figure 3). On the basis of the SAXS data and TEM analyses (Figure 4), we assign this sample as a disorganized microphase-separated structure between cylinders and lamellae. Temperature-dependent SAXS analyses up to 220 °C indicate that this sample remains microphase separated (see Supporting Information Figure S3), with no indications of additional peak broadening associated with proximity to the order−disorder transition (ODT). SM-51

Figure 4. TEM micrographs of (a) SM-46 exhibiting a disorganized yet microphase separated morphology with both lamellar and cylindrical features with only short-range order and (b) the lamellar phase of SM-51. The dark regions in the TEM micrograph correspond to poly(styrene) stained with RuO4. E

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exhibits scattering maxima located at q* = 0.0220 Å−1 (d = 28.6 nm at 140 °C), √4q*, √9q*, and √16q* consistent with a lamellar morphology, as confirmed by TEM (Figure 4). The suppressed intensities of the Bragg reflections corresponding to √4q* and √16q* are expected based on the near-symmetric composition of the sample ( f M ∼ 0.51).57 A complete list of the observed morphologies and principal domain spacings (d) of these SM samples is given in Table 2.

Using expressions developed by Matsen and Bates,58 we calculated the theoretical domain spacings (dthy in Table 2) for perfectly monodisperse SM copolymers with equivalent N values (or Mn,total) in the intermediate segregation regime for comparison with our experimental observations. These calculated values agree well with the scaling relation reported by Anastasiadis et al.,60 upon using the corrected value of χS‑MMA (vide supra). On this basis, we conclude that the domain spacings for the broad dispersity SM copolymers are 20−70% larger than those expected for the narrow dispersity analogues. One might argue that dispersity in the M block implies the presence of long M chains that dictate the characteristic ordering length scale of these microphase-separated melts. Replacing Mn,total with Mw,total in the theoretical domain spacing calculation at intermediate segregation improves the agreement with our experimental observations. However, we do not quantitatively recover the observed domain spacing through this modification. Therefore, we conclude that M segmental dispersity substantially stabilizes the microphase-separated melts and dilates the microdomains. Known theories42,61 qualitatively capture the behaviors of these broad dispersity block copolymers, yet quantitative comparisons cannot be drawn due to differences between the specific copolymers reported herein and those studied in previous theoretical treatments. The composition-dependent phase window shifts and microphase-separated melt stabilization for broad dispersity SM diblock copolymers may be readily rationalized in terms of the intrinsic composition and chain length dispersities of these copolymer samples. Some SM chains have very short M blocks with respect to the average value Mn,MMA, while other chains have M blocks substantially larger than Mn,MMA. For chains bearing M blocks shorter than some critical molecular weight determined by χSM(T), it becomes entropically favorable for them to desorb (pull out) from the microphase-separated interface. These chains act as effective “S homopolymers” that partition into and dilate the S-rich microdomains.42 This effect decreases the effective amount of M in the morphology relative to that expected based on the bulk chemical composition of the copolymer. Of the SM chains remaining at the domain interface, chains with intermediate length M blocks crowd the interface and shield the longer M chains in the distribution from unfavorable enthalpic contacts with the S chains. This interfacial crowding or “co-surfactancy”62,63 alleviates the need for the longer M chains to stretch, allowing them to relax. This relief of long M chain stretching increases the configurational entropy of the ensemble, while further decreasing the volume filled by the broad dispersity M block. Therefore, the composition-dependent phase windows shift in a manner consistent with the broad dispersity M block acting as if it is much smaller than expected based on the copolymer chemical composition. Consequently, the lamellar window shifts to higher values of f M. The increased configurational entropy associated with the relief of chain stretching also entropically stabilizes the microphase-separated melt.34,42 Therefore, the (χN)ODT required for microphase separation decreases as compared to narrow dispersity copolymer melts. Since the χS‑MMA(T) is inversely dependent on temperature, the latter result implies that TODT increases in these systems. SM microdomain dilation likely stems from a combination of the packing constraints imposed by the longest M chains in the molar mass distribution and the presence of effective “S homopolymer” that swells the S domains. We noted previously

Table 2. Morphological Characteristics and Homopolymer Swelling of Polydisperse SM Diblock Copolymers at 140 °C sample

morphologya

db (nm)

dthyc (nm)

wt % Sd

SM-35 SM-39 SM-44 SM-46 SM-51 SM-56 SM-63

H H H H/L* L L L

31.8 37.7 37.4 36.8 28.6 27.0 27.0

20.6 22.1 24.4 25.3/19.9 19.6 20.0 22.6

10.4 8.0 5.7 5.0/5.1 4.2 3.1 1.8

a H = hexagonally packed cylinders; L = lamellae; H/L* = disorganized morphology between H and L. bPrincipal domain spacing determined from synchrotron SAXS at 140 °C. cdthy = theoretical principal domain spacing for a perfectly monodisperse SM diblock copolymer calculated from ref 58 at T = 140 °C. dCalculated weight fraction of polymer chains acting as effective “S homopolymer” (see text and Supporting Information for details).

The different phase assignments for SM-44 and SM-51, along with the disorganized structure exhibited by SM-46, suggest that the cylinders/lamellae phase boundary is located near f M ∼ 0.45−0.50. When f M < 0.45, we observe a cylindrical morphology by SAXS in which the microphase-separated interfaces preferentially curve toward the broad dispersity M block. This observation is consistent with previous experiments that described asymmetric shifts in the lamellar phase window for asymmetrically polydisperse AB diblocks.24 Self-consistent mean-field theory (SCMFT) predictions of the phase behavior of conformationally symmetric, AB diblock copolymers with χN = 15 anticipate that the cylinders/lamellae phase boundary shifts from f B ∼ 0.40 for perfectly monodisperse copolymers to f B ∼ 0.50 upon broadening the B block dispersity to ĐB = 1.5.42 Therefore, broad M block dispersity shifts the cylinders/ lamellae phase boundary to higher values of f M in a manner consistent with previous reports. In spite of the weakly segregated nature of the broad dispersity SM diblock copolymers reported in this study (Table 1), we note that their order−disorder transition temperatures (TODT) are in excess of 240 °C and that the microphaseseparated domain spacings are quite large (d = 27−38 nm). For conformationally symmetric, monodisperse diblock copolymers, SCMFT predicts that (χN)ODT ∼ 12 at f = 0.40 (neglecting fluctuation effects).8 Thus, it is somewhat surprising that broad dispersity SM diblocks with f M = 0.35−0.44 and χN = 10.4−12.1 microphase separate at all. A few previous reports have noted that segmental dispersity can stabilize microphase separation of block copolymer melts at unusually low segregation strengths.25,32,35,59 The observed microdomain spacings are also quite large, in spite of the modest overall molecular weights of these SM copolymers (Table 2). Domain dilation has been experimentally observed in other broad dispersity block copolymer systems, particularly as the overall segregation strength χN of the system decreases.25,32−34,38 F

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Figure 5. SEM micrographs of thin films of (a) SM-46 and (b) SM-51 on Si substrates modified with hydroxyl-terminated P(S-r-MMA) brushes where the styrene mole fractions ranges FSt = 0.50−0.78. Thin films exhibiting the best alignment and order for each diblock copolymer are indicated by a red box in (a) and a blue box in (b). All scale bars are 200 nm.

that theoretical domain spacings calculated using Mw,total (instead of Mn,total) for the SM copolymers underestimate the experimentally observed domain spacings, suggesting that other effects contribute to domain dilation. In order to account for the additional domain swelling, we used Matsen’s formalism42 for calculating the amount of effective “S homopolymer” present in each of our broad dispersity SM samples (see Supporting Information for details). In our specific system, this formalism allows for the calculation of the critical chain length below which the broad dispersity M blocks pull out of the microphase-separated interface and act as “S homopolymer”. The calculated amount of “homopolymer” present in each SM diblock copolymer sample is listed in Table 2. Close inspection of the domain spacings listed in Table 2 reveals that the ratio d/ dthy decreases with increasing segregation strength χN. Consistent with this trend, the calculated amount of effective “homopolymer” decreases as the f M and χS‑MMAN increase. Thus, we ascribe the observed domain dilation to the presence of this substantial “homopolymer” fraction coupled with the presence of high molecular weight M segments due to the broad segmental dispersity. SM Thin Film Self-Assembly. The self-assembly behavior of block copolymer thin films can be subtly controlled by tailoring the interfacial energy between the surface and each block.10 Mansky et al. showed that modifying a substrate with a P(S-r-MMA) random copolymer brush furnishes a simple method to precisely control the polymer−surface interactions of narrow dispersity SM diblock copolymers.9 By varying the random copolymer composition, one can minimize the interfacial energies for both blocks to access a nonpreferential or neutral wetting surface. These neutral surfaces allow perpendicular alignment of the block copolymer microdomains with respect to the substrate. In this initial study, we utilized hydroxyl-terminated P(S-r-MMA) random copolymers to tailor the interfacial energy of Si substrates for studies of the thin film

self-assembly of the polydisperse SM diblock copolymers. Thin films of SM-46 and SM-51 were cast on the random copolymer-modified substrates. The mole fraction of styrene (FSt) in the random copolymer was varied from FSt = 0.50− 0.78. Figures 5a and 5b present SEM micrographs of SM-46 and SM-51, respectively, assembled on these various random copolymer-modified substrates after annealing at 220 °C for 3 h. The thickness of the diblock copolymer films was maintained below 1L0 (or d (nm), listed in Table 2) to eliminate film thickness effects. The film thickness of SM-46 was measured by ellipsometry to be 16.0 nm, which is roughly half of the characteristic spacing L0,H = 36.8 nm. The SM-51 film had a thickness of 15.5 nm, which is ∼0.6 times that of the bulk characteristic lamellar period, L0,L = 28.6 nm. Because these thicknesses are incommensurate with the natural periodicity of the diblocks, the microdomains either orient perpendicularly with respect to the neutral substrate or dewet on preferential substrates. On substrates with 0.57 ≤ FSt ≤ 0.68, SM-46 exhibits perpendicularly oriented M cylinders in an S matrix with a number of defects, specifically, cylinders lying parallel to the surface (Figure 5a). At higher FSt = 0.68−0.73 predominantly parallel cylinders are observed, indicating that the substrate is preferential for the S block. At FSt ∼ 0.78, a “hole and island” morphology is seen due to polymer dewetting from the surface. SM-51 exhibits perpendicularly oriented lamellar microdomains with some defects (lamellae lying parallel to the surface) on substrates with 0.57 ≤ FSt ≤ 0.63 (Figure 5b). As the styrene content of the underlying brush increases, we observe a higher degree of parallel lamellar domain orientation. In comparing the bulk morphology of SM-46 (Figure 4) to that observed in the thin film, we find that surface confinement induced significant ordering of the disorganized H/L* morphology into a perpendicularly oriented, hexagonally packed cylinders morphology. In the case of SM-51, the bulk lamellar G

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morphology appears unperturbed by thin film confinement and thus no significant reconstruction is observed. The thin film domain spacing (L0) in the thin films of SM-46 only deviates by about 10% from that of the bulk diblock, while the thin film lamellar period of SM-51 reflects the same L0 observed in the bulk. Previous work by Han et al. showed that the perpendicular window, defined as the composition of the underlying brush that induces perpendicular domain orientation, for narrow dispersity SM diblock copolymers using hydroxyl-terminated random copolymer-modified substrates is 0.59 ≤ FSt ≤ 0.72 for M cylinders and 0.45 ≤ FSt ≤ 0.57 for lamellae.10 Compared to these monodisperse analogues, the perpendicular window for SM-46 is similar, whereas it occurs at higher FSt for SM-51. Notably, the composition range of the perpendicular window is narrower for SM-51 compared to monodisperse SM diblock copolymers. Sriprom et al. previously observed a shift in the lamellar phase boundary in thin films of polydisperse poly(methyl methacrylate-b-n-butyl acrylate) diblock copolymers toward asymmetric compositions near f PBA = 0.30.45 While they suggested that this composition window shift arises from surface reconstruction due to surface wetting effects, they did not compare the bulk morphological behavior of their copolymers to the thin films. This missing comparison does not rule out the nontrivial role of polydispersity shifting the morphological phase windows. Our studies demonstrate that the composition-dependent phase window shift in thin films arises from the observed shifts in the bulk morphologies. Additionally, our data imply that the random brush compositions associated with neutral surfaces favoring perpendicular microdomain alignment reflect the compositions and not the morphologies of the block copolymer template. Finally, we have shown that surface confinement and wetting effects can significantly influence the ordering of broad dispersity samples such as SM-46, which exhibit poorly ordered microphase-separated morphologies.

system. In summary, these studies clearly demonstrate that broad segmental dispersity does not preclude well-defined, long-range microphase-separated order in block copolymers.

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ASSOCIATED CONTENT

S Supporting Information *

Extended physical characterization data for samples and details of calculations presented in the text. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (M.K.M.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully acknowledge financial support from the UWMadison NSF NSEC (DMR-0832760) for personnel and core facilities. M.K. also acknowledges funding from the Wisconsin Alumni Research Foundation. Synchrotron SAXS data were acquired at the Advanced Photon Source DuPontNorthwestern-Dow Collaborative Access Team beamline (Sector 5), which is supported by E. I. DuPont de Nemours & Co., the Dow Chemical Company, the State of Illinois, and the U.S. Department of Energy, Office of Basic Energy Sciences (Contract DE-AC02-06CH11357).

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REFERENCES

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CONCLUSION In this study, we examined the phase behavior of weakly segregated, polydisperse SM diblock copolymers in the melt and in thin films. In the bulk, SM-44 ( f M = 0.44) exhibited M cylinders in an S matrix, SM-46 ( f M = 0.46) exhibited a disorganized morphology between cylinders and lamellae, and SM-51 ( f M = 0.51) formed lamellae. These observations indicate that M segmental dispersity shifts the cylinders/ lamellae phase boundary to near f M ∼ 0.45−0.50, while also stabilizing the microphase-separated structure at lower than expected segregation strengths (χN). We also found that the domain spacings for the broad dispersity SM diblocks are substantially larger than their monodisperse analogues by up to 70%. Thin films of these SM diblock copolymers on substrates modified with hydroxyl-terminated P(S-r-MMA) brushes exhibited morphologies similar to those observed in bulk samples. Perpendicularly oriented microdomains were observed at mol % styrene brush compositions of 0.59 ≤ FSt ≤ 0.72 for hexagonally packed M cylinders and 0.45 ≤ FSt ≤ 0.57 for lamellae. These results demonstrate that the neutral surface conditions for these block copolymers depend primarily on the block copolymer composition and not the microphaseseparated morphology. Further studies are necessary to determine the precise interplay between polydispersity and brush composition in this polydisperse diblock copolymer H

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