Tailored CO2-philic Gas Separation Membranes ... - ACS Publications

Jul 2, 2019 - batteries,6−8 and gas separation membranes.9−11 Possessing .... (Brockmann Grade I, 58 Å) was purchased from Alfa Aesar. The .... T...
0 downloads 0 Views 3MB Size
Download PDF
Article Cite This: Macromolecules XXXX, XXX, XXX−XXX

pubs.acs.org/Macromolecules

Tailored CO2‑philic Gas Separation Membranes via One-Pot Thiol− ene Chemistry Tao Hong,† Peng-Fei Cao,*,§ Sheng Zhao,† Bingrui Li,§ Connor Smith,§ Michelle Lehmann,‡,§ Andrew J. Erwin,§,∥ Shannon M. Mahurin,§ Surendar R. Venna,⊥ Alexei P. Sokolov,†,§ and Tomonori Saito*,‡,§

Downloaded via BUFFALO STATE on July 26, 2019 at 07:00:09 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.

†

Department of Chemistry and ‡The Bredesen Center for Interdisciplinary Research and Graduate Education, University of Tennessee, Knoxville, Tennessee 37996, United States § Chemical Sciences Division, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831, United States ∥ School of Material Science and Engineering, Georgia Institute of Technology, Atlanta, Georgia 30332, United States ⊥ National Energy Technology Laboratory/AECOM, 626 Cochrans Mill Road, Pittsburgh, Pennsylvania 15236, United States S Supporting Information *

ABSTRACT: Thiol−ene chemistry draws much attention nowadays in the construction of functional polymer materials due to its versatility and fast reaction kinetics, though only a few studies have been reported on its utilization in the fabrication of elastic polymer materials. Herein, a series of elastic, poly(dimethylsiloxane)−poly(ethylene glycol) methyl ether acrylate (PDMS−PEGMEA)-based co-polymer membranes are synthesized via a one-pot thiol−ene reaction. These membranes are highly stable and exhibit tunable thermal/ mechanical properties by tailoring the cross-linker and sidechain functionality. When used for gas separation application, all grafted elastomer membranes show excellent gas permeability and selectivity, and the membrane with an optimal composition (PDMS−PEGMEA30−EOPDMS10) has reached the Robeson upper bound (CO2 permeability ∼800 barrer and α[CO2/N2] ∼39). The high permeability originates from the extremely fast chain mobility of PDMS molecules at the ambient temperature. Tailoring the PEGMEA content allows control of the α[CO2/ N2] ranging from 21 to 39 by enhancing gas solubility within the membrane matrix. This study provides a promising strategy to be utilized for the gutter layer, selective layer, or their combination in the industrial gas separation modules.

■

INTRODUCTION Elastic polymer membranes have been attracting significant research attentions due to their wide range of applications, from electronic skin devices,1,2 epidermal/wearable sensors,3−5 batteries,6−8 and gas separation membranes.9−11 Possessing elasticity,12 good thermal/chemical stability,13,14 biocompatibility,15,16 and affordable price, poly(dimethylsiloxane) (PDMS) is one of the most useful elastic materials to be exploited. PDMS possesses unique properties arising from the extremely low glass transition temperature (∼−125 °C), which allows the material to remain rubbery/elastic state even at very low operating temperatures.11,17 In the past decades, various approaches, including the incorporation of inorganic fillers,18,19 blending with other polymers,20,21 and the introduction of Hbonding units,22−25 have been reported to tune the properties of PDMS for different application purposes. To fully utilize the unique capability of PDMS, developing the rational design principle of PDMS-based polymers on a molecular level is imperative, which contributes to deepen the fundamental understanding of the field. Furthermore, by taking into account the cost-effective requirements for industry, developing a facile and efficient method, especially with the capability to be © XXXX American Chemical Society

processed at the ambient environment, will be appealing from practical aspects. With the versatile selection of “thiol” and “ene” functionality, as well as the fast reaction kinetics, the thiol−ene chemistry has become a promising approach for constructing various polymeric materials since its discovery.26−28 Moreover, the majority of thiol−ene reactions are viable at the ambient environment even with the presence of oxygen, which allows designed polymeric materials to be produced at high conversion in a fast and reliable fashion, providing great potential for future large-scale manufacturing. Despite large numbers of previous thiol−ene works, only a few studies have been reported for PDMS-based elastomers.29−32 Nazarenko and his co-workers have pioneered gas-separation membranes fabricated via thiol−ene chemistry.33,34 Various multifunctional thiols were employed to cross-link the poly(ethylene glycol) (PEG)-based acrylate via photopolymerization, and the polymerization mechanisms were also discussed in detail. Received: March 11, 2019 Revised: July 2, 2019

A

DOI: 10.1021/acs.macromol.9b00497 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules

(∼10 layers total). The polymer was then added to the syringe, steadily pushed through using the plunger, and collected in a vial. The process was repeated once more with a new syringe and a new vial. Synthesis of N-(2-(Diethylamino)ethyl)acrylamide (DEAEA). In a typical process, N,N-diethylethylenediamine (11.62 g, 100 mmol) was dissolved in 100 mL of anhydrous chloroform in a 250 mL roundbottom flask under a nitrogen atmosphere, then the flask was submerged into an ice-water bath. After stirring for 20 min, acryloyl chloride (9.70 mL, 120 mmol) solution (in 50 mL of anhydrous chloroform) was added drop-wise to the N,N-diethylethylenediamine solution. The acryloyl chloride solution was added in about 1 h, and the mixture solution was kept stirring for another hour at room temperature. The mixture solution was washed by NaOH (100 mL, 1 mol/L) once and deionized water twice. The organic layer was collected and dried by anhydrous MgSO4. The solution was rotaryevaporated under reduced pressure to remove the solvent, and 15.5 g of DEAEA was obtained (yield 90%). Proton (1H) and carbon (13C) nuclear magnetic resonance (NMR) spectra were collected using a Bruker ADVANCE III spectrometer operating at 400 MHz for 1H and 100 MHz for 13C. Chemical shifts were reported in ppm downfield from tertamethylsilane. CDCl3 (deuterated chloroform) was used as the solvent for all NMR samples. The spectra shown in Figure S1 indicate the successful synthesis of DEAEA. 1H NMR (CDCl3, 400 MHz): δ ppm: 6.50 (d, J = 16.4 Hz, 1H), 6.27−6.16 (m, 1H), 6.10 (ddd, J = 17.2, 10.2, 3.0 Hz, 1H), 5.57 (d, J = 9.1 Hz, 1H), 3.34 (q, J = 6.4, 5.9 Hz, 2H), 2.51 (pd, J = 7.9, 5.0, 4.5 Hz, 6H), 0.98 (dq, J = 7.6, 3.8, 3.2 Hz, 6H). 13C NMR (CDCl3, 100 MHz): δ ppm: 165.49, 131.13, 125.89, 76.75, 51.26, 46.67, 36.88, 11.79. Membrane Fabrication. All membranes were synthesized via the thiol−ene click reaction using Thiol-PDMS, EOPDMS/VTPDMS/ PEGDA, and PEGMEA/DEAEA. In a typical process, Thiol-PDMS (128 mg, 2.32 × 10−5 mol), EOPDMS (406 mg, 2.32 × 10−4 mol), PEGMEA (334 mg, 6.96 × 10−4 mol), and TPO (32 mg, 9.28 × 10−5 mol) were added to a 20 mL scintillation vial and dissolved in DCM (4.5 mL). The vial was shaken using a vortex mixer until the solution was homogeneous (∼10 s). The solution was allowed to rest for ∼1 min to allow bubbles to dissipate, then poured into an 8 cm diametered Teflon (PTFE) dish. The dish was placed into the bottom rack of a UV irradiation box and irradiated for 3 min (3″ distance from UV light, 92 mW/cm2, Uvitron, IntelliRay 400 system). The resulting membrane was left in the fume hood for 30 min and then dried in a vacuum oven overnight at room temperature. Finally, the cross-linked free-standing polymer membrane was detached from the PTFE dish and cut into pieces for further testing. All membranes have thickness between 150 and 200 μm. Fourier-Transform Infrared (FT-IR) Spectroscopy. Transmission FT-IR was conducted using a Nicolet iS50 FT-IR spectrometer equipped with a deuterated triglycine sulfate detector. A minimum of 128 scans was signal-averaged to obtain the final spectrum. Tensile Analysis. Rectangular stripes were cut from synthesized films into ≈7.0 × 5.0 × 0.5 mm3 specimens for tensile analysis using an Instron 3343 Universal Tensile Meter. Samples were elongated at the rate of 1 mm/min till break at ambient conditions. Rheology. Small-amplitude oscillatory shear measurements of membrane samples were carried out on an AR2000ex rheometer (TA Instruments) using 4 mm plates parallel plate geometry. The temperature was controlled by an environmental test chamber with nitrogen as a gas source. Prior to the measurement, the sample was purged at 25 °C for 1 h under a nitrogen atmosphere to ensure that thermal equilibrium was achieved. All of the samples were measured at 25 °C with the angular frequency sweep from 100 to 0.01 rad/s. The cross-linking densities of the membranes were calculated using the measured shear modulus. Differential Scanning Calorimetry (DSC). Differential scanning calorimetry (DSC) measurements were performed under argon atmosphere using a TA Instruments Q1000 with aluminum hermetic pans. Temperature-modulated DSC measurements were performed using the following procedure: the sample was equilibrated at 120 °C, isothermal for 10 min to erase the thermal history, and then cooled to

The advantages of thiol−ene reactions, which proceeded via a step-growth free-radical polymerization process rather than chain-growth acrylate polymerization process, were presented. Later, they also reported a PEG-based membrane based on trithiol cross-linker/PEG diene/PEG dithiol networks, and the resulting membranes exhibited tunable mechanical properties and gas-separation performance. Following their work, Kusuma et al. reported a series of poly(ethylene glycol) diacrylate (PEGDA)−siloxane membranes and achieved high gas permeabilities with good CO2/N2 selectivity.29 The best sample showed a CO2 permeability of 110 barrer and a CO2/N2 selectivity of 43. Herein, a series of multigrafted PDMS elastomers with variable functionalities, including poly(ethylene glycol) methyl ether acrylate and N-(2-(diethylamino)ethyl)acrylamide (see Table S1 for detailed information), are synthesized via a onepot thiol−ene reaction. Grafted PDMS elastomers were reported via other synthesis techniques, such as hydrosilylation, atom-transfer radical polymerization, ring-opening metathesis polymerization, etc.35−37 Changing the cross-linkers and grafting side-chains allows tuning of their thermal/ mechanical properties. However, as far as we know, this is the first report on the grafted PDMS elastomers combined with the advantage of both fast thiol−ene kinetics and tunable properties/functionalities. These elastic membranes are expected to be adopted in a wide range of applications, and herein, we demonstrate one of the potential industrial applications, i.e., gas separation. In a recent paper published in Science, Freeman and his co-workers emphasized the importance of highly permeable membranes for gas-separation applications.38 The similar viewpoint was also reported earlier with a more careful estimation of practical carbon capture cost.39−41 PDMS membranes are excellent candidates due to their highly permeable nature originating from the extremely fast segmental mobility at room temperature.9,11,42 Moreover, these rubbery membranes are free from aging phenomena that usually hinder the performance of many glassy polymeric membranes.43−46 To demonstrate the tunable functionalities, different types of functional units and cross-linkers are used to graft onto the polymer matrix via the thiol−ene reaction.

■

EXPERIMENTAL SECTION

Materials. (Mercaptopropyl)methylsiloxane homopolymer (Thiol-PDMS) with an average molecular weight ranged from 4000 to 7000 g/mol, acryloxy-terminated ethyleneoxide dimethylsiloxaneethyleneoxide ABA-typed block co-polymer (1700−1800 g/mol) (EOPDMS), and vinyl-terminated poly(dimethylsiloxane) (6000 g/ mol) (VTPDMS) were purchased from Gelest Inc. Poly(ethylene glycol) methyl ether acrylate (480 g/mol) (PEGMEA), poly(ethylene glycol) diacrylate (700 g/mol) (PEGDA), diethylenediamine, diphenyl(2,4,6-trimethylbenzoyl)phosphine oxide (TPO), azobisisobutyronitrile (AIBN), and inhibitor removers (for removing hydroquinone and monomethyl ether hydroquinone) were purchased from Sigma-Aldrich. Dichloromethane (DCM) and chloroform were purchased from Fisher Scientific. Activated aluminum oxide (Brockmann Grade I, 58 Å) was purchased from Alfa Aesar. The EOPDMS, VTPDMS, PEGMEA, and PEGDA had their inhibitors removed, as described below. All other chemicals were used as received. To remove the inhibitor from EOPDMS, VTPDMS, PEGMEA, and PEGDA, a column was prepared using a 20 mL syringe, cotton, inhibitor removers, and aluminum oxide. The bottom of the syringe was plugged with a small amount of cotton, then alternating layers of inhibitor removers and aluminum oxide were deposited into the syringe. Layers were added until a final volume of 10 mL was reached B

DOI: 10.1021/acs.macromol.9b00497 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules Scheme 1. Synthesis of PDMS Elastomer Membranes via Thiol−ene Click Reactiona

The photo shows the free-standing and elastic nature of the highest performing PDMS−PEGMEA30−EOPDMS10 membrane.

a

−160 °C at 3 °C/min with a modulation of ±1 °C/min, and then heated to 120 °C. The glass transition temperature (Tg) value was taken as the midpoint of the transition step in the heating process of the reversible heat flow signals. Thermogravimetric Analysis (TGA). The thermal stability of the membranes was examined using a TA Instruments Q-50 TGA. About 10 mg of the sample was placed in the platinum pan and equilibrated at room temperature. TGA measurements were conducted from room temperature to 800 °C under a nitrogen atmosphere at a rate of 20 °C/min. Small-Angle X-ray Scattering (SAXS). Small-angle X-ray scattering (SAXS) measurements were performed at the DND-CAT Sector 5ID-D beamline at Advanced Photon Source, Argonne National Laboratory. SAXS images were collected with a Mar CCD detector at a resolution of 1024 × 1024 pixels with E = 9 keV and a sample-to-detector distance of 2.42 m. The exposure time for each image is 0.2 s. The intensity I as a function of the scattering vector q is determined by averaging I over the two-dimensional image. Atomic Force Microscopy (AFM). The surface morphology of the dried membranes was examined using an AFM in the tapping and soft tapping modes using standard 8 nm n-type silicon tips, according to usual procedure. Cantilevers had a force constant in the range of ∼1.1−5.6 N/m (resonant frequency 60−100 kHz). The scanning speed was maintained in the range of 0.5−1.0 Hz for all scan sizes. Images were processed using the Gwyddion 2.49 software. Gas Permeability Measurements. The permeation measurements were performed at room temperature with single-gas tests using a custom test chamber following the constant-volume variablepressure method. Before loading, the membrane sample was mounted using a 47 mm nonporous brass disc with a hole (10 mm diameter) cut in the center and sealed with epoxy (Devcon). The mounted membrane sample was then placed on a highly porous stainless-steel support, which was for mechanical stability and shows negligible resistance to gases, and the entire assembly was installed into the test chamber. The chamber was evacuated with a mechanical pump to a base pressure of 20 mTorr. The membrane was allowed to remain in the test chamber overnight to fully degas the solvent residue and reach a steady base pressure. All permeability data are based on duplicate measurements on two different samples, and the average value is reported. The permeance and selectivity of mixed gas testing were determined using the constant pressure−variable volume (isobaric) system. The flow rate of the feed and permeate was measured using a

digital flow meter (Brooks Instrument). The feed pressure and the permeate pressure were measured using the Honeywell pressure transducer and accompanying readout. The temperature was recorded using a type K thermocouple. The thicknesses of the membranes were measured using a micrometer (Marathon electronic digital micrometer) several times, and the average value was used for the calculation of permeability. The membrane was loaded in a module made of ConFlat vacuum Flanges (from Kurt Lesker). The feed gas composition was 20 mol % CO2, 20 mol % N2, and balance Argon and had a flow rate of 10.0 mL/min. The sweep gas was pure argon and had a flow rate of 3.0 mL/min. After the introduction of the feed gas, the system was allowed to reach the steady state (approximately 2 h), and data was subsequently collected for at least another 2 h. The composition of both permeate and retentate was measured using a gas chromatograph (PerkinElmer Clarus 500). In these measurements, the permeance, (P/l), of a particular gas species, a, is defined as PA =

xpA Q p l A (Pf xfA − PpxpA )

where PA is the permeability, l is the membrane thickness, Qp is the flow rate of permeate, A is the membrane surface area, Pf is the feed pressure, Pp is the permeate pressure, xfA is the mole fraction of gas A in the feed, and xpA is the mole fraction of gas A in the permeate stream. Selectivity, α, of gas species A with respect to species B is the ratio of permeability of each gas. α=

■

PA PB

RESULTS AND DISCUSSION Synthesis of PDMS−PEGMEA-Based Membranes via Thiol−ene Reaction. The thiol−ene-based synthesis of the grafted PDMS elastomer is presented in Scheme 1. The membrane properties are tailored by the careful selection of side chains and cross-linkers. The PEGMEA and DEAEA functionalities are employed to demonstrate the effect of CO2philic groups on gas separation performance. Three kinds of cross-linkers are studied in this work: EOPDMS, VTPDMS, and PEGDA. Among them, EOPDMS and VTPDMS crossC

DOI: 10.1021/acs.macromol.9b00497 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules

and S3) was employed to monitor the conversion of thiol−ene reaction. For the PDMS−PEGMEA−EOPDMS membrane series, the significant intensity decrease of CC stretching vibration peak at 1638 cm−1 and the S−H stretching vibration of the thiol functional groups at 2560 cm−1 indicates the successful thiol−ene reaction.48,49 For the PDMS−PEGMEA− VTPDMS membrane series, due to the high MW (6000 g/ mol) of the VTPDMS, the peak intensity corresponding to the CC stretching is too low to be observed, as reported in a previous study.50 However, the decreased intensity of the S−H peak could still manifest the successful thiol−ene reaction. For the PDMS−DEAEA30−EOPDMS10 sample, the presence of the N−H stretching (3300 cm−1) and N−H bending (1548 cm−1) peaks demonstrates the incorporation of DEAEA functionality into the co-polymer matrix. Although the CC stretching vibration peak is covered by the CO stretching in the DEAEA functionality (1615−1755 cm −1),25,51 the disappearance of the alkene C−H stretching peak at 3076 cm−1 and the S−H stretching at 2560 cm−1 still indicate the formation of a cross-link network. For the PEGMEA moiety in PDMS−PEGMEA−EOPDMS, the CH2 scissoring and CH2 asymmetric bending bands are observed at 1448 and 1350 cm−1, respectively.52−54 For PDMS, the CH3 asymmetric deformation is observed at 1410 cm−1.55 The change of IR signals in the PEGMEA and PDMS can be observed in Figure 1. With the increased feed ratio of PEGMEA to EOPDMS/VTPDMS, the intensities of CH2 scissoring and CH2 asymmetric bending peaks show an increase relative to the CH3 asymmetric deformation peaks of PDMS, confirming the higher PEGMEA content in the copolymer membranes. As illustrated by thermogravimetric analysis (TGA) in Figure S4, the PDMS−PEGMEA−EOPDMS co-polymers demonstrate good thermal stability and are stable until 330 °C (indicated by 5 wt % loss). With higher PEGMEA composition, the decomposition process shows a slight shift toward a lower temperature. Based on previous studies, PDMS is typically stable until around 430 °C, whereas PEGMEA is likely to be completely decomposed at this temperature.56,57 Thus, the thermal stability of the co-polymer membrane toward a lower temperature is likely influenced by the PEGMEA content. The decomposition temperature (Td) results are summarized in Table 1. Moreover, the Td is also influenced by the cross-linker materials. For the PDMS− PEGMEA−VTPDMS membranes containing no PEG moiety in the cross-linker, the decomposition temperature shows a significant shift (10−60 °C) toward a higher temperature. Moreover, since the PDMS−DEAEA30−EOPDMS10 membrane contains DEAEA, the decomposition temperature shifts toward a lower range. The differential scanning calorimetry (DSC) data for PDMS−PEGMEA−EOPDMS membranes is shown in Figures 2a and S5. In Figure S5, for PDMS−PEGMEA0−EOPDMS20, all PEGMEA content are from the PEGMEA-PDMS backbone (no grafting PEGMEA side chains). Thus, the sparsely distributed PEGMEA moiety and short PEGMEA repeating chain in the cross-linker prevent the formation of PEGMEA crystallinity in the polymer matrix. For the samples with a grafted PEGMEA side chain, a minor crystallization peak (around −40 °C) was observed followed by a melting peak (around −13 °C), which is formed by the high grafting density of the PEGMEA moiety in the films, and is consistent with previous studies.29,58 However, no melting behavior of

linkers generate mechanically robust elastic membranes (see Figure S2). The PEGDA-cross-linked polymer membranes show poor mechanical performance and tend to crack during the testing process due to excessively high cross-link density. Therefore, only the EOPDMS- and VTPDMS-cross-linked membranes are selected for further studies. They are labeled as PDMS−PEGMEAx−EOPDMSy and PDMS−PEGMEAm− VTPDMSn, respectively. The first PDMS in the label represents the thiol-PDMS backbone; the second PEGMEA represents the side chain functionality; the last EOPDMS or VTPDMS represents the cross-linker. The subscripts x, y, m, and n are based on the feed ratio of the starting material (see Table S1) and represent the number of chains/cross-linker per Thiol-PDMS backbone. For example, PDMS−PEGMEA30− EOPDMS10 means the thiol-PDMS-co-EOPDMS with 30 PEGMEA chains and 10 EOPDMS cross-linkers per thiolPDMS. The co-polymer membrane of EOPDMS and DEAEA is labeled as PDMS−DEAEA30−EOPDMS10. According to the previous studies by Nazarenko and his coworkers, the thiol−ene polymerization of thiols and acrylates should result in a mixed mode polymerization, involving both thiol−ene step-growth polymerization between the thiol and acrylate and acrylate homopolymerization.33,34 Therefore, Scheme 1 presents two polymerization modes of the polymer matrix: (1) the connection via cross-linkers (red curves) between two thiol-PDMS backbones (black curves) and (2) the acrylate−acrylate cross-linking arising from chain-growth homopolymerization. Given the fact that the reactions were performed in the air, we expect the chain-growth acrylate homopolymerization to be hindered by oxygen due to the formation of unreactive peroxy radicals.34,47 Thus, the majority of the network is formed by cross-linkers/thiol-PDMS connection, with a small portion of EOPDMS−EOPDMS and EOPDMS−PEGMEA connecting chains. On the other hand, for VTPDMS cross-linker, since it is unlikely to undergo homopolymerization, the network structures are different compared with that of EOPDMS (see Scheme S1). In this network, we expect that cross-linker and functional side chains are predominantly connected via step-growth thiol−ene reaction with thiol-PDMS. Thermal, Morphological, and Mechanical Properties. Fourier-transform infrared (FT-IR) spectroscopy (Figures 1

Figure 1. Zoomed in FT-IR spectra (1300−3100 cm−1) of (a) PDMS−PEGMEA−EOPDMS series, (b) PDMS−PEGMEA− VTPDMS series, (c) PDMS−DEAEA30−EOPDMS10. The figures demonstrate the disappearance of the CC stretching vibration peak at 1638 cm−1, alkene C−H stretching peak at 3076 cm−1, and the S− H stretching vibration of the thiol functional groups at 2560 cm−1. D

DOI: 10.1021/acs.macromol.9b00497 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules Table 1. Summary of Thermal and Mechanical Properties of Multigrafting Elastic Membranes sample code

Tg,1 (°C)

Tg,2 (°C)

Tda (°C)

G′ (kPa)

Cx (mol/cm3) × 106

Mx (kg/mol)

Ree (nm)

d-spacing (nm)

PDMS−PEGMEA0−EOPDMS20 PDMS−PEGMEA10−EOPDMS15 PDMS−PEGMEA20−EOPDMS10 PDMS−PEGMEA30−EOPDMS10 PDMS−PEGMEA0−VTPDMS20 PDMS−PEGMEA30−VTPDMS10 PDMS−DEAEA30−EOPDMS10

−120 −127 −127 −126

−63 −61 −63 −63

346 341 337 334 395 355 284

140.0 85.6 60.0 48.1 43.4 17.2 79.2

28.5 17.4 12.2 9.8 8.8 3.5 16.1

17.3 28.4 40.5 50.5 55.9 141.2 30.6

6.8 8.8 10.5 11.7 12.3 19.5 9.1

8.1 8.2 9.9 10.3 20.3 9.0

Decomposition temperature at ∼5% weight loss.

a

peak at 0.6−0.8 nm−1 confirmed the presence of phase separation of the two polymer components. The d-spacing of PDMS−PEGMEA−EOPDMS membranes was calculated using the Bragg relation (d = 2π/qmax), where qmax is the peak position.59,60 The d-spacing values are varying in the range 8−10 nm (Table 1). The shapes of SAXS peaks remain similar (no significant broadening), whereas the d-spacing values show a slight increase with a higher PEGMEA content. However, no long-range order in their morphology was observed due to the absence of any higher order peaks in the SAXS data. The SAXS profile of PDMS−DEAEA30− EOPDMS10 and PDMS−PEGMEA−VTPDMS is shown in Figure S6. The PDMS−DEAEA30−EOPDMS10 showed similar phase separation behavior (a broad peak at 0.7 nm−1) compared with the PDMS−PEGMEA−EOPDMS series. However, for PDMS−PEGMEA−VTPDMS membranes, the PDMS−PEGMEA0−VTPDMS20 sample did not show any phase separation behavior, whereas a broad peak is present for PDMS−PEGMEA30−VTPDMS10. Consistent with DSC and SAXS data, AFM measurements reveal that all PDMS−PEGMEA−EOPDMS membrane are amorphous, and no evidence of crystallization has been observed at room temperature (Figures 3 and S7). Moreover, there are indications of phase-separated microstructures, particularly for PDMS−PEGMEA30−EOPDMS10 (Figure 3). Specifically, the phase-contrast images reveal a percolating network of cylindrical or “wormlike” domains that are reminiscent of the compartmentalized morphologies frequently

Figure 2. (a) Reversible heat flow curve of PDMS−PEGMEA− EOPDMS membranes. The inset is the derivative of reversible heat flow change. All curves have been shifted vertically to illustrate the systematic change of the peak shape and position. The Tg is determined from both the transition process in reversible heat flow curves and the peak position in the derivative curves. (b) SAXS profile of PDMS−PEGMEA−EOPDMS membranes.

PEGMEA is observed near the room-temperature range. This is very important for gas permeation applications, since the suppression of crystallization at room temperature could significantly improve the gas diffusivity of PDMS−PEGMEA−EOPDMS membranes. The glass transition temperatures are estimated from both the transition process in reversible heat flow curves (Figure 2a) and the peak position in the derivative curves (the inset of Figure 2a). The noticeable Tg of PEGMEA around −60 °C agrees well with previous studies,17 and the minor Tg peak at around −124 °C is assigned to PDMS. The presence of two Tg values suggests some phase separation between the two chemical compositions. The morphology of the PDMS−PEGMEA−EOPDMS membranes was further investigated by small-angle X-ray scattering (SAXS). In Figure 2b, the appearance of a broad

Figure 3. AFM height (a, b) and phase (c, d) images of the PDMS− PEGMEA30−EOPDMS10 membrane. The Z scales for the former are 200 nm (a) and 25 nm (b); for the latter, they are 20° (c) and 10° (d). E

DOI: 10.1021/acs.macromol.9b00497 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules Table 2. Summary of Gas Transport Properties of Multigrafting Elastic Membranes permeability (barrer) sample code PDMS−PEGMEA0−EOPDMS20 PDMS−PEGMEA10−EOPDMS15 PDMS−PEGMEA20−EOPDMS10 PDMS−PEGMEA30−EOPDMS10 PDMS−PEGMEA30−EOPDMS10a PDMS−PEGMEA30−EOPDMS10b PDMS−PEGMEA0−VTPDMS20 PDMS−PEGMEA30−VTPDMS10 PDMS−DEAEA30−EOPDMS10

CO2 1200 1000 930 820 810 930 4600 3350 970

± ± ± ± ± ± ± ± ±

N2 30 40 30 20 15 15 50 60 20

58 35 25 21 20 42 430 225 61

± ± ± ± ± ± ± ± ±

selectivity CH4

2 2 1 1 1 1 7 5 2

180 140 85 75 73

± ± ± ± ±

CO2/N2 4 5 5 2 2

21.2 28.6 37.5 39.0 39.9 22.1 10.6 14.9 15.9

1300 ± 20 840 ± 15 140 ± 5

± ± ± ± ± ± ± ± ±

0.6 1.0 0.9 1.2 1.4 0.9 0.4 0.6 0.8

CO2/CH4 6.7 7.3 10.9 11.0 11.1

± ± ± ± ±

0.3 0.4 0.6 0.5 0.5

3.5 ± 0.1 4.0 ± 0.2 7.0 ± 0.3

a

Re-measured after 90 days. bMixed gas data.

observed in branched block co-polymers.61 These distinct, interconnected regions cannot be accounted for the membrane morphology and are likely arised from the compositional contrast, e.g., the nanophase separation between EO and DMS functionalities. The cylinder radii (ca. 10−30 nm) are of the same scale as the d-spacings derived from the characteristic SAXS peaks (8−10 nm). The PDMS−PEGMEA−EOPDMS and PDMS−PEGMEA−VTPDMS membranes are fabricated via the UVtriggered thiol−ene cross-linking reaction of liquid starting materials, and the obtained membranes remain in rubbery state at room temperature. The mechanical property of grafted elastomer membranes is tested by tensile analysis and rheology, where the results of stress−strain and storage shear modulus (G′) are shown in Figures S8 and S9, respectively. The PDMS−PEGMEA−EOPDMS membrane series shows 20−35% elongation at breaks and ultimate tensile stress of 20− 200 kPa. The higher PEGMEA content leads to lower tensile stress due to the lower cross-link densities of the polymer network. For rheology tests, a rubbery plateau could be observed for all of the membranes, and the data are summarized in Table 1. The G′ values of rubbery materials may be in the range of 0.01−0.7 MPa,11,25,37,62,63 and the G′ values of our samples are around 0.05−0.1 MPa, which are comparable with typical rubbery materials. The G′ value of PDMS−DEAEA30−EOPDMS10 (79.2 kPa) falls within the G′ range of PDMS−PEGMEA−EOPDMS series (48.1−140 kPa, see Table 1), due to the same cross-linker matrix. The G′ values of the PDMS−PEGMEA−VTPDMS membranes are lower than that of the PDMS−PEGMEA−EOPDMS membranes. This is explained by the difference of cross-link densities of the membranes, which is estimated using the G′ value11 cx =

small portion of longer dangling chains formed by acrylate− acrylate homopolymerization. The difference between the side chain could also contribute to the difference in mechanical properties. Based on the obtained Mx values, the distance between the cross-link junctions is estimated when Gaussian-like chains are assumed. The end-to-end distance is calculated using the following equation64 ⟨R ee2⟩ = C∞ × l 2 ×

2Mx M0

where Ree is the end-to-end distance between the cross-link junctions, C∞ is the characteristic ratio at an infinite chain length, l is the Si−O bond length, Mx/M0 gives the number of repeating units. In our estimation, for PDMS chains, C∞ value of 6.6,64 l value of 0.164 nm,65 and M0 of 133, which is the MW of the thiol-PDMS monomer unit, are used. The calculated Ree values agree reasonably well with the d-spacing values estimated from SAXS data (Table 1). Given the fact that no crystallization is observed from the DSC tests, we believe that the phase separation behaviors observed from SAXS and AFM are likely originated from the contrast of PEGMEA functionality and PDMS backbones. Gas Transport Properties. Gas permeabilities of the elastic polymer membranes for three different gases (CO2, N2, and CH4) are tested and summarized in Table 2. Although not obtained experimentally, the O2 permeability values of PDMS−PEGMEA−EOPDMS, PDMS−PEGMEA−VTPDMS, and PDMS−DEAEA30−EOPDMS10 are expected to be within the permeability range of pure PDMS and pure PEO9,10,66 and may vary slightly depending on the composition ratio. The membranes made of PEO and PDMS normally show O2/N2 selectivity ranged from 2 to 3. The raw pressure rise data is presented in Figure S10, which clearly demonstrates that the slope of CO2 pressure is much higher than that of N2 and CH4. The utilization of in situ thiol−ene reactions generates mechanically robust PDMS−PEGMEA−EOPDMS multigrafting co-polymer networks. The elastic PDMS backbones provide fast segmental mobility at the ambient temperature, which yields high gas permeabilities. The relatively low MW (480 Da) of PEGMEA and the introduction of the cross-linked network prevent the formation of the crystalline regime at room temperature, which ensures the minimum loss of gas permeabilities.17,23 Moreover, as we reported in our previous cross-linked (bicycloheptenyl)ethyl-terminated poly(dimethylsiloxane) (XLPDMSPNB) work, lower cross-link

ρ G′ = 2RT 2Mx

where cx is the number of moles of cross-links per unit volume, R is the gas constant, T is the temperature in kelvin, ρ is the density, Mx is the number-average molecular weight of the polymer segments between cross-links. The VTPDMS (6000 g/mol) has much higher MW than EOPDMS (1700−1800 g/mol). The steric hindrance of the longer chains reduces the cross-link density of the polymer matrixes and, therefore, causes a lower G′. As mentioned before, VTPDMS is unlikely to undergo homopolymerization. Thus, the PEGMEA chains are present as short dangling ends in the network. However, in the EOPDMS matrix, there is a F

DOI: 10.1021/acs.macromol.9b00497 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules densities (5.7 × 10−6 to 3.17 × 10−5 mol/cm3) of the polymer network could further contribute to faster chain mobility, and therefore, the higher gas permeabilities.11 In this study, by careful selection of the MW and reaction ratio of the starting materials, we are able to obtain lightly cross-linked networks (3.5 × 10−6 to 2.85 × 10−5 mol/cm3), which are comparable with XLPDMSPNB. We believe that it is a crucial design parameter to obtain highly permeable rubbery membranes. The incorporation of CO2-philic PEGMEA in both crosslinkers and side chains contributed to the significant improvement of gas selectivities. By comparison with the membranes containing no PEG moiety in the cross-linker, i.e., PDMS−PEGMEA0−VTPDMS20 and PDMS−PEGMEA30− VTPDMS10, the PDMS−PEGMEA30−EOPDMS10 showed more than 3 times improvement in terms of CO2 /N 2 selectivity. The co-polymer membrane with DEAEA functionality, i.e., PDMS−DEAEA30−EOPDMS10, showed a much lower selectivity of CO2 over N2 and CH4, due to the weaker CO2-philicity of DEAEA functionality. This comparison clearly demonstrates the importance of the EO group in both the cross-linker and side chain. Moreover, for PDMS−PEGMEA−EOPDMS series, as the overall PEGMEA/EOPDMS ratio increases from 0/20 to 30/ 10, the CO2/N2 selectivity of PDMS−PEGMEA−EOPDMS showed about 2-fold increase, while maintaining high CO2 permeability (over 800 barrer). The drop of gas permeabilities with a higher PEGMEA composition originates from a much slower segmental mobility of PEGMEA, which was reported in previous studies.17 The mixed gas permeability of the highest performing membrane, i.e., PDMS−PEGMEA30−EOPDMS10, is also studied, and the data are presented in Table 2. The CO2 permeability showed a slight increase in the mixed gas test, whereas the CO2/N2 selectivity decreased. Similar behavior was reported by other researchers, and it is proposed to be caused by the plasticization effect of the condensable CO2 molecules.66 The PDMS−PEGMEA−VTPDMS series contains no PEGMEA functionality in the backbone. Therefore, the higher PDMS content offers increased gas permeabilities, which could be used in the areas where more permeable but less selective materials are needed. The gas separation performance of synthesized membranes is summarized in 2008 Robeson plot (Figure 4).67,68 All grafted elastomer membranes show excellent gas permeability and selectivity, and the membrane with an optimal composition (PDMS−PEGMEA 30 − EOPDMS10) has reached the Robeson upper bound (CO2

permeability ∼800 barrer and α[CO2/N2] ∼39). Compared with other EO-containing CO2/N2 separation polymers, e.g., pebax69−71 and polyactive,72−74 PDMS−PEGMEA−EOPDMS series showed a lower CO2/N2 selectivity and higher gas permeability. This is attributed to the presence of PDMS segments, which usually offer a higher permeating and lower selective nature due to their fast dynamics. Moreover, in current membrane gas separation field, the majority of highperforming glassy membranes, e.g., polymers of intrinsic microporosity, poly(1-trimethylsilyl-1-propyne), suffers from aging, which is caused by the collapse of excess free volume at room temperature,75,76 leading to the significant drop in permeability.77−80 In our study, the highest performing PDMS−PEGMEA30−EOPDMS10 is retested after 90 days, and the data is presented in Table 2. Within the experimental uncertainty range (±5%), the gas permeabilities of the “aged” sample are comparable with the “fresh” one. This result confirms the long-term stability of the grafted PDMS elastic membranes. The high-performance elastic membranes with no aging issues potentially offer a game-changing platform as a durable and versatile candidate for the industrial applications.

■

CONCLUSIONS In conclusion, we have developed a series of multigrafting copolymer based elastic membranes via the one-pot thiol−ene reaction in 3 min. The synthesized membranes show a freestanding and elastic nature, with slight phase separation behavior among the different components (PEGMEA/DEAEA and PDMS). The co-polymer membranes show noncrystalline behavior at room temperature, guaranteeing minimum gas permeability loss. The PEGMEA content in both cross-linker and side chain contributes to high CO2-philicity of the synthesized membranes. The polymer membrane with the PEGMEA/EOPDMS ratio of 30/10 (PDMS−PEGMEA30− EOPDMS10), exhibiting CO2 permeability ∼800 barrer and CO2/N2 selectivity ∼39, has reached the Robeson upper bound. The versatility of the thiol−ene technique allows us to substitute a variety of cross-linker/side chain into the rubbery system, providing tunable gas separation performance for future development. The developed membrane can be utilized as a gutter layer, selective layer, or both in the industrial gas separation modules. These findings provide strong support to practical CO2 separation material development as well as fostering a fundamental understanding of the gas separation process using polymer membranes.

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.9b00497.

■

Synthesis scheme and additional characterization data (PDF)

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (P.-F.C.). *E-mail: [email protected] (T.S.). ORCID

Peng-Fei Cao: 0000-0003-2391-1838 Bingrui Li: 0000-0002-4974-5826

Figure 4. Summary of the gas separation performance of synthesized membranes in 2008 Robeson plot. G

DOI: 10.1021/acs.macromol.9b00497 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules

(15) Bélanger, M. C.; Marois, Y. Hemocompatibility, biocompatibility, inflammatory and in vivo studies of primary reference materials low-density polyethylene and polydimethylsiloxane: A review. J. Biomed. Mater. Res. 2001, 58, 467−477. (16) Peterson, S. L.; McDonald, A.; Gourley, P. L.; Sasaki, D. Y. Poly (dimethylsiloxane) thin films as biocompatible coatings for microfluidic devices: cell culture and flow studies with glial cells. J. Biomed. Mater. Res. 2005, 72A, 10−18. (17) Hong, T.; Lai, S.; Mahurin, S. M.; Cao, P. F.; Voylov, D. N.; Meyer, H. M., III; Jacobs, C. B.; Carrillo, J. M. Y.; Kisliuk, A.; Ivanov, I. N.; et al. Highly Permeable Oligo (ethylene oxide)-co-poly (dimethylsiloxane) Membranes for Carbon Dioxide Separation. Adv. Sustainable Syst. 2018, 2, No. 1700113. (18) Rezakazemi, M.; Shahidi, K.; Mohammadi, T. Sorption properties of hydrogen-selective PDMS/zeolite 4A mixed matrix membrane. Int. J. Hydrogen Energy 2012, 37, 17275−17284. (19) Wang, N.; Liu, J.; Li, J.; Gao, J.; Ji, S.; Li, J.-R. Tuning properties of silicalite-1 for enhanced ethanol/water pervaporation separation in its PDMS hybrid membrane. Microporous Mesoporous Mater. 2015, 201, 35−42. (20) Lue, S. J.; Ou, J. S.; Kuo, C. H.; Chen, H. Y.; Yang, T.-H. Pervaporative separation of azeotropic methanol/toluene mixtures in polyurethane−poly (dimethylsiloxane)(PU−PDMS) blend membranes: correlation with sorption and diffusion behaviors in a binary solution system. J. Membr. Sci. 2010, 347, 108−115. (21) Martinelli, E.; Suffredini, M.; Galli, G.; Glisenti, A.; Pettitt, M. E.; Callow, M. E.; Callow, J. A.; Williams, D.; Lyall, G. Amphiphilic block copolymer/poly (dimethylsiloxane)(PDMS) blends and nanocomposites for improved fouling-release. Biofouling 2011, 27, 529− 541. (22) Zha, R. H.; de Waal, B. F.; Lutz, M.; Teunissen, A. J.; Meijer, E. End groups of functionalized siloxane oligomers direct blockcopolymeric or liquid-crystalline self-assembly behavior. J. Am. Chem. Soc. 2016, 138, 5693−5698. (23) Cao, P.-F.; Li, B.; Hong, T.; Xing, K.; Voylov, D. N.; Cheng, S.; Yin, P.; Kisliuk, A.; Mahurin, S. M.; Sokolov, A. P.; Saito, T. Robust and Elastic Polymer Membranes with Tunable Properties for Gas Separation. ACS Appl. Mater. Interfaces 2017, 9, 26483−26491. (24) Xing, K.; Tress, M.; Cao, P.; Cheng, S.; Saito, T.; Novikov, V. N.; Sokolov, A. P. Hydrogen-bond strength changes network dynamics in associating telechelic PDMS. Soft Matter 2018, 1235− 1246. (25) Cao, P. F.; Li, B.; Hong, T.; Townsend, J.; Qiang, Z.; Xing, K.; Vogiatzis, K. D.; Wang, Y.; Mays, J. W.; Sokolov, A. P.; Saito, T. Superstretchable, Self-Healing Polymeric Elastomers with Tunable Properties. Adv. Funct. Mater. 2018, 28, No. 1800741. (26) Hoyle, C. E.; Bowman, C. N. Thiol−ene click chemistry. Angew. Chem., Int. Ed. 2010, 49, 1540−1573. (27) Lowe, A. B. Thiol−ene “click” reactions and recent applications in polymer and materials synthesis: a first update. Polym. Chem. 2014, 5, 4820−4870. (28) Lowe, A. B. Thiol-ene “click” reactions and recent applications in polymer and materials synthesis. Polym. Chem. 2010, 1, 17−36. (29) Kusuma, V. A.; Roth, E. A.; Clafshenkel, W. P.; Klara, S. S.; Zhou, X.; Venna, S. R.; Albenze, E.; Luebke, D. R.; Mauter, M. S.; Koepsel, R. R.; et al. Crosslinked poly (ethylene oxide) containing siloxanes fabricated through thiol-ene photochemistry. J. Polym. Sci., Part A: Polym. Chem. 2015, 53, 1548−1557. (30) Mongkhontreerat, S.; Ö berg, K.; Erixon, L.; Löwenhielm, P.; Hult, A.; Malkoch, M. UV initiated thiol−ene chemistry: a facile and modular synthetic methodology for the construction of functional 3D networks with tunable properties. J. Mater. Chem. A 2013, 1, 13732− 13737. (31) Campos, L. M.; Meinel, I.; Guino, R. G.; Schierhorn, M.; Gupta, N.; Stucky, G. D.; Hawker, C. J. Highly Versatile and Robust Materials for Soft Imprint Lithography Based on Thiol-ene Click Chemistry. Adv. Mater. 2008, 20, 3728−3733. (32) Ye, S.-H.; Jang, Y.-S.; Yun, Y.-H.; Shankarraman, V.; Woolley, J. R.; Hong, Y.; Gamble, L. J.; Ishihara, K.; Wagner, W. R. Surface

Michelle Lehmann: 0000-0003-1323-9785 Shannon M. Mahurin: 0000-0003-3792-1631 Surendar R. Venna: 0000-0003-1094-4534 Tomonori Saito: 0000-0002-4536-7530 Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS The research is sponsored by U.S. Department of Energy, Office of Fossil Energy, Carbon Capture Program and by the Oak Ridge National Laboratory Technology Innovation Program using technology transfer license royalties. Oak Ridge National Laboratory is managed by UT-Battelle, LLC, for the U.S. Department of Energy under Contract No. DEAC05-00OR22725. The authors also acknowledge Dr. Zhe Qiang (Department of Chemical and Biological Engineering, Northwestern University) for performing the SAXS measurements.

■

REFERENCES

(1) Wang, X.; Gu, Y.; Xiong, Z.; Cui, Z.; Zhang, T. Silk-molded flexible, ultrasensitive, and highly stable electronic skin for monitoring human physiological signals. Adv. Mater. 2014, 26, 1336−1342. (2) Wang, X.; Dong, L.; Zhang, H.; Yu, R.; Pan, C.; Wang, Z. L. Recent progress in electronic skin. Adv. Sci. 2015, 2, No. 1500169. (3) Jeong, S. H.; Zhang, S.; Hjort, K.; Hilborn, J.; Wu, Z. PDMSBased Elastomer Tuned Soft, Stretchable, and Sticky for Epidermal Electronics. Adv. Mater. 2016, 28, 5830−5836. (4) Pang, C.; Lee, G.-Y.; Kim, T.-i.; Kim, S. M.; Kim, H. N.; Ahn, S.H.; Suh, K.-Y. A flexible and highly sensitive strain-gauge sensor using reversible interlocking of nanofibres. Nat. Mater. 2012, 11, 795. (5) Wang, Y.; Wang, L.; Yang, T.; Li, X.; Zang, X.; Zhu, M.; Wang, K.; Wu, D.; Zhu, H. Wearable and highly sensitive graphene strain sensors for human motion monitoring. Adv. Funct. Mater. 2014, 24, 4666−4670. (6) Weng, W.; Sun, Q.; Zhang, Y.; He, S.; Wu, Q.; Deng, J.; Fang, X.; Guan, G.; Ren, J.; Peng, H. A Gum-Like Lithium-Ion Battery Based on a Novel Arched Structure. Adv. Mater. 2015, 27, 1363− 1369. (7) Zhu, B.; Jin, Y.; Hu, X.; Zheng, Q.; Zhang, S.; Wang, Q.; Zhu, J. Poly (dimethylsiloxane) Thin Film as a Stable Interfacial Layer for High-Performance Lithium-Metal Battery Anodes. Adv. Mater. 2017, 29, No. 1603755. (8) Yang, X.; Xu, J.; Bao, D.; Chang, Z.; Liu, D.; Zhang, Y.; Zhang, X. B. High-Performance Integrated Self-Package Flexible Li−O2 Battery Based on Stable Composite Anode and Flexible Gas Diffusion Layer. Adv. Mater. 2017, 29, No. 1700378. (9) Merkel, T.; Bondar, V.; Nagai, K.; Freeman, B.; Pinnau, I. Gas sorption, diffusion, and permeation in poly (dimethylsiloxane). J. Polym. Sci., Part B: Polym. Phys. 2000, 38, 415−434. (10) Lin, H.; Freeman, B. D. Gas solubility, diffusivity and permeability in poly (ethylene oxide). J. Membr. Sci. 2004, 239, 105−117. (11) Hong, T.; Niu, Z.; Hu, X.; Gmernicki, K.; Cheng, S.; Fan, F.; Johnson, J. C.; Hong, E.; Mahurin, S.; Jiang, D.; et al. Effect of CrossLink Density on Carbon Dioxide Separation in PolydimethylsiloxaneNorbornene Membranes. ChemSusChem 2015, 8, 3595−3604. (12) Lötters, J. C.; Olthuis, W.; Veltink, P.; Bergveld, P. The mechanical properties of the rubber elastic polymer polydimethylsiloxane for sensor applications. J. Micromech. Microeng. 1997, 7, 145. (13) Camino, G.; Lomakin, S.; Lageard, M. Thermal polydimethylsiloxane degradation. Part 2. The degradation mechanisms. Polymer 2002, 43, 2011−2015. (14) Camino, G.; Lomakin, S.; Lazzari, M. Polydimethylsiloxane thermal degradation Part 1. Kinetic aspects. Polymer 2001, 42, 2395− 2402. H

DOI: 10.1021/acs.macromol.9b00497 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules modification of a biodegradable magnesium alloy with phosphorylcholine (PC) and sulfobetaine (SB) functional macromolecules for reduced thrombogenicity and acute corrosion resistance. Langmuir 2013, 29, 8320−8327. (33) Kwisnek, L.; Goetz, J.; Meyers, K. P.; Heinz, S. R.; Wiggins, J. S.; Nazarenko, S. PEG containing thiol−ene network membranes for CO2 separation: effect of cross-linking on thermal, mechanical, and gas transport properties. Macromolecules 2014, 47, 3243−3253. (34) Kwisnek, L.; Heinz, S.; Wiggins, J. S.; Nazarenko, S. Multifunctional thiols as additives in UV-cured PEG-diacrylate membranes for CO2 separation. J. Membr. Sci. 2011, 369, 429−436. (35) Cai, L. H.; Kodger, T. E.; Guerra, R. E.; Pegoraro, A. F.; Rubinstein, M.; Weitz, D. A. Soft Poly (dimethylsiloxane) Elastomers from Architecture-Driven Entanglement Free Design. Adv. Mater. 2015, 27, 5132−5140. (36) Daniel, W. F.; Burdyńska, J.; Vatankhah-Varnoosfaderani, M.; Matyjaszewski, K.; Paturej, J.; Rubinstein, M.; Dobrynin, A. V.; Sheiko, S. S. Solvent-free, supersoft and superelastic bottlebrush melts and networks. Nat. Mater. 2016, 15, 183. (37) Pesek, S. L.; Lin, Y.-H.; Mah, H. Z.; Kasper, W.; Chen, B.; Rohde, B. J.; Robertson, M. L.; Stein, G. E.; Verduzco, R. Synthesis of bottlebrush copolymers based on poly (dimethylsiloxane) for surface active additives. Polymer 2016, 98, 495−504. (38) Park, H. B.; Kamcev, J.; Robeson, L. M.; Elimelech, M.; Freeman, B. D. Maximizing the right stuff: The trade-off between membrane permeability and selectivity. Science 2017, 356, No. eaab0530. (39) Lin, H.; He, Z.; Sun, Z.; Kniep, J.; Ng, A.; Baker, R. W.; Merkel, T. C. CO 2-selective membranes for hydrogen production and CO 2 capture−Part II: Techno-economic analysis. J. Membr. Sci. 2015, 493, 794−806. (40) Lin, H.; He, Z.; Sun, Z.; Vu, J.; Ng, A.; Mohammed, M.; Kniep, J.; Merkel, T. C.; Wu, T.; Lambrecht, R. C. CO2-selective membranes for hydrogen production and CO2 capture−Part I: Membrane development. J. Membr. Sci. 2014, 457, 149−161. (41) Merkel, T. C.; Lin, H.; Wei, X.; Baker, R. Power plant postcombustion carbon dioxide capture: an opportunity for membranes. J. Membr. Sci. 2010, 359, 126−139. (42) Hong, T.; Chatterjee, S.; Mahurin, S. M.; Fan, F.; Tian, Z.; Jiang, D.-e.; Long, B. K.; Mays, J. W.; Sokolov, A. P.; Saito, T. Impact of tuning CO 2-philicity in polydimethylsiloxane-based membranes for carbon dioxide separation. J. Membr. Sci. 2017, 530, 213−219. (43) Bernardo, P.; Bazzarelli, F.; Tasselli, F.; Clarizia, G.; Mason, C.; Maynard-Atem, L.; Budd, P.; Lanč, M.; Pilnácě k, K.; Vopička, O.; et al. Effect of physical aging on the gas transport and sorption in PIM-1 membranes. Polymer 2017, 113, 283−294. (44) Harms, S.; Rätzke, K.; Faupel, F.; Chaukura, N.; Budd, P.; Egger, W.; Ravelli, L. Aging and free volume in a polymer of intrinsic microporosity (PIM-1). J. Adhes. 2012, 88, 608−619. (45) Dorkenoo, K. D.; Pfromm, P. H. Accelerated physical aging of thin poly [1-(trimethylsilyl)-1-propyne] films. Macromolecules 2000, 33, 3747−3751. (46) Starannikova, L.; Khodzhaeva, V.; Yampolskii, Y. Mechanism of aging of poly [1-(trimethylsilyl)-1-propyne] and its effect on gas permeability. J. Membr. Sci. 2004, 244, 183−191. (47) Kloosterboer, J. G. Network Formation by Chain Crosslinking Photopolymerization and Its Applications in Electronics. In Electronic Applications; Springer, 1988; pp 1−61. (48) Chiou, B.-S.; Khan, S. A. Real-time FTIR and in situ rheological studies on the UV curing kinetics of thiol-ene polymers. Macromolecules 1997, 30, 7322−7328. (49) Zhu, L.; Zimudzi, T. J.; Li, N.; Pan, J.; Lin, B.; Hickner, M. A. Crosslinking of comb-shaped polymer anion exchange membranes via thiol−ene click chemistry. Polym. Chem. 2016, 7, 2464−2475. (50) Kim, M. S.; Lee, G. H.; Hong, J.-M.; Lee, H. Synthesis of monodisperse PS-co-PDMS microspheres by dispersion polymerization. Mater. Sci. Eng., C 2007, 27, 1247−1251.

(51) Cordier, P.; Tournilhac, F.; Soulié-Ziakovic, C.; Leibler, L. Selfhealing and thermoreversible rubber from supramolecular assembly. Nature 2008, 451, 977−980. (52) Abdelghany, A.; Abdelrazek, E.; Badr, S.; Morsi, M. Effect of gamma-irradiation on (PEO/PVP)/Au nanocomposite: Materials for electrochemical and optical applications. Mater. Des. 2016, 97, 532− 543. (53) Chapi, S.; Raghu, S.; Devendrappa, H. Enhanced electrochemical, structural, optical, thermal stability and ionic conductivity of (PEO/PVP) polymer blend electrolyte for electrochemical applications. Ionics 2016, 22, 803−814. (54) Kumar, K.; Ravi, M.; Pavani, Y.; Bhavani, S.; Sharma, A.; VVR, N. R. Electrical conduction mechanism in NaCl complexed PEO/PVP polymer blend electrolytes. J. Non-Cryst. Solids 2012, 358, 3205− 3211. (55) Bodas, D.; Khan-Malek, C. Formation of more stable hydrophilic surfaces of PDMS by plasma and chemical treatments. Microelectron. Eng. 2006, 83, 1277−1279. (56) Grainger, D.; Okano, T.; Kim, S.; Castner, D.; Ratner, B.; Briggs, D.; Sung, Y. Poly (dimethylsiloxane)-poly (ethylene oxide)heparin block copolymers III: Surface and bulk compositional differences. J. Biomed. Mater. Res. 1990, 24, 547−571. (57) Ç akmak, G.; Kücu̧ ̈kyavuz, Z.; Kücu̧ ̈kyavuz, S.; Ç akmak, H. Mechanical, electrical and thermal properties of carbon fiber reinforced poly (dimethylsiloxane)/polypyrrole composites. Composites, Part A 2004, 35, 417−421. (58) Lin, H.; Kai, T.; Freeman, B. D.; Kalakkunnath, S.; Kalika, D. S. The effect of cross-linking on gas permeability in cross-linked poly (ethylene glycol diacrylate). Macromolecules 2005, 38, 8381−8393. (59) Patel, R.; Kim, S. J.; Roh, D. K.; Kim, J. H. Synthesis of amphiphilic PCZ-r-PEG nanostructural copolymers and their use in CO2/N2 separation membranes. Chem. Eng. J. 2014, 254, 46−53. (60) Park, C. H.; Lee, J. H.; Jung, J. P.; Jung, B.; Kim, J. H. A highly selective PEGBEM-g-POEM comb copolymer membrane for CO2/ N2 separation. J. Membr. Sci. 2015, 492, 452−460. (61) Erwin, A. J.; Korolovych, V. F.; Iatridi, Z.; Tsitsilianis, C.; Ankner, J. F.; Tsukruk, V. V. Tunable Compartmentalized Morphologies of Multilayered Dual Responsive Star Block Polyampholytes. Macromolecules 2018, 51, 4800−4812. (62) Ramesh, S.; Winie, T.; Arof, A. Investigation of mechanical properties of polyvinyl chloride−polyethylene oxide (PVC−PEO) based polymer electrolytes for lithium polymer cells. Eur. Polym. J. 2007, 43, 1963−1968. (63) Shankar, R.; Ghosh, T. K.; Spontak, R. J. Electroactive nanostructured polymers as tunable actuators. Adv. Mater. 2007, 19, 2218−2223. (64) Hiemenz, P. C.; Lodge, T. P. Polymer Chemistry; CRC press, 2007. (65) Liu, Y. Silicone Dispersions; CRC Press, 2017. (66) Reijerkerk, S. R.; Knoef, M. H.; Nijmeijer, K.; Wessling, M. Poly (ethylene glycol) and poly (dimethyl siloxane): Combining their advantages into efficient CO2 gas separation membranes. J. Membr. Sci. 2010, 352, 126−135. (67) Robeson, L. M. The upper bound revisited. J. Membr. Sci. 2008, 320, 390−400. (68) Robeson, L. M. Correlation of separation factor versus permeability for polymeric membranes. J. Membr. Sci. 1991, 62, 165−185. (69) Nafisi, V.; Hägg, M.-B. Development of dual layer of ZIF-8/ PEBAX-2533 mixed matrix membrane for CO2 capture. J. Membr. Sci. 2014, 459, 244−255. (70) Car, A.; Stropnik, C.; Yave, W.; Peinemann, K.-V. Pebax/ polyethylene glycol blend thin film composite membranes for CO2 separation: Performance with mixed gases. Sep. Purif. Technol. 2008, 62, 110−117. (71) Car, A.; Stropnik, C.; Yave, W.; Peinemann, K.-V. PEG modified poly (amide-b-ethylene oxide) membranes for CO2 separation. J. Membr. Sci. 2008, 307, 88−95. I

DOI: 10.1021/acs.macromol.9b00497 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules (72) Yave, W.; Car, A.; Funari, S. S.; Nunes, S. P.; Peinemann, K.-V. CO2-philic polymer membrane with extremely high separation performance. Macromolecules 2010, 43, 326−333. (73) Rahman, M. M.; Filiz, V.; Shishatskiy, S.; Abetz, C.; Georgopanos, P.; Khan, M. M.; Neumann, S.; Abetz, V. Influence of poly (ethylene glycol) segment length on CO2 permeation and stability of polyactive membranes and their nanocomposites with PEG POSS. ACS Appl. Mater. Interfaces 2015, 7, 12289−12298. (74) Karunakaran, M.; Shevate, R.; Kumar, M.; Peinemann, K.-V. CO 2-selective PEO−PBT (PolyActive)/graphene oxide composite membranes. Chem. Commun. 2015, 51, 14187−14190. (75) Pfromm, P.; Koros, W. Accelerated physical ageing of thin glassy polymer films: evidence from gas transport measurements. Polymer 1995, 36, 2379−2387. (76) Struik, L. C. E. Physical Aging in Amorphous Polymers and Other Materials; Elsevier: Amsterdam, 1978; Vol. 106. (77) Huang, Y.; Wang, X.; Paul, D. R. Physical aging of thin glassy polymer films: Free volume interpretation. J. Membr. Sci. 2006, 277, 219−229. (78) Rowe, B. W.; Freeman, B. D.; Paul, D. R. Physical aging of ultrathin glassy polymer films tracked by gas permeability. Polymer 2009, 50, 5565−5575. (79) Nagai, K.; Nakagawa, T. Effects of aging on the gas permeability and solubility in poly (1-trimethylsilyl-1-propyne) membranes synthesized with various catalysts. J. Membr. Sci. 1995, 105, 261−272. (80) Lin, W.-H.; Chung, T.-S. Gas permeability, diffusivity, solubility, and aging characteristics of 6FDA-durene polyimide membranes. J. Membr. Sci. 2001, 186, 183−193.

J

DOI: 10.1021/acs.macromol.9b00497 Macromolecules XXXX, XXX, XXX−XXX