Structure and Properties of Sulfonated Pentablock Terpolymer Films

Article Cite This: Macromolecules XXXX, XXX, XXX−XXX

Structure and Properties of Sulfonated Pentablock Terpolymer Films as a Function of Wet−Dry Cycles Phuc V. Truong,† Stacy Shingleton,§ Mejdi Kammoun,‡ Rephayah L. Black,∥ Marc Charendoff,§ Carl Willis,§ Haleh Ardebili,‡ and Gila E. Stein*,∥ †

Department of Chemical and Biomolecular Engineering and ‡Department of Mechanical Engineering, University of Houston, Houston, Texas 77204, United States § Kraton Performance Polymers, Inc., 16400 Park Row, Houston, Texas 77084, United States ∥ Department of Chemical and Biomolecular Engineering, The University of Tennessee at Knoxville, Knoxville, Tennessee 37996, United States S Supporting Information *

ABSTRACT: The structure and properties of poly(tert-butylstyrene-b-hydrogenated isoprene-b-sulfonated styrene-b-hydrogenated isoprene-b-tert-butylstyrene) (tBS-HI-SS-HI-tBS) films were investigated as a function of “wet−dry cycles”, where one “cycle” is defined as a 24 h soak in deionized water followed by a 24 h drying period in air. Films were characterized with a variety of complementary measurements that include X-ray scattering, infrared spectroscopy, water uptake, impedance spectroscopy, and tensile tests. We find that cycling drives a structural transition toward increasingly interconnected SS domains, which is favorable for water and ion transport. However, cycling can also induce mechanical deformations that reduce ductility, swelling, and water uptake. The significance of this trade-off is illustrated by comparing the properties for two film thicknesses as a function of cycle number: The ductility of thinner films (15 μm) is lost after four cycles, an effect that is correlated with the appearance of macroscale buckles, and the extent of swelling is also reduced. Therefore, the transport properties reflect a balance between the increased SS domain interactions and reduced water content. The ductility in thick films (30 μm) also declines with cycling, but to a lesser extent, and these systems retain their ability to swell through six cycles. Therefore, the transition to a network-like SS structure enhances both water uptake and transport. These systematic studies demonstrate that successive wet−dry cycles can lead to complex changes in the performance of amphiphilic block copolymer films, which may complicate their design for applications in water treatment or protonconducting layers in electrochemical devices. desalination membranes.16,17 However, at the high sulfonation levels needed for viable salt sorption and water transport behavior, they become water-soluble and lose their mechanical stability. The end-segments of poly(styrene-b-ethylene butadiene-b-styrene) (SEBS) triblock copolymers can be sulfonated to a high level without losing mechanical strength in a dry state or when hydrated by water vapor.18 However, the mechanical properties decline when hydrated in liquid water due to the uncontrolled swelling of the sulfonated styrene end-blocks. For applications such as direct methanol fuel cells, sulfonated poly(styrene-b-isobutylene-b-styrene) (SIBS) triblock copolymers offer good conductivities and do not dissolve in water at high ion contents. However, they suffer from extreme liquid water uptakes (ca. 350 wt %) and high liquid methanol permeability.19−21

1. INTRODUCTION Ionic block copolymers show promise for applications in membranes and batteries because they can transport water and ions with good mechanical stability in both dry and hydrated states.1−3 The morphology of these materials can be tuned by altering the volume fraction, molecular weight, and chemistry of each block, providing a means to optimize performance for different applications. Potential applications for ionic block copolymers include gas purification,1,2 reverse osmosis membranes for water purification,3−5 polyelectrolyte membrane for fuel cells,6,7 lithium ion batteries,8,9 and supercapacitors.10−12 Sulfonated polymers are a widely explored class of ionic polymers that show promise in the aforementioned applications.3,13−15 However, it is challenging to optimize their transport properties without compromising the mechanical properties. As examples, sulfonated poly(arylene ether sulfones) have been studied in liquid water as potential © XXXX American Chemical Society

Received: January 25, 2018

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DOI: 10.1021/acs.macromol.8b00194 Macromolecules XXXX, XXX, XXX−XXX

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in the structure when exposed to water. However, swelling and deswelling can drive nanoscale restructuring in block copolymer films31−33 and induce surface deformations in other types of polymer films.34,35 Such behaviors, if present, would likely impact both transport and mechanical properties. The effects of “cyclic” environmental conditions on polymer electrolyte membranes are also relevant to energy conversion and storage. As an example, humidity-induced swelling and deswelling can lead to mechanical degradation of Nafion in fuel cells.35 In this work, we show that “cycles” of swelling in liquid water and drying in humid air can progressively change the structure and properties of tBS-HI-SS-HI-tBS pentablock terpolymers. Cycling increases the interactions between SS domains but also degrades the mechanical properties. These behaviors lead to complex structure−property−processing relations: an SS network should facilitate water uptake and transport, but extensive strain will limit swelling and water uptake. The competition between these attributes is evident when comparing the performance of different film thicknesses, as thinner films are more susceptible to mechanical instabilities. These studies highlight the challenges in designing amphiphilic block copolymers for water treatment and electrochemical devices, as a changing environment can affect structure, water uptake, transport, and mechanics in different ways.

A promising type of sulfonated material is the pentablock terpolymer poly(tert-butylstyrene-b-hydrogenated isoprene-bsulfonated styrene-b-hydrogenated isoprene-b-tert-butylstyrene) (tBS-HI-SS-HI-tBS). The material can be dissolved in an organic solvent and processed into films using low cost roll-toroll techniques such as slot-die coating. This block copolymer system can resist extreme swelling and maintain its mechanical integrity in liquid water due to the glassy tBS end-blocks.18 The rubbery HI segments impart toughness to the dry films and can deform to facilitate water uptake. Transport of ions, water, and polar gases is mediated by the SS middle block. Numerous studies have demonstrated that tBS-HI-SS-HI-tBS copolymers can offer high proton conductivity and water permeability, good mechanical stability under hydrated conditions, and high water uptake with swelling resistance.1,18,22−24 In addition, film morphology and transport properties can be tuned through sulfonation level,22,23,25,26 molecular weight,4 and casting conditions.4,23,27 Recent studies have examined the structure of this sulfonated pentablock copolymer in solutions and thick films (10−100 μm) as well as important film properties such as ion transport, water transport, and mechanics. When dissolved in a cyclohexane/heptane mixture, which is selective to the tBS endsegments, this block copolymer assembles into spherical micelles having an SS core and HI/tBS corona.25 This micellar structure is observed at sulfonation levels of 26 and 52 mol %, producing ion exchange capacities (IEC) of 1.0 and 2.0 mequiv/g, respectively. The as-cast film structure has been measured with transmission electron microscopy (TEM), revealing discrete SS microdomains for IEC of 1.0 mequiv/g and a near-percolating network of SS microdomains at an IEC of 2.0 mequiv/g.26,27 The high IEC and dense packing of SS cores lead to large increases in water sorption, salt permeability, and water/proton transport.4,22,26 Transport properties of asprepared films are also sensitive to the casting methods, likely due to the formation of different kinetically trapped morphologies.4 Furthermore, the mechanical properties of asprepared films in humid air and water have been examined. Tensile tests show that films measured in humid air exhibit yielding, which is absent in water, and are stronger and stiffer than wet films.24 Biaxial creep tests suggest the hydrated deformation behavior of these films are strongly dependent on the IEC.23 The total strain increases by a factor of 2 as the IEC is increased from 1.0 to 1.5 mequiv/g and by a factor of 3 as the IEC is increased from 1.0 to 2.0 mequiv/g. To address the critical worldwide shortage of clean water, membrane-based technologies such as reverse osmosis are among the most effective and economic ways available to purify brackish groundwater or seawater. Currently, the dominant membrane materials are aromatic polyamide-based thin film composite membranes.28,29 However, these membranes must be prewetted by manufacturers and shipped in a hydrated state to end-users. If the membrane dries out, then it is difficult to rewet the hydrophobic polysulfone support layer, causing a large decline in water flux.30 Membranes based on tBS-HI-SSHI-tBS films show promise in desalination:18 the material is easily wet and rewet, which addresses one current challenge in water purification, while still maintaining high water permeability and good ion selectivity. To our knowledge, the effects of wetting, drying, and rewetting on copolymer structure and performance have not been studied. In the case of tBS-HI-SSHI-tBS pentablock terpolymers, one might expect that the glassy tBS end-segments will act as physical cross-links and lock

2. EXPERIMENTAL METHODS 2.1. Material. The sulfonated pentablock copolymer (Figure 1) was synthesized and characterized following previously reported

Figure 1. Structure of sulfonated pentablock copolymer. The blocks tBS, HI, and SS are labeled. The films are in the “acid” form and have no counterion. procedures.36 A pentablock copolymer of tBS-I-S-I-tBS, where tBS = poly(t-butylstyrene), I = poly(isoprene), and S = polystyrene, was synthesized via anionic polymerization followed by hydrogenation of the polyisoprene block (HI). The molecular weight of the unsulfonated pentablock copolymer is approximately 15−10−28− 10−15 kg/mol. The middle polystyrene block of the pentablock copolymer was then selectively sulfonated (SS) to an ion exchange capacity (IEC) of 2.0 mequiv/g (52 mol % sulfonation). Sulfonated pentablock copolymer films with different thicknesses h (15 and 30 μm at 0% relative humidity) were received from Kraton Polymers LLC. These films were machine-cast onto Mylar sheets from a solution of 11 wt % polymer in a cyclohexane and heptane mixture (72:28 by weight) at ambient temperature and humidity. Film thickness was controlled by adjusting the knife height and angle at the coating head. Films were then immediately transferred to an oven that has increasing temperature zones from 66 to 93 °C for 1 h. The relative humidity (RH) was not controlled inside the oven but was likely low at these temperatures. The resulting films were stored inside a glovebox at 25 °C and 0% RH prior to use. 2.2. Sample Preparation. Each wet−dry cycle involves fully soaking film samples at ambient conditions in deionized (DI) water (18.2 MΩ) for 24 h. Samples were spread flat on an absorbent cloth (clean room wipe) and dried in ambient air at 25 °C and ∼50−55% B


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resolution of 8 cm−1 with an average of 128 scans. A “point-and-click” procedure in MATLAB was used to subtract the baseline. This software allows the user to identify the baseline points in the spectra and connect neighboring points with a line. Areas of IR absorbance peaks were determined with numerical integration in Igor Pro. Three independent film samples were measured for each cycle and thickness, and the average values with error bars (±1 standard deviation) are reported. 2.5. Water Vapor Transfer Rate (WVTR). The water vapor transfer rate was measured by the upright and inverted cup methods at 25 °C and 50% RH in a humidity-controlled chamber (Testequity model 123H). The vials were filled with 10 mL of DI water. Film samples were tightly fastened inside the vial caps to create a permeable seal. Water was not in direct contact with the film in the upright cup method, but the vial was turned upside-down in the inverted cup method so water directly contacted the films. A small fan was installed inside the humidity chamber to facilitate removal of possible water accumulation at the film/air interface during the test. The vials were weighed four times every 2 h to determine the weight of water that evaporated. Three independent film samples were measured for each cycle and thickness, and the average and error bar (±1 standard deviation) are reported. Water vapor transfer rate (kg/day/m2) for each sample was determined using linear regression from the change in weight as a function of time. The coefficient of determination R2 value ranges from 0.97 to 0.99. Note that in the inverted test the WVTR calculation assumes that a constant surface area is available for mass transfer of water. During actual experiments, this area slightly expands as water pushes down on the membrane and causes a small bulge. A schematic of the experimental geometry is included in Figure S1. 2.6. Film Thickness Measurements. Each film was sandwiched between two steel disks of known thickness, and the thickness of the full stack was measured with a digital caliper that has an accuracy of 2 μm. The caliper was tightened until resistance was felt. “Dry” film thicknesses were measured at 25 °C and 0% RH inside a nitrogenpurged glovebox. “Humid” film thicknesses were measured at 25 °C and 55% RH after the films were equilibrated at the same conditions for at least 30 min. Through thickness measurements, we verified that these films swell to their equilibrium thickness within minutes after being transferred from glovebox (0% RH) to laboratory ambient conditions (∼55% RH). 2.7. Proton Conductivity. Films were transferred from a nitrogen-purged glovebox and equilibrated in air at 25 °C and 55% RH (sensor indicated 55% RH during testing with small fluctuations of ±2%) for 30 min. Proton conductivity measurements were then performed at the same ambient conditions by means of impedance spectroscopy under potentiostatic mode using a Metrohm Autolab Frequency Response Analysis (FRA 2). Three films were measured for each cycle and thickness. Each film was cut into a disk-like shape (area A = 2.84 cm2) and sandwiched between two stainless steel disks that have the same area, corresponding to a through-plane conductivity measurement. This configuration provides proper contact between the film sample and the electrodes, eliminating noise in the measurements. The swollen film thickness was measured as described above. Nyquist impedance plots were collected over the frequency range of 10 Hz to 1 MHz at 10 mV of ac amplitude. Data from Nyquist plots were fit using the electrical circuits (Figures S2 and S3) to determine the bulk film resistance Rb. Proton conductivity (S/cm) was then calculated using eq 3:

RH for another 24 h. The cycle number indicates how many times this process was repeated. Film samples were transferred to a fresh volume of DI water before each new cycle. At cycle number 0, there was no soaking and films were measured as provided. Note that the pH of DI water was approximately 5.5, which was measured with pH paper. The slight acidity is from dissolved carbon dioxide that forms carbonic acid. 2.3. Small-Angle X-ray Scattering (SAXS) and Wide-Angle Xray Scattering (WAXS). Small-angle X-ray scattering experiments were performed using a Rigaku S-MAX-3000. Cu Kα X-rays (0.154 nm) were generated from a RA-Micro7 HFM tabletop rotating anode generator operating at 40 kV and 30 mA. The point-collimated beam was obtained using Rigaku Confocal Max-Flux (CMF) optics. Film samples were mounted on an aluminum sample holder and illuminated at normal incidence (transmission through the film thickness) under vacuum. SAXS data were recorded using a two-dimensional single photon counting detector with a pixel size of 100 μm at a sample-todetector distance of 1 m. The detection q-range with this setup is 0.08−4.0 nm−1. Three independent film samples were measured for each cycle and thickness. Each sample was measured for 10 min, and then the sample was removed and background data were recorded for 10 min. 2D spectra were converted into 1D plots with azimuthal angle integration in SAXSGUI software, producing the scattering intensity profile I(q). The 1D background profile was subtracted from 1D film data set. Wide-angle X-ray scattering experiments were performed with the same beamline. WAXS data were recorded using an image plate detector (18 cm × 24 cm) with a spatial resolution of 10 pixels/mm at a sample-to-detector distance of 70 mm. The detection range with this setup is 2.5−40 nm−1. One film sample was measured for each cycle. Each sample was measured for 12 h, and then the sample was removed and background data were recorded for 12 h. 2D spectra were converted into 1D plots as described above. These data are presented in the Supporting Information. 2.4. Chemical Functionality. 2.4.1. Water Content. The water vapor uptake of the films was calculated by eq 1, where W55RH and W0RH are the weights of the film at ∼55% RH (in humid air at ambient laboratory conditions) and 0% RH (inside glovebox), respectively, in grams. Both W55RH and W0RH were measured at room temperature (25 °C). Three independent film samples were measured for each cycle, and the average and error bar (±1 standard deviation) are reported. water content =

W55RH − W0RH × 100 W55RH

(1)

2.4.2. Ion Exchange Capacity. The ion exchange capacity (IEC; mequiv/g) of the films was measured by titration of the polymer following two approaches. In the first method, 0.1 g of the film was dissolved in 50 mL of tetrahydrofuran (THF). THF does not participate in reactions and serves only as a medium.37 In the second method, the films were cut into small pieces (approximate mass of 0.1 g) and immersed in 50 mL of 1.0 N NaCl solution at room temperature for 2 days. During this time, the H+ ions in the film were replaced by the Na+ ions.38,39 After completion of ion exchange, films were removed from the solution. In both methods, the solutions were titrated with a standard 0.1 N NaOH using phenolphthalein as an indicator. The IEC (mequiv/g) was calculated from the titration data using eq 2: IEC =

VNaOHC NaOH W0RH

(2)

σ=

where VNaOH is the volume of NaOH consumed (mL), CNaOH is the concentration of NaOH (N), and W0RH is the weight of the film at 0% RH (inside glovebox). Six independent film samples were measured for each cycle, and the average and error bar (±1 standard deviation) are reported. 2.4.3. Attenuated Total Reflectance (ATR). Each film sample was secured on an ATR crystal (diamond ZnSe). The penetration depth of the infrared (IR) beam into the sample is approximately 2 μm. IR absorbance was measured under laboratory ambient conditions (∼52% RH) with a Nicolet NEXUS 670 FTIR spectrometer and recorded at a

h R bA

(3)

Error bars for conductivity were calculated from a propagation of errors analysis that included the uncertainty in measuring Rb and h. 2.8. Tensile Property Characterization. Tensile mechanical properties were measured according to ASTM D 412.40 Samples were cut from polymer films using a die and a press. Each dumbbell sample was 1 in. long and 0.1 in. wide. Both “humid” and “wet” tensile tests were performed using a dual column, tabletop Instron 3366 instrument. Tensile tests of “humid” samples were performed in an C


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Macromolecules environmentally controlled laboratory at 25 °C and 50% RH. “Wet” samples were prepared by equilibrating films in DI water 24 h before measurement, and “wet” measurements were implemented in a custom-built chamber. The wet chamber was filled with DI water such that samples were fully submerged, and samples were soaked for 5 min before initiating the test. Samples were mounted between the tester grips under a load capacity of 10 N. The initial separation between the test grips was 1 in. The crosshead speed was 2 in./min, and the strain rate was fixed at 0.0167 s−1. Data acquisition and analysis were performed automatically using Bluehill 3 software. Seven independent film samples were tested for each cycle and thickness. Because of possible imperfections, some samples failed prematurely during both “wet” and “humid” testing, as illustrated by data in Figure S5. Therefore, results of the three best-performing samples for each cycle are reported in the discussion. 2.9. Optical Microscopy. Film topography was imaged using a Nikon Eclipse LV100 upright microscope at ambient laboratory conditions (∼52% RH). Films were placed on a reflective silicon substrate and imaged in bright-field mode.

scattering spectra exhibit similar features: a strong and broad primary peak that reflects the types of interactions (or correlations) between domains, followed by a weak “shoulder” that is determined by the sizes and shapes of SS cores. SAXS is a common tool for characterization of disordered media, but interpretation of these data is very challenging because many different models can offer a quantitative fit. However, using hypotheses about the film structure at each cycle that is partly based on microscopy work from other reports,27,42 we selected a model that has few adjustable parameters and can quantitatively describe all data sets over the full range of q. First, the structure of this sulfonated pentablock copolymer has been studied in solution and in as-cast films. When dissolved in a nonpolar solvent, such as the cyclohexane/heptane mixture used for commercial machine casting, the block copolymer assembles into micelles having an SS core and HI/tBS corona.41 This micelle structure is partly preserved in the ascast films: transmission electron microscopy (TEM) reveals a dense packing of disordered micelles that are near percolation.27,42 Second, immersion in water will swell the SS domains42 and may induce connectivity between SS cores. The network may be “locked in” as water evaporates because the tBS end-segments will rapidly vitrify and act as physical cross-links. Such swelling-induced transitions from isolated domains to network-like phases have been reported for diblock copolymers when a glassy insoluble block is linked to a soluble block.31,33,45 The scattering from spherical block polymer micelles is welldescribed by the analytical structure factor for hard-sphere interactions S(q), which captures the intermicelle length scales and packing density, and a spherical form factor P(q) that accounts for the size of micelle cores (SS domains).46−48 The model parameters are the radius of the micelle core Rc, the hard-sphere interaction radius Rhs, and the micelle volume fraction N. The length scales Rc and Rhs are illustrated in Figure S6. We incorporate polydispersity in Rc using a Gaussian probability density with standard deviation σ. Polydispersity is implemented with the local monodisperse approximation (LMA), which assumes that micelle size and position are coupled.46 We employ a simple proportional relation, Rhs = kRc with k ≥ 1. When k = 1, the micelle cores are touching and the model predicts a network-like structure. The SAXS intensity for this hard-sphere model with polydispersity in core size is

3. RESULTS AND DISCUSSION The objective of these studies is to examine the structure and properties of the sulfonated pentablock copolymer tBS-HI-SSHI-tBS after a series of swelling and deswelling steps in liquid DI water. Two sulfonated pentablock copolymer films with different dry thicknesses h (15 and 30 μm at 0% RH) were investigated. All films were machine cast from a nonpolar solvent, so the as-prepared structure resembles densely packed micelles with an SS core and HI/tBA corona.27,41,42 In the following discussion, a “cycle” is defined as one 24 h soak in DI water (swelling) followed by 24 h of drying in air at 25 °C and 30% RH (deswelling). The films are “free” when soaked in DI water, meaning they are not tethered or clamped to any surface. Cycle 0 designates the as-prepared film, which is stored inside a nitrogen-purged glovebox prior to use. 3.1. Nanoscale Structure. The film structure was analyzed as a function of cycle number and film thickness using SAXS. The samples were stored in a nitrogen-purged glovebox prior to the SAXS measurement, and SAXS data were acquired under vacuum, so there is little or no moisture in the SS domains during these measurements. The contrast for SAXS measurements is due to spatial variations in scattering length density (SLD). The SLD of each block is calculated for 0.154 nm photons from the chemical formula and mass density (ρ), and the results are reported in Table 1. The tBS/HI contrast is Table 1. Stoichiometry, Mass Density (ρ),43 and Scattering Length Density (SLD)44 for Each Block polymer

formula

ρ (g/cm3)

SLD (10−4 nm−2)

tBS HI SS

C12H16 C5H10 C8H8/C8H7SO3

0.95 0.85 1.26

8.85 + 0.01i 8.25 + 0.01i 11.37 + 0.01i

I(q) = A 0

⎧ (R − R )2 ⎫ c ⎬S(q , R hs = kR c , N )P(q , R ) dR 2σ 2 ⎭

∫ σ√12π exp⎨⎩

(4)

The five adjustable parameters for model refinement are the standard deviation σ (nm), the mean core radius Rc (nm), the proportionality constant k ≥ 1, the volume fraction of micelles N, and the prefactor A0 (arbitrary units). The structure factor and form factor functions are included in the Supporting Information along with details that pertain to the fitting algorithm and experimental resolution function. Other SAXS models were also considered and are discussed in the Supporting Information. Figure 2 and Figure S16 show representative scattering data (open symbols) for 15 and 30 μm films at each cycle. The model of eq 4 can describe the scattering from 15 μm films with near-quantitative agreement. However, the scattering from 30 μm films can only be fit with a “two-layer” model, meaning eq 4 is applied with two different sets of parameters, and we cannot determine unique values for all 10 unknown parameters.

much weaker than the tBS/SS or HI/SS contrast, and attempts to fit a three-phase core−shell−matrix (SS-HI-tBS) model were not informative, as these merely converged to a two-phase core−matrix solution. Therefore, the models described herein assume a two-phase system of spherical SS domains surrounded by an HI/tBS matrix. All SAXS measurements were performed with the incident beam perpendicular to the film surface, which is most sensitive to structures within the plane of the film. Representative SAXS data of 15 and 30 μm films for each cycle are summarized in Figure 2, where I is the intensity in arbitrary units and q is the scattering vector in nm−1. All D


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Figure 2. Representative SAXS data on log I vs log q scale (a) 15 μm films and (b) 30 μm films. Lines are best fits to the hard-sphere model described in the text and Supporting Information.

Figure 3. Optimized parameters for hard-sphere model as a function of cycle number for 15 and 30 μm films. (a) Core radius Rc, (b) proportionality k = Rhs/Rc, (c) volume fraction of cores ϕc, and (d) σ, standard deviation of Rc. The optimized parameters of 15 μm films are marked by red triangles. The 30 μm films are modeled as two equal layers, with a constrained layer (red triangles) and an optimized layer (blue squares). Symbols represent the average value from three independent measurements. Error bars represent the statistically derived uncertainty from nonlinear regression or the sample-to-sample variability (whichever is larger).

fraction of SS cores, ϕc = NRc3/Rhs3. The parameters demonstrate a progressive structural transition toward a more disordered, network-like state with increasing cycle number. The model describes a percolating network when k ≈ 1 and ϕc > 0.34,49 which occurs near cycle 3 in dehydrated 15 μm films. It is difficult to say if or when a network is formed throughout thicker dehydrated films, but these data predict a similar trend toward greater disorder with increasing cycle number. Scattering from disordered, two-phase systems is often analyzed with the Teubner−Strey (TS) structure factor.50,51 This model describes a bicontinuous structure and is restricted to low-q data. The TS model parameters are a characteristic periodicity d and a correlation length ξ. The optimized values of these parameters are summarized in Table S2 as a function of cycle number. These data also support a progressive change in structure, specifically a transition to a more disordered,

Therefore, one “layer” is constrained based on the best fit from 15 μm films at the same cycle number, and the parameters that describe the second “layer” are optimized. We assume each layer equally contributes to the observed scattering. This description agrees with the data, and the trend is qualitatively correct, but the numbers should be interpreted with caution. While the data for each film thickness show similar characteristics, the structures in 15 and 30 μm films are different: 15 μm films are characterized by smaller length scales than 30 μm films, which is seen by visual inspection of the structure factor peak in Figure 2 (marked by solid vertical line). The films are prepared from the same solution on the same casting line, but using different blade heights, so this outcome is likely a consequence of different drying rates. Figure 3a−d summarizes the optimized SAXS model parameters for each cycle along with the calculated volume E


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reports water uptake under “wet” conditions, meaning after submersion in water. These data show that, on average, 30 μm films take up approximately 5−10% more liquid water than 15 μm films, but the results are independent of cycle number. ATR measurements at 55% RH confirm the water uptake data and also show the effects of water content on sulfonic acid dissociation. Figure 5a,b reports the IR absorbance as a function

network-like state with increasing cycle number and decreasing film thickness. As discussed in subsequent sections of this article, the transition to a state of increased disorder is correlated with enhanced rates of transport. 3.2. Chemical Functionality. IEC and water uptake are important characteristics of ion exchange membranes. The IEC results obtained from two titration methods are plotted in Figure S7 as a function of cycle number. In the first method, the polymer is dissolved in THF so all acid sites are accessible and directly detected through titration. This measurement determines an “intrinsic” IEC that is set by the sulfonation level in the polymer. In the second method, polymer films are soaked in a sodium chloride (NaCl) solution for 2 days, and titration detects the products of that ion exchange reaction. Both methods show similar trends as a function of cycle number for 15 and 30 μm films: the IEC decreases from 2.2 mequiv/g at cycle 0 to 2.0 mequiv/g at cycle 1 and then stays constant at 2.0 mequiv/g from cycle 1 to cycle 6. The drop in IEC after the first soak is likely due to the removal of byproducts, such as residual isobutyric acid, that remain in the polymer from the sulfonation step. The IEC value of ca. 2.0 mequiv/g is consistent with other reports where the films were soaked to remove byproducts before titration.4,18,42 This outcome demonstrates that there is no loss in the functionality of the polymer as a result of repeated soaking and drying. Figure 4 shows water content as a function of cycle number at 55% RH. Opposite trends are observed for each film

Figure 5. IR spectra of membranes for (a) 15 μm films and (b) 30 μm films. All spectra are normalized to the polystyrene aromatic absorbance at 828 cm−1 in 15 μm films. Area under IR absorbance peak for (c) water band at 1630 cm−1 as a function of cycle number. Three independent films were measured per cycle number, and the average value and error bars (±1 standard deviation) are reported.

Figure 4. Water content at 55% RH for 15 and 30 μm films. Three independent film samples were measured for each cycle. Each data point represents the average value, and error bars are ±1 standard deviation.

thickness: For the 15 μm films, the water content of the as-cast films is approximately 14 wt % (cycle 0) but drops to 9 wt % at cycle 1 and gradually decreases to 6 wt % by cycle 4. For the 30 μm films, the water content of the as-cast films is 19 wt % and gradually increases up to 30 wt % by cycle 4. No changes are observed in either film thickness after cycle 4. We also find that thicker films initially take up more water than thinner films. This may reflect a surface “skin” layer with different properties than the bulk, or it could be associated with the larger domain sizes in thicker films (Figures 2 and 3). Figure S11 reports the swollen thickness of each film at 55% RH. The trend mirrors that of water uptake: with each successive cycle, thin films swell less and thick films swell more. Figure S8 reports the hydration number λ at 55% RH (number of water molecules/sulfonic acid) for both film thicknesses as a function of cycle number. The results show the same trend as water content: λ dropped from approximately 3.5 to 1.7 over 4 cycles in 15 μm films and increased from 4.0 to 8.5 over 4 cycles in 30 μm films; λ stabilizes after 4 cycles in both film thicknesses. Figure S9

Table 2. FTIR Peak Assignments peak

cm−1

peak assignment

i ii iii iv v

828 1010 1130 1440 1630

PS aromatic ring SS aromatic ring R−SO3− undissociated HSO3 OH bending band of H2O

of cycle number, and Table 2 summarizes the peak assignments in the range of 800−1800 cm−1. All spectra are normalized to the polystyrene aromatic absorbance at 828 cm−1 (peak i) in 15 μm films.52 The peaks at 1010 cm−1 (peak ii) and 1130 cm−1 (peak iii) are attributed to the CH stretching/bending mode for SS aromatic ring and the asymmetric stretching mode for (SO3−) concentration, respectively.51 Previous studies of the same material showed that these peaks are sensitive to IEC and F


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Macromolecules type of cation.22,53 Based on our IEC data (Figure S7), no changes are anticipated in these signals as a function of cycle number, and no changes were observed (Figure S10). Figure 5c shows that the water band at 1630 cm−1 (peak v) mirrors the trend in water uptake for each film thickness, as it decreases with cycle number for 15 μm films and increases with cycle number for 30 μm films. In previous studies of Nafion, the absorbance at 1440 cm−1 (peak iv) was attributed to heatactivated cross-linking through condensation of sulfonic acids54−56 or undissociated sulfonic acid in dehydrated samples.57 Cross-linking is not anticipated with cycling at ambient temperature, and titration of the dissolved polymer confirms a constant level of sulfonic acid in all samples. As such, we attribute the peak at 1440 cm−1 to undissociated sulfonic acid. An increase in water content should correspond to a lower amount of undissociated acid (more acid dissociation), and vice versa. Figure S10c shows the ratio of the area under the 1440 cm−1 peak and the area under the 1630 cm−1 peak for both film thicknesses. As expected, this ratio increases with increasing cycle number for 15 μm films, while the opposite trend is observed for the 30 μm films. This is consistent with water uptake reported in Figure 4. There are two significant outcomes of the IEC, water uptake, and IR measurements: First, we find that sulfonated styrene functionality is preserved through the cycling protocol, so changes in water uptake are not associated with chemical degradation. Second, cycling reduces water uptake in thinner films but enhances water uptake in thicker films, with a corresponding trend in the extent of acid dissociation. The decline in water uptake for thin films seems inconsistent with the SAXS data, as a transition from isolated SS domains to a network-like structure should produce the opposite effect. However, as discussed in later sections, these films are susceptible to mechanical instabilities that increase in severity with decreasing thickness and successive cycles. The films are strained as a result of these instabilities, which can reduce their ability to swell and take up water. 3.3. Transport Properties. Water vapor transport rates (WVTR) and proton conductivity were evaluated as a function of cycle number through evaporative mass loss and impedance spectroscopy, respectively. Both of these techniques probe transport through the film thickness. Measurements of WVTR at 50% RH are summarized in Figure 6a (inverted) and Figure 6b (upright). Using the inverted cup method, where liquid water is in direct contact with the film, the measured fluxes for as-cast (cycle 0) 15 and 30 μm films are in the range of 20−30 kg/m2/day. These values are consistent with a previous study that used that same method to measure WVTR from films of the same polymer.42 Beyond cycle 0, there is no clear trend in the measured fluxes of the 15 μm films, as the results fluctuate between 20 and 30 kg/m2/day. On the other hand, the thicker films exhibit a gradual increase in WVTR (30−34 kg/m2/day from cycle 0 to cycle 6). We note that the area of the film is assumed constant when calculating the flux, but there is some bulging that may vary with film thickness and cycle number. The upright cup method, where liquid water is not in direct contact with the membrane, yields similar results (∼0.5 kg/m2/ day) for both films at cycle 0. However, there is dramatic change in WVTR for both film thicknesses after the first soak in water: the fluxes increase from approximately 0.4 kg/m2/day at cycle 0 to 1.2 kg/m2/day at cycle 1, which demonstrates that soaking and drying the as-prepared films can open a transport pathway. Subsequent soak/dry cycles produce a slight decrease

Figure 6. Water vapor transport rates of 15 and 30 μm films measured as a function of cycle number from the (a) inverted geometry and (b) upright geometry. Ambient was controlled at 50% RH. Each data point reflects the average value and uncertainty (±1 standard deviation) from three independent measurements.

in WVTR from 1.2 to 1.0 kg/m2/day for 15 μm films and a slight increase in WVTR from 1.2 to 1.4 kg/m2/day for 30 μm films. The nature of the “transport pathway” induced by the first cycle could reflect changes in surface structure, bulk structure, or both. Soaking in water could drive the segregation of SS blocks to the membrane surface, thereby reducing the barrier for water diffusion into the film. The membrane surfaces are too rough to analyze with contact angle measurements, but all samples (as-cast and cycled) resist immersion in water, which means all films have some tBS/HI segments at the surface. The SAXS measurements of dehydrated films show a gradual restructuring process, rather than a transition to a network after a single cycle. However, it is worth noting that a subtle structural change, either at the surface or in the bulk, could be amplified through swelling. As an example, an atomic force microscopy (AFM) study of a similar material was able to image a humidity-induced transition from isolated SS micelles to a continuous SS network.23 Furthermore, tensile tests (section 3.4) also show a large change in mechanical properties at 50% RH after a single cycle, and these measurements are more sensitive to bulk structure than surface properties. The trends in WVTR are consistent with water uptake data. As a first approximation, the measured water flux in a solutiondiffusion framework is proportional to permeability and inversely proportional to film thickness.58 Therefore, if each film exhibits the same permeability, then the flux for 30 μm films should be half that of 15 μm films. Instead, using both upright and inverted measurements, we find similar or enhanced transport rates in the thicker films that take up more water (Figure 4 and Figure S9). Similarly, the trends in WVTR as a function of cycle number generally match the trends in water uptake: the exception is the dramatic increase in “upright” measurements after a single cycle, which reflects a structural change. Measurements of bulk proton conductivity σ at ∼55% RH (ambient laboratory conditions) are summarized in Figure 7a,b as a function of cycle number and water content. The measured conductivities of 15 μm films range from 0.002 to 0.003 mS/cm G


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with no clear dependence on cycle number or water content, at least within the uncertainty of the measurements. In the thicker films, where water content increases by a factor of 1.6 (19 to 30 wt %) from cycle 0 to cycle 6, the ionic conductivity increases by a factor of 2.6 (0.006 to 0.016 mS/cm). This behavior is expected as water facilitates dissociation and transport of ions.19,59,60 To distinguish between the effects of water content and nanoscale structure on transport, Figure 7c shows the ionic conductivity σ normalized by the water mass fraction xw. These data illustrate that cycling enhances ion transport in both 15 and 30 μm films, consistent with the progressive structural transition to a more network-like structure. The normalized conductivity is approximately 40% higher in 30 μm films compared with 15 μm films. The effects of film thickness on normalized conductivity can arise from morphological effects: as films become thicker, bulk properties will dominate surface properties, so water uptake and ion conduction are likely to overshadow roughness and contact impedance at the interfaces. However, another potential source of the discrepancy is straindependent ion conductivity. As discussed in section 3.4, the thinner films are more susceptible to cycling-induced mechanical stabilities that introduce buckles. The internal compressive stresses that drive extensive buckling in thinner films could also depress the conductivity.61 3.4. Mechanical Properties. Mechanical properties were evaluated with tensile tests. Membranes were measured in humid air (50% RH) and immersed in liquid water. Figure 8a,b shows representative stress−strain curves as a function of cycle number. The data collected for each sample using the stress− strain curve include elongation at break and stress at break, taken as the strain and stress when the film snaps, respectively; tensile modulus, which is the initial slope of the curve; and yield stress, the point at which deformation becomes permanent and irreversible. We observe different mechanical responses under each measurement environment, which are consistent with other reports:24 As examples, only the films measured in humid air exhibit yielding, while immersion in water significantly

Figure 7. Proton conductivity σ at 55% RH as a function of water content for (a) 15 μm films and (b) 30 μm films. Each data point reflects a single measurement, and error bars reflect the uncertainty of conductivity calculations based on propagation of errors. (c) Proton conductivity at 55% RH normalized by mass fraction of water (xw) as a function of cycle number. Each data point reflects the average value, and error bars reflect the uncertainty (±1 standard deviation) from three independent measurements.

Figure 8. Representative tensile tests of 15 μm samples measured (a) at 50% RH in air (“humid”) and (b) submerged in liquid DI water (“wet”). Representative tensile tests of 30 μm samples measured (c) at 50% RH in air (“humid”) and (d) submerged in liquid DI water (“wet”). H


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Figure 9. (a) Elongation at break, (b) stress at break, (c) Young’s modulus, and (d) yield stress of 15 and 30 μm films at 50% RH (“humid”). The average value and error bars (±1 standard deviation) are calculated from three independent measurements. Note that the 15 μm films become too brittle for tensile tests after cycle 4.

further cycling, the modulus of each film steadily decreases and approaches 380 MPa (15 μm films) and 630 MPa (30 μm films). The stress at break for as-cast 15 μm films is approximately 7 MPa, then increases, and holds near 12 MPa after cycle 1. The thicker films also experience an increase in stress at break from cycle 0 to 1 (10 to 17 MPa), which stays relatively constant at 15 MPa after cycle 1. The only property that is not increased by the first soak/dry cycle is ductility (percent elongation), which steadily declines with increasing cycle number: in 15 μm films, ductility declines from 40% to 10% over 4 cycles, while 30 μm films see a decline from 160% to 80% over 6 cycles. The extreme loss of ductility in 15 μm films is also seen by visual observation, as the films become increasingly delicate (brittle) and difficult to handle. The initial increases in modulus and strength at cycle 1 are consistent with greater interactions between sulfonated domains, as suggested by SAXS data analysis and WVTR measurements. It is worth noting that the nanoscale structure is partly controlled by processing, so a change in film thickness will result in different domain sizes/length scales, as shown in Figures 2 and 3, and this may account for the some differences in the performance of 15 and 30 μm films. Figure 10a−c summarizes the outcomes of tensile tests for samples that were fully submerged in DI water (“wet”). The liquid water content is approximately the same for all film thicknesses and cycle numbers, as summarized in Figure S9. Both 15 and 30 μm films are softer and more elastic in liquid water. The modulus of 15 μm films is approximately 20 MPa as-cast and declines to 13 MPa with cycling, while the modulus of 30 μm films is 39 MPa as-cast and declines to 25 MPa with cycling. Elongation at break and stress at break are similar for both film thicknesses from cycle 0 to 3; by cycle 4, the thinner films start failing, as in the case of humid measurements. The tensile tests under humid and wet conditions demonstrate that cycling has severe consequences on mechanics, particularly in 15 μm films, that are not explained by changes in the nanoscale structure or water content. Therefore, we consider other changes that might lead to mechanical failures and cannot be detected through scattering or IR. As examples: (1) If the SS domains contract during the

enhances the ductility. Both measurement environments reveal changes in film mechanics with increasing cycle number, such as reductions in ductility. Tensile data for 15 μm films were not acquired beyond cycle 4, as those films were extremely brittle and could not be reliably characterized in humid air or liquid water. Figure 9a−d summarizes the outcomes of tensile tests performed in air at 50% RH (“humid”). First, we look at the thickness dependence of mechanical properties of the as-cast (cycle 0) films. Figure 9a shows that 30 μm as-cast films have a significantly higher elongation at break (160%) than 15 μm films (40%), and Figure 9b shows that stress at break is improved in 30 μm films (10 MPa) compared with 15 μm films (7 MPa). The as-cast films have different structures and water uptakes at each thickness that could partly explain these behaviors: greater water uptake could enhance ductility in thicker films, and larger domains (greater number of chains) in thicker films might provide some resistance against failure. More generally, the increased amount of bulk material from doubling the film thickness will support a greater strain and stress,62−66 and this effect likely dominates over other attributes. Figures 9c and 9d report the effects of film thickness on as-cast modulus and yield stress, respectively. The 30 μm films have a modulus of 350 MPa and yield stress of 11 MPa, while the 15 μm films have a modulus of 430 MPa and a similar yield stress of 13.5 MPa. Similar to ductility, the reduced modulus in thicker films is consistent with the elevated water content. Next, we look at the effects of hydration cycling on mechanical properties when measured at “humid” conditions. The yield strength of the as-cast 15 μm film is 13.5 MPa (cycle 0). After the first soak/dry step (cycle 1), the yield strength increases to just over 18 MPa. With further cycling, the average strength decreases back down toward 14 MPa. A similar trend is observed for yield strength of as-cast 30 μm films, where we see an increase from 11.5 to 24 MPa from cycle 0 to 1, and a gradual decrease to 21 MPa from cycle 1 to 6. Young’s modulus follows the same trend as yield strength: The modulii of as-cast 15 μm film and 30 μm films are 430 and 350 MPa, and they increase to 550 and 700 MPa, respectively, after cycle 1. With I


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well-documented for swelling of constrained polymer films, such as a surface-anchored gels 34,79 or edge-clamped sheets,35,80,81 where the surface will crease or buckle when the internal compressive stress exceeds the polymer’s critical stress. In the present study, the films are briefly clamped at one end for transferring to and from the water bath, but are otherwise unconstrained through the remainder of the 24 h soak and 24 h dry. (During the drying step, the films are placed on an absorbent wipe; we believe this is closer to a state of “free” drying than constrained, as there is no adhesion or conformal contact between these layers.) In such free films, spatial variations in stress can drive buckling, wrinkling, and rolling82−85 Figure 11 shows low-magnification optical microscopy images of as-cast and cycled films. First, we observe that the

Figure 10. (a) Elongation at break, (b) stress at break, and (c) Young’s modulus of 15 and 30 μm films when submerged in DI water (“wet”). No yield stress is observed in “wet” measurements. The average value and error bars (±1 standard deviation) are calculated from three independent measurements. Note that the 15 μm films become too brittle for tensile tests after cycle 4.

drying process, then the films could develop nanoscale porosity as documented for other amphiphilic systems.31,33,45 There is no evidence of uniform nanoscale porosity in SAXS data, which would increase the scattering contrast. However, scattering measurements are not very sensitive to low populations of uncorrelated voids. (2) During the drying process, the tBS endsegments will vitrify before the SS domains are dry. As water continues to escape, the SS domains might collapse and densify, placing the rubbery HI segments under tension. Distorted chain formations have been reported in other types of block copolymers after swelling and deswelling in selective solvents.67 A composite X-ray scattering plot of 15 μm films in SAXS and WAXS regions is shown in Figure S13. Neither SAXS nor WAXS offers support for this type of mechanism: SAXS data are consistent with a slight expansion of the SS domains rather than a contraction. WAXS measurements detect a peak near 12.6 nm−1 that likely corresponds to the HI interchain distance,68,69 and this does not change as a function of cycle number (Figure S14). (3) Ionic aggregation is welldocumented in polymer ionomers and can impact a variety of dynamical and mechanical properties.70,71 SAXS measurements detect a peak near 2.9 nm−1 that is consistent with the formation of ionic aggregates,72−74 but there is no indication that the aggregate structure changes with cycling (Figure S15). The likely cause of failure in these materials is cyclinginduced mechanical instabilities. In polymer electrolyte membranes (PEMs), water absorption/desorption behaviors can lead to changes in residual stress and buckling.35,75−77 In elastoplastic materials, buckling may produce high levels of local strains that drive plastic deformation.35,78 This behavior is

Figure 11. Optical microscopy of 15 and 30 μm films at different cycles. Thin films exhibit numerous sharp folds and buckles within the field of view after a single cycle.

as-cast films are very flat, but cycling induces buckles with a periodicity on the order of millimeters. We suspect that spatial variations in water content during swelling and/or deswelling are driving the buckling process: the initial hydrophobicity makes it difficult to submerge the films uniformly, leading to uneven water uptake, and during drying the concentration of water may vary laterally and with depth. Second, the thinner films are far more susceptible to mechanical instabilities than thicker films, and show a dense population of buckles and sharp folds after a single cycle. As a result of these instabilities, we believe the thinner films are heavily strained within a few cycles, which can lead to reduced water uptake (Figure 4), reduced extent of swelling (Figure S11), and severe deterioration of ductility (Figure 9). This hypothesis is also supported by other studies, as a high degree of plastic strain in polymer membranes is known to limit swelling and water uptake.76,77 The thicker films also show signs of buckling after a single cycle, but these deformations are far more subtle than in thinner films: gentle rolls with a long wavelength rather than a high density of large bends and sharp folds. J


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Polymer Science. J. Polym. Sci., Part B: Polym. Phys. 2010, 48 (15), 1685−1718. (4) Geise, G. M.; Freeman, B. D.; Paul, D. R. Characterization of a Sulfonated Pentablock Copolymer for Desalination Applications. Polymer 2010, 51 (24), 5815−5822. (5) Park, M. J.; Kim, S. Y. Ion Transport in Sulfonated Polymers. J. Polym. Sci., Part B: Polym. Phys. 2013, 51 (7), 481−493. (6) Wang, F.; Hickner, M.; Kim, Y. S.; Zawodzinski, T. A.; McGrath, J. E. Direct Polymerization of Sulfonated Poly(Arylene Ether Sulfone) Random (Statistical) Copolymers: Candidates for New Proton Exchange Membranes. J. Membr. Sci. 2002, 197 (1−2), 231−242. (7) Elabd, Y. A.; Hickner, M. A. Block Copolymers for Fuel Cells. Macromolecules 2011, 44 (1), 1−11. (8) Patel, S. N.; Javier, A. E.; Balsara, N. P. Electrochemically Oxidized Electronic and Ionic Conducting Nanostructured Block Copolymers for Lithium Battery Electrodes. ACS Nano 2013, 7 (7), 6056−6068. (9) Inceoglu, S.; Rojas, A. A.; Devaux, D.; Chen, X. C.; Stone, G. M.; Balsara, N. P. Morphology−Conductivity Relationship of Single-IonConducting Block Copolymer Electrolytes for Lithium Batteries. ACS Macro Lett. 2014, 3 (6), 510−514. (10) Lufrano, F.; Staiti, P. Performance Improvement of Nafion Based Solid State Electrochemical Supercapacitor. Electrochim. Acta 2004, 49 (16), 2683−2689. (11) Kim, B. C.; Kwon, J. S.; Ko, J. M.; Park, J. H.; Too, C. O.; Wallace, G. G. Preparation and Enhanced Stability of Flexible Supercapacitor Prepared from Nafion/Polyaniline Nanofiber. Synth. Met. 2010, 160 (1−2), 94−98. (12) Park, K.-W.; Ahn, H.-J.; Sung, Y.-E. All-Solid-State Supercapacitor Using a Nafion® Polymer Membrane and Its Hybridization with a Direct Methanol Fuel Cell. J. Power Sources 2002, 109 (2), 500− 506. (13) Hickner, M. A. Ion-Containing Polymers: New Energy & Clean Water. Mater. Today 2010, 13 (5), 34−41. (14) Strathmann, H.; Grabowski, A.; Eigenberger, G. Ion-Exchange Membranes in the Chemical Process Industry. Ind. Eng. Chem. Res. 2013, 52 (31), 10364−10379. (15) Xu, T. Ion Exchange Membranes: State of Their Development and Perspective. J. Membr. Sci. 2005, 263 (1−2), 1−29. (16) Park, H. B.; Freeman, B. D.; Zhang, Z.-B.; Sankir, M.; McGrath, J. E. Highly Chlorine-Tolerant Polymers for Desalination. Angew. Chem., Int. Ed. 2008, 47 (32), 6019−6024. (17) Paul, M.; Park, H. B.; Freeman, B. D.; Roy, A.; McGrath, J. E.; Riffle, J. S. Synthesis and Crosslinking of Partially Disulfonated Poly(Arylene Ether Sulfone) Random Copolymers as Candidates for Chlorine Resistant Reverse Osmosis Membranes. Polymer 2008, 49 (9), 2243−2252. (18) Willis, C. L.; Handlin, Jr., D. L. Sulfonated Block Copolymers, Method for Making Same, and Various Uses for Such Block Copolymers. 20070021569. (19) Elabd, Y. A.; Napadensky, E.; Walker, C. W.; Winey, K. I. Transport Properties of Sulfonated Poly(Styrene-b-Isobutylene-bStyrene) Triblock Copolymers at High Ion-Exchange Capacities. Macromolecules 2006, 39 (1), 399−407. (20) Elabd, Y. A.; Napadensky, E.; Sloan, J. M.; Crawford, D. M.; Walker, C. W. Triblock Copolymer Ionomer Membranes: Part I. Methanol and Proton Transport. J. Membr. Sci. 2003, 217 (1−2), 227−242. (21) Elabd, Y. A.; Napadensky, E. Sulfonation and Characterization of Poly(Styrene-Isobutylene-Styrene) Triblock Copolymers at High Ion-Exchange Capacities. Polymer 2004, 45 (9), 3037−3043. (22) Fan, Y.; Zhang, M.; Moore, R. B.; Cornelius, C. J. Structure, Physical Properties, and Molecule Transport of Gas, Liquid, and Ions within a Pentablock Copolymer. J. Membr. Sci. 2014, 464, 179−187. (23) Laprade, E. J.; Liaw, C.-Y.; Jiang, Z.; Shull, K. R. Mechanical and Microstructural Characterization of Sulfonated Pentablock Copolymer Membranes. J. Polym. Sci., Part B: Polym. Phys. 2015, 53 (1), 39−47. (24) Geise, G. M.; Willis, C. L.; Doherty, C. M.; Hill, A. J.; Bastow, T. J.; Ford, J.; Winey, K. I.; Freeman, B. D.; Paul, D. R. Characterization

4. CONCLUSIONS We investigated the structure and properties of sulfonated pentablock terpolymer films through multiple hydration cycles. We demonstrate that cycling drives a progressive transition of SS domains into a more network-like state, which should increase the rates of water and proton transport. However, enhanced transport might not be observed, as changes to film mechanics can reduce swelling and water uptake. The competition between these effects is apparent when comparing different film thicknesses: Thin films lose ductility within a few cycles, an effect that is correlated with buckling, and both swelling and water uptake (at ambient humidity) are similarly diminished. However, the proton conductivity is nearly invariant with increasing number of cycles, which shows that structural rearrangements can compensate for the loss of water. Thick films also lose ductility through cycling, but to a lesser extent, and the films retain their ability to swell. In these systems, water uptake and proton conductivity are both increased through cycling, which is the anticipated trend based on structural analysis. Clearly, these pentablock copolymers are very challenging to study: the ill-defined structures are controlled by processing conditions, and film thickness can have a large impact on the evolution of certain properties (mechanics, swelling, and water uptake). However, by employing a large suite of complementary experimental techniques, it is possible to interrogate the complex structure− property−processing relations as a function of time and cyclic environmental conditions.

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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.8b00194. Figures S1−S17; Tables S1 and S2 (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected] (G.E.S.). ORCID

Gila E. Stein: 0000-0002-3973-4496 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS G.E.S. and P.V.T. acknowledge partial financial support from Kraton Polymers LLC and NSF DMR 1731292. R.L.B. acknowledges an NSF REU supplement from NSF CMMI 1727517. M.K. and H.A. acknowledge partial financial support from TcSUH. The authors also thank Karen Winey, Phil Griffin, and Richard Blackwell for helpful discussions.

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REFERENCES

(1) Fan, Y.; Zhang, M.; Moore, R. B.; Lee, H.-S.; McGrath, J. E.; Cornelius, C. J. Viscoelastic and Gas Transport Properties of a Series of Multiblock Copolymer Ionomers. Polymer 2011, 52 (18), 3963− 3969. (2) Nobile, M. A. D.; Mensitieri, G.; Nicolais, L.; Weiss, R. A. Gas Transport through Ethylene-Acrylic Acid Ionomers. J. Polym. Sci., Part B: Polym. Phys. 1995, 33 (8), 1269−1280. (3) Geise, G. M.; Lee, H.-S.; Miller, D. J.; Freeman, B. D.; McGrath, J. E.; Paul, D. R. Water Purification by Membranes: The Role of K


Article

Macromolecules of Aluminum-Neutralized Sulfonated Styrenic Pentablock Copolymer Films. Ind. Eng. Chem. Res. 2013, 52 (3), 1056−1068. (25) Choi, J.-H.; Kota, A.; Winey, K. I. Micellar Morphology in Sulfonated Pentablock Copolymer Solutions. Ind. Eng. Chem. Res. 2010, 49 (23), 12093−12097. (26) Choi, J.-H.; Willis, C. L.; Winey, K. I. Structure−property Relationship in Sulfonated Pentablock Copolymers. J. Membr. Sci. 2012, 394−395, 169−174. (27) Mineart, K. P.; Jiang, X.; Jinnai, H.; Takahara, A.; Spontak, R. J. Morphological Investigation of Midblock-Sulfonated Block Ionomers Prepared from Solvents Differing in Polarity. Macromol. Rapid Commun. 2015, 36 (5), 432−438. (28) Lonsdale, H. K. The Growth of Membrane Technology. J. Membr. Sci. 1982, 10 (2), 81−181. (29) Petersen, R. J. Composite Reverse Osmosis and Nanofiltration Membranes. J. Membr. Sci. 1993, 83 (1), 81−150. (30) Suzuki, T.; Tanaka, R.; Tahara, M.; Isamu, Y.; Niinae, M.; Lin, L.; Wang, J.; Luh, J.; Coronell, O. Relationship between Performance Deterioration of a Polyamide Reverse Osmosis Membrane Used in a Seawater Desalination Plant and Changes in Its Physicochemical Properties. Water Res. 2016, 100, 326−336. (31) Wang, Y. Nondestructive Creation of Ordered Nanopores by Selective Swelling of Block Copolymers: Toward Homoporous Membranes. Acc. Chem. Res. 2016, 49 (7), 1401−1408. (32) Mei, S.; Jin, Z. Mesoporous Block-Copolymer Nanospheres Prepared by Selective Swelling. Small 2013, 9 (2), 322−329. (33) Zavala-Rivera, P.; Channon, K.; Nguyen, V.; Sivaniah, E.; Kabra, D.; Friend, R. H.; Nataraj, S. K.; Al-Muhtaseb, S. A.; Hexemer, A.; Calvo, M. E.; Miguez, H. Collective Osmotic Shock in Ordered Materials. Nat. Mater. 2012, 11 (1), 53−57. (34) Chen, D.; Yoon, J.; Chandra, D.; Crosby, A. J.; Hayward, R. C. Stimuli-Responsive Buckling Mechanics of Polymer Films. J. Polym. Sci., Part B: Polym. Phys. 2014, 52 (22), 1441−1461. (35) Uchiyama, T.; Kato, M.; Yoshida, T. Buckling Deformation of Polymer Electrolyte Membrane and Membrane Electrode Assembly under Humidity Cycles. J. Power Sources 2012, 206, 37−46. (36) Willis, C.; Handlin, D.; Trenor, S.; Mather, B. Sulfonated Block Copolymers, Method for Making Same, and Various Uses for Such Block Copolymers. US20070021569 A1, January 25, 2007. (37) Dado, G. P.; Handlin, J.; Willis, C. L. United States Patent Application: 0100048817-Method of Sulfonation of Block Polymers. US20100048817 A1, January. (38) Lee, K. H.; Chu, J. Y.; Kim, A. R.; Nahm, K. S.; Yoo, D. J. Highly Sulfonated Poly(Arylene Biphenylsulfone Ketone) Block Copolymers Prepared via Post-Sulfonation for Proton Conducting Electrolyte Membranes. Bull. Korean Chem. Soc. 2013, 34 (6), 1763−1770. (39) Zhao, Y.; Yin, J. Synthesis and Evaluation of All-BlockSulfonated Copolymers as Proton Exchange Membranes for Fuel Cell Application. J. Membr. Sci. 2010, 351 (1−2), 28−35. (40) ASTM Standard Test Methods for Vulcanized Rubber and Thermoplastic Elastomers - Tension. 2006, ASTM D 412 (41) Choi, J.-H.; Kota, A.; Winey, K. I. Micellar Morphology in Sulfonated Pentablock Copolymer Solutions. Ind. Eng. Chem. Res. 2010, 49 (23), 12093−12097. (42) Choi, J.-H.; Willis, C. L.; Winey, K. I. Structure-Property Relationship in Sulfonated Pentablock Copolymers. J. Membr. Sci. 2012, 394-395, 169−174. (43) Polymer Handbook, 4th ed.; Brandrup, J., Immergut, E. H., Grulke, E. A., Eds.; Wiley-Interscience: 1999; Vol. 2. (44) Neutron Activation Calculator; http://www.ncnr.nist.gov/ resources/activation/ (accessed Jun 1, 2015). (45) Wang, Z.; Yao, X.; Wang, Y. Swelling-Induced Mesoporous Block Copolymer Membranes with Intrinsically Active Surfaces for Size-Selective Separation. J. Mater. Chem. 2012, 22 (38), 20542. (46) Pedersen, J. S. Determination of Size Distribution from SmallAngle Scattering Data for Systems with Effective Hard-Sphere Interactions. J. Appl. Crystallogr. 1994, 27 (4), 595−608.

(47) Kinning, D. J.; Thomas, E. L.; Fetters, L. J. Morphological Studies of Micelle Formation in Block Copolymer/Homopolymer Blends. J. Chem. Phys. 1989, 90 (10), 5806. (48) Pedersen, J. S. Analysis of Small-Angle Scattering Data from Colloids and Polymer Solutions: Modeling and Least-Squares Fitting. Adv. Colloid Interface Sci. 1997, 70, 171−210. (49) Heyes, D. M.; Cass, M.; Brańka, A. C. Percolation Threshold of Hard-Sphere Fluids in between the Soft-Core and Hard-Core Limits. Mol. Phys. 2006, 104 (20−21), 3137−3146. (50) Schubert, K.-V.; Strey, R.; Kline, S. R.; Kaler, E. W. Small Angle Neutron Scattering near Lifshitz Lines: Transition from Weakly Structured Mixtures to Microemulsions. J. Chem. Phys. 1994, 101 (6), 5343. (51) Teubner, M.; Strey, R. Origin of the Scattering Peak in Microemulsions. J. Chem. Phys. 1987, 87 (5), 3195−3200. (52) Socrates, G. Infrared and Raman Characteristics Group Frequencies: Tables and Charts. In Infrared and Raman Characteristics Group Frequencies: Tables and Charts; Wiley: Chichester, NY, 2001. (53) Fan, Y.; Cornelius, C. Raman Spectroscopic and Gas Transport Study of a Pentablock Ionomer Complexed with Metal Ions and Its Relationship to Physical Properties. J. Mater. Sci. 2013, 48 (3), 1153− 1161. (54) Frieberg, B. R.; Page, K. A.; Graybill, J. R.; Walker, M. L.; Stafford, C. M.; Stafford, G. R.; Soles, C. L. Mechanical Response of Thermally Annealed Nafion Thin Films. ACS Appl. Mater. Interfaces 2016, 8 (48), 33240−33249. (55) Collette, F. M.; Thominette, F.; Mendil-Jakani, H.; Gebel, G. Structure and Transport Properties of Solution-Cast Nafion® Membranes Subjected to Hygrothermal Aging. J. Membr. Sci. 2013, 435, 242−252. (56) Shi, S.; Dursch, T. J.; Blake, C.; Mukundan, R.; Borup, R. L.; Weber, A. Z.; Kusoglu, A. Impact of Hygrothermal Aging on Structure/Function Relationship of Perfluorosulfonic-Acid Membrane. J. Polym. Sci., Part B: Polym. Phys. 2016, 54 (5), 570−581. (57) Buzzoni, R.; Bordiga, S.; Ricchiardi, G.; Spoto, G.; Zecchina, A. Interaction of H2O, CH3OH, (CH3)2O, CH3CN, and Pyridine with the Superacid Perfluorosulfonic Membrane Nafion: An IR and Raman Study. J. Phys. Chem. 1995, 99 (31), 11937−11951. (58) Wijmans, J. G.; Baker, R. W. The Solution-Diffusion Model: A Review. J. Membr. Sci. 1995, 107 (1), 1−21. (59) Shi, Z.; Holdcroft, S. Synthesis and Proton Conductivity of Partially Sulfonated Poly([Vinylidene Difluoride-Co-Hexafluoropropylene]-b-Styrene) Block Copolymers. Macromolecules 2005, 38 (10), 4193−4201. (60) Kammoun, M.; Lundquist, L.; Ardebili, H. High Proton Conductivity Membrane with Coconut Shell Activated Carbon. Ionics 2015, 21 (6), 1665−1674. (61) Kelly, T.; Ghadi, B. M.; Berg, S.; Ardebili, H. In Situ Study of Strain-Dependent Ion Conductivity of Stretchable Polyethylene Oxide Electrolyte. Sci. Rep. 2016, 6, 20128. (62) Gong, T.; Lam, D. V.; Liu, R.; Won, S.; Hwangbo, Y.; Kwon, S.; Kim, J.; Sun, K.; Kim, J.-H.; Lee, S.-M.; Lee, C. Thickness Dependence of the Mechanical Properties of Free-Standing Graphene Oxide Papers. Adv. Funct. Mater. 2015, 25 (24), 3756−3763. (63) Bonnaillie, L. M.; Zhang, H.; Akkurt, S.; Yam, K. L.; Tomasula, P. M. Casein Films: The Effects of Formulation, Environmental Conditions and the Addition of Citric Pectin on the Structure and Mechanical Properties. Polymers 2014, 6 (7), 2018−2036. (64) Longares, A.; Monahan, F. J.; O’Riordan, E. D.; O’Sullivan, M. Physical Properties and Sensory Evaluation of WPI Films of Varying Thickness. LWT - Food Sci. Technol. 2004, 37 (5), 545−550. (65) Sobral, P. J. D. A. Influência Da Espessura de Biofilmes Feitos à ́ Miofibrilares Sobre Suas Propriedades Funcionais. Base de Proteinas Pesqui. Agropecu. Bras. 2000, 35 (6), 1251−1259. (66) Cuq, B.; Gontard, N.; Cuq, J.-L.; Guilbert, S. Functional Properties of Myofibrillar Protein-Based Biopackaging as Affected by Film Thickness. J. Food Sci. 1996, 61 (3), 580−584. L


Article

Macromolecules (67) Albalak, R. J.; Capel, M. S.; Thomas, E. L. Solvent Swelling of Roll-Cast Triblock Copolymer Films. Polymer 1998, 39 (8), 1647− 1656. (68) Zuo, J.; Shi, G. M.; Wei, S.; Chung, T.-S. The Development of Novel Nexar Block Copolymer/Ultem Composite Membranes for C2−C4 Alcohols Dehydration via Pervaporation. ACS Appl. Mater. Interfaces 2014, 6 (16), 13874−13883. (69) Alvarez, F.; Colmenero, J.; Zorn, R.; Willner, L.; Richter, D. Partial Structure Factors of Polyisoprene: Neutron Scattering and Molecular Dynamics Simulation. Macromolecules 2003, 36 (1), 238− 248. (70) Cheng, S.; Zhang, M.; Wu, T.; Hemp, S. T.; Mather, B. D.; Moore, R. B.; Long, T. E. Ionic Aggregation in Random Copolymers Containing Phosphonium Ionic Liquid Monomers. J. Polym. Sci., Part A: Polym. Chem. 2012, 50 (1), 166−173. (71) Weiss, R. A.; Fitzgerald, J. J.; Kim, D. Viscoelastic Behavior of Lightly Sulfonated Polystyrene Ionomers. Macromolecules 1991, 24 (5), 1071−1076. (72) Fitzgerald, J. J.; Weiss, R. A. Synthesis, Properties, and Structure of Sulfonate Ionomers. J. Macromol. Sci., Polym. Rev. 1988, 28 (1), 99− 185. (73) Lu, X.; Steckle, W. P.; Weiss, R. A. Ionic Aggregation in a Block Copolymer Ionomer. Macromolecules 1993, 26 (22), 5876−5884. (74) Song, J.-M.; Shin, J.; Sohn, J.-Y.; Nho, Y. C. Ionic Aggregation Characterization of Sulfonated PEEK Ionomers Using by X-Ray and DMA Techniques. Macromol. Res. 2012, 20 (5), 477−483. (75) Kusoglu, A.; Karlsson, A. M.; Santare, M. H.; Cleghorn, S.; Johnson, W. B. Mechanical Response of Fuel Cell Membranes Subjected to a Hygro-Thermal Cycle. J. Power Sources 2006, 161 (2), 987−996. (76) Kusoglu, A.; Karlsson, A. M.; Santare, M. H.; Cleghorn, S.; Johnson, W. B. Mechanical Behavior of Fuel Cell Membranes under Humidity Cycles and Effect of Swelling Anisotropy on the Fatigue Stresses. J. Power Sources 2007, 170 (2), 345−358. (77) Kusoglu, A.; Santare, M. H.; Karlsson, A. M.; Cleghorn, S.; Johnson, W. B. Numerical Investigation of Mechanical Durability in Polymer Electrolyte Membrane Fuel Cells. J. Electrochem. Soc. 2010, 157 (5), B705−B713. (78) Tallinen, T.; Åström, J. A.; Timonen, J. The Effect of Plasticity in Crumpling of Thin Sheets. Nat. Mater. 2009, 8 (1), 25−29. (79) Trujillo, V.; Kim, J.; Hayward, R. C. Creasing Instability of Surface-Attached Hydrogels. Soft Matter 2008, 4 (3), 564−569. (80) Liu, Z.; Hong, W.; Suo, Z.; Swaddiwudhipong, S.; Zhang, Y. Modeling and Simulation of Buckling of Polymeric Membrane Thin Film Gel. Comput. Mater. Sci. 2010, 49, S60−S64. (81) Mora, T.; Boudaoud, A. Buckling of Swelling Gels. Eur. Phys. J. E: Soft Matter Biol. Phys. 2006, 20 (2), 119−124. (82) Sharon, E.; Roman, B.; Marder, M.; Shin, G.-S.; Swinney, H. L. Mechanics: Buckling Cascades in Free Sheets. Nature 2002, 419 (6907), 579. (83) Kim, J.; Hanna, J. A.; Byun, M.; Santangelo, C. D.; Hayward, R. C. Designing Responsive Buckled Surfaces by Halftone Gel Lithography. Science 2012, 335 (6073), 1201−1205. (84) Byun, M.; Santangelo, C. D.; Hayward, R. C. Swelling-Driven Rolling and Anisotropic Expansion of Striped Gel Sheets. Soft Matter 2013, 9 (34), 8264−8273. (85) Luchnikov, V.; Sydorenko, O.; Stamm, M. Self-Rolled Polymer and Composite Polymer/Metal Micro- and Nanotubes with Patterned Inner Walls. Adv. Mater. 2005, 17 (9), 1177−1182.

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