Hydrophilic Triblock Copolymers: Synthesis and

Jan 16, 2015 - State Key Laboratory of Fine Chemicals, Department of Polymer Science and Engineering, School of Chemical Engineering, Dalian Universit...
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Hydrophobic/Hydrophilic Triblock Copolymers: Synthesis and Properties of Physically Cross-Linked Hydrogels Hui Niu,†,‡ Fei Wang,† and R. A. Weiss*,† †

Department of Polymer Engineering, University of Akron, 250 South Forge Street, Akron, Ohio 44325-0301, United States State Key Laboratory of Fine Chemicals, Department of Polymer Science and Engineering, School of Chemical Engineering, Dalian University of Technology, Dalian, Liaoning 116024, China

‡

S Supporting Information *

ABSTRACT: Hydrophobic/hydrophilic triblock copolymers of poly(2-(N-ethylperfluorooctanesulfonamido)ethylmethyl acrylate) and poly(N,N′-dimethylacrylamide) (PD) were synthesized by sequential reversible addition−fragmentation chain transfer polymerization. Physically cross-linked hydrogels were produced by immersing compression-molded triblock copolymers into water. The copolymers and their hydrogels were characterized by differential scanning calorimetry, thermogravimetric analysis, thermal desorption-GC/MS analysis, swelling isotherms, wide- and small-angle X-ray scattering, and dynamic mechanical analysis. The equilibrium water sorption of the hydrogels depended on the length of the water-soluble polymer block (PD), and the block copolymers swelled more in water than a random copolymer of the same composition. The block copolymer hydrogels were viscoelastic, though the frequency dependence of the dynamic modulus was weak. The dynamic modulus of the block copolymer hydrogels ranged from ∼103 to 4 × 104 Pa, which was much lower than the modulus of a random copolymer hydrogel of the same composition.

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INTRODUCTION

Hydrogels are water-soluble polymers that have been crosslinked such that they absorb large quantities of water but are not water-soluble.1 Applications for such materials include biomedical (drug delivery,2 wound healing,3 tissue engineering4), personal care products (cosmetics, skin care, diapers), food (encapsulate, time release, candy),5 and agriculture (synthetic soils, encapsulates).6 The network in hydrogels may be formed by covalent cross-links, supramolecular physical cross-links, or both. Physical hydrogels that have been of particular interest in our research are random copolymers of dimethylacrylamide (DMA) and a fluoro(meth)acrylate (FA), which form a network as a consequence of nanoaggregation of hydrophobic fluorocarbon side chains when swollen with water.7−12 Earlier work on random DMA/FA copolymers used 2-(N -et hy l perfluo ro oct anesulfam ido)et hyl acrylate (FOSA).8−10 Hydrogels formed from random DMA/FA copolymers using 2-(N-ethylperfluorooctanesulfamido)ethyl acrylate (FOSM) have also been reported, and there does not appear to be any significant difference in the structure or properties of the hydrogels prepared with the two FA monomers. This paper describes the synthesis and characterization of triblock copolymer hydrogels with similar compositions as the reported random DMA/FA physical hydrogels. Reversible addition−fragmentation chain transfer (RAFT) polymerization was used to prepare FA-b-DMA-b-FA triblock copolymer hydrogels using FOSM as the FA. © 2015 American Chemical Society

EXPERIMENTAL DETAILS

Materials. N,N-Dimethylacrylamide (DMA) (Sigma-Aldrich) was purified by distillation under reduced pressure to remove the inhibitor. 2-(N-Ethylperfluorooctanesulfamido)ethyl methacrylate, FOSM (BOC Sciences), was recrystallized twice from methanol (Fisher) before use. The initiator 2,2′-azobis(isobutyronitrile) (AIBN) (SigmaAldrich) was purified by recrystallization from methanol. The RAFT agent 2-cyano-2-propyl benzodithioate (CPBDT), 1,4-dioxane, α,α,αtrifluorotoluene (TFT), and diethyl ether were obtained from SigmaAldrich and used without further purification. 1,1,2-Trichlorotrifluoroethane (CFC-113) was obtained from Fisher Scientific and used as received. RAFT Polymerization. Triblock poly(FOSM)−poly(DMA)− poly(FOSM) (FxDyFz, where x, y, and z represent the molecular weight of each block in units of kg/mol) was synthesized by sequential RAFT polymerizations of the three blocks. The polymerization procedure is shown in Scheme 1. The poly(FOSM), PF, homopolymer was not completely soluble in any common organic solvent, but it was soluble in trifluorotoluene (TFT), which was used as the solvent for the polymerization of the first PF block, carried out at 80 °C. AIBN was used as the initiator and CPBDT as the RAFT agent. The second and third blocks were synthesized at 80 °C using AIBN and CPBDT and a 3:1 v/v dioxane/TFT mixed solvent to prevent the polymers from precipitation. The triblock copolymers were asymmetric, which was a consequence of the synthesis by three sequential RAFT polymerizations. The molecular weight of the first PF end-block was held constant at 16.0 kg/mol for each block copolymer produced, but the Received: October 18, 2014 Revised: January 2, 2015 Published: January 16, 2015 645

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Macromolecules Scheme 1. Synthesis Route for FxDyFz Triblock Copolymers Using Sequential RAFT Polymerization

Table 1. Homopolymers and Triblock Copolymers Synthesized by RAFT Polymerization block Mn (kg/mol) sample

Mn (kg/mol)

PF PD FDr F16D20F25 F16D38F19 F16D50F19 F16D71F15

16.0 24.8 117 61.4 72.8 85.6 103

F

16.0 16.0 16.0 16.0

D

19.8 37.6 50.4 71.3

DMA

FOSM

F

mol %

ϕg

mol %

ϕg

25.6 19.2 19.2 15.3

0 100 89.3 74.3 87.0 90.0 93.5

0 1.00 0.685 0.430 0.636 0.701 0.785

100 0 10.7 25.7 13.0 10.0 6.5

100 0 0.315 0.579 0.364 0.299 0.215

Tg1a (°C)

Tcpb (°C)

ΔHchc (J/g)

Tccd (°C)

74

101

2.67

95

Tg2e (°C) 117

107 81 70 69 70

98 97 98 98

112 120 120 117

Tdf (°C) 188 375 337 292 300 304 305

a

Glass transition temperature of PF block (DSC). bEndothermic transition temperature upon heating (DSC). cEnthalpy of endothermic transition upon heating (DSC). dExothermic transition temperature upon cooling (DSC). eGlass transition temperature of PD block (DSC). fDegradation temperature defined as the lowest temperature peak in the derivative TGA curve. gϕ = volume fraction. molecular weight of the other PF end-block varied from 15 to 26 kg/ mol (see Table 1). The composition of the FDF triblock copolymer was varied mainly by changing the molecular weight of the poly(DMA), PD, midblock between 20 and 71 kg/mol. The compositions and thermal properties of the four triblock copolymers studied are summarized in Table 1. A PF homopolymer, a PD homopolymer, and a random copolymer of FOSM and DMA random copolymer containing 10.7 mol % FOSM, FDr, were also synthesized for comparison purposes. The FDr synthesis followed the procedure described by Tian et al.8 The molecular weights listed in Table 1 were calculated from 1H NMR. Size exclusion chromatography was not used because of the insolubility of PF and the block copolymers in commonly used solvents. Example Synthesis of Triblock Copolymer (F16D20F25). For the first PF block, FOSM monomer (20 g, 31.3 mmol), the CPBDT RAFT agent (346 mg, 1.6 mmol), and AIBN initiator (51 mg, 0.3 mmol) were dissolved in 100 mL of TFT, and the solution was stirred and purged with N2 for 1 h at room temperature. The polymerization was run for 12 h at 80 °C, and then the solution was cooled to 0 °C to terminate the reaction. The polymer was precipitated and washed with excess diethyl ether and then dried under vacuum at room temperature for 24 h. 13.2 g of PF powder with a molecular weight of 16.0 kg/mol was obtained, with a yield of 66%.

A FOSM−DMA diblock copolymer, denoted as F16D20, was synthesized by dissolving the PF sample with a CPBDT chain end from the preceding paragraph (3 g, 0.1 mmol), denoted as F16, DMA monomer (4.1 g, 41.4 mmol), and AIBN (7 mg, 0.04 mmol) in a mixed solvent of 20 mL of TFT and 20 mL of dioxane. After purging with N2 for 1 h at room temperature, the polymerization was run for 24 h at 80 °C. The reaction was terminated by cooling the solution to 0 °C. The diblock polymer was precipitated and washed with diethyl ether and then dried under vacuum at room temperature for 24 h. 6.3 g of F16D20 diblock copolymer powder was obtained (yield = 89%). Diblock copolymers F16D38, F16D50, and F16D71 were synthesized by the same method, but using different amounts of the DMA comonomer or polymerization time. The triblock copolymer F16D20F25 was synthesized from the diblock copolymer F16D20. F16D20, which had a RAFT agent chain end (5 g, 0.12 mmol), FOSM (5.5 g, 8.6 mmol), and AIBN (7 mg, 0.04 mmol) were dissolved in mixed solvent of 40 mL of TFT and 30 mL of dioxane. The solution was purged with N2 while stirring at room temperature for 1 h. The solution was then heated to 80 °C to initiate the polymerization. After 48 h, the solution was cooled to 0 °C and the triblock copolymer was precipitated in diethyl ether. The triblock copolymer was washed with diethyl ether and dried at room temperature, and 7.0 g of triblock copolymer powder F16D20F25 was isolated (yield = 67%). Triblock copolymers F16D38F19, F16D50F19, and 646

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F16D71F15 were similarly synthesized from the corresponding diblock copolymers (see Table 1). Polymer Characterization. 1H NMR analysis was conducted with a Varian Mercury 300 MHz spectrometer using 1,1,2-trichloro-1,2,2trifluoroethane (CFC-113) as the solvent. Deuterated acetone was added to “lock” the magnetic field.13 Thermal properties were measured with a TA Instruments Q2000 differential scanning calorimeter (DSC) using heating and cooling rates of 10 °C/min. The glass transition temperature was defined as the midpoint of the change in heat capacity associated with the glass transition. The transition temperature for first-order endothermic or exothermic transitions was defined at the peak of the transition, and the enthalpy of the transition was calculated from the area under the endotherm or exotherm. Thermal stability was determined with a TA Instruments Q50 thermogravimetric analyzer (TGA) using a nitrogen atmosphere and a heating rate of 20 °C/min. The thermal degradation temperature was defined as the peak in the derivative of the first mass loss process. The thermal decomposition products at 190 °C were analyzed by combined gas chromatography/mass spectroscopy (GC/MS) measurements using a Agilent 7890-A GC/MS system with a dynamic thermal desorption device,14 DTD, mounted on the GC injection port. The DTD was used to thermally extract the decomposition products from samples and inject them directly into the capillary column of the GC. Equilibrium swelling measurements were made using compressionmolded films. A preweighed specimen was immersed into deionized water, and the gel was allowed to equilibrate at room temperature. During this process, the gel was taken out at various times and weighed to within ±1 mg to determine the water concentration. The water sorption of the triblock copolymer hydrogels is expressed as a swelling ratio, S, defined as the mass of water absorbed per mass dry polymer; i.e., S = 2 corresponds to 2 g water/g polymer. Wide-angle X-ray diffraction (WAXD) was carried out with a Bruker X-ray diffractometer using Cu Kα radiation (λ = 0.154 nm) and a generator voltage of 40 kV and current of 40 mA. Small-angle X-ray scattering (SAXS) measurements were performed with a Rigaku rotating anode X-ray generator operating at 4 kW with Cu Kα radiation (λ = 0.154 nm) and a Ni filter. This was interfaced with a Bruker AXS SAXS system, including a Hi-Star area detector. The sample geometry used allowed for data collection over a range of scattering vector q = 4π sin θ/λ, where θ is one-half the scattering angle, from 0.1 to 3.3 nm−1. That q-range corresponds to sizes in real space, d = 2π/q, of 1.9−63 nm. Compression-molded triblock polymer films were annealed at 170 °C under vacuum for 2 days before the structure and morphology characterization. Measurements were made on dry-annealed, water-swollen, and dehydrated films. The hydrated samples were sandwiched between Kapton polyimide (DuPont de Nemours Co.) films to prevent the sample from dehydrating during the experiment. All the samples were measured at ambient temperature. Transmission electron microscopy (TEM) was carried out with a Jeol JEM-1230 transmission electron microscope using a 120 kV illumination system. Thermally annealed samples for TEM were microtomed into ultrathin (∼100 nm) films that were then stained by osmium tetroxide for 12 h prior to TEM imaging. The pore morphology of dehydrated gels was observed with a Jeol JSM-6700F field emission scanning electron microscope (FEG-SEM) using an accelerating voltage of 5 kV. Gel samples that reached equilibrium in deionized water were freeze-dried for 48 h, and the specimens were cryo-fractured in liquid nitrogen. The fracture surface was coated with platinum and observed with the FEG-SEM at room temperature. Dynamic shear rheological measurements were made with a TA Instruments ARES-G2 rheometer equipped with 8 mm parallel plates and a water trap that maintained the sample hydration. Dynamic experiments used a frequency range of 0.1−100 rad/s and a strain within the linear response region, which was determined from a strain sweep.

Article

RESULTS AND DISCUSSION Thermal Analysis. The DSC thermograms of homopolymers PF, PD, and copolymers F16D50F19 and FDr are compared in Figure 1. The PF exhibited a glass transition

Figure 1. DSC heating thermograms polymer for (a) PF, (b) PF/PD 1:1 (w/w) blend, (c) PD, (d) F16D50F19, and (e) FDr. Arrows for curves b and d indicate Tg’s. Inset is a magnification of the block copolymer data (curve d).

(Tg1) at 74 °C and an endotherm at 101 °C (ΔH = 2.67 J/g) upon heating from room temperature (curve a in Figure 1). The PF forms a liquid crystalline mesophase above Tg, and the endotherm at 101 °C represents the clearing point (Tcp), i.e., the liquid crystal to isotropic transition. The latter result is consistent with the report by Wang et al.15 that a similar polymer, poly(2-N-methylperfluorohexanesulfamido)ethyl methacrylate) with a C6 perfluorinated side chain, formed a smectic A liquid crystalline phase with Tcp = 125 °C and ΔHcp = 3.0 J/g. The PD homopolymer was completely amorphous with a glass transition temperature (Tg2) of 117 °C (curve c in Figure 1). The random copolymer FDr was also amorphous. Although SAXS analysis indicates that this polymer exhibited microphase separation of FOSM nanodomains, the DSC curve only resolves a single Tg at 107 °C (curve e in Figure 1). A 50/50 (w/w) blend of the two homopolymers (curve b in Figure 1) exhibited two Tg’s at 73 and 120 °C, which corresponded to the Tg’s of the individual components, and the clearing point of the mesophase in the PF at 101 °C. Those results for the blend indicated that the PF and PD were completely immiscible. The DSC heating thermogram of F16D50F19 (curve d in Figure 1 and the inset) showed two Tg’s, at 69 and 120 °C (also see Figure S1 of the Supporting Information, which more clearly shows the Tg of the PF phase), that corresponded to the two homopolymer Tgs, though these were much broader than for either the homopolymers or the immiscible blend of the two homopolymers. The broadening of the glass transitions was a consequence of the nanometer-scale microphase separation of the two polymer phases in the block copolymer, which will be discussed later in this paper. The F16D50F19 block copolymer also exhibited the same order−disorder transition of a liquid crystal phase, Tcp, as the PF homopolymer, though the temperature of the transition was several degrees lower (98 °C) than for the PF homopolymer and the enthalpy of the transition was too small to accurately measure. The thermal transitions for the four triblock copolymers are summarized in Table 1. The DMA concentration varied from about 74 to 94 mol %, and with the exception of the block copolymer with the lowest DMA concentration, F16D20F25, the 647

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Macromolecules two Tg’s were close to those of the homopolymers. The two Tg’s for the F16D20F25 increased for the poly(FOSM) blocks and decreased for the poly(DMA) block, which indicates there may have been some phase mixing in that sample. Tcp was observed in the DSC scans for all of the block copolymers, but it was very weak. The TGA data for the PF, PD, their blend, and F16D38F19 are shown in Figure 2, and the temperatures of the peak of the

for direct comparison of the properties of hydrophobically modified block copolymer hydrogels to random copolymer hydrogels with the same composition that were studied previously.7−10 Water Sorption. The block copolymers were composed of a water-soluble PD midblock and hydrophobic PF end-blocks, and they formed water-insoluble hydrogels due to a selfassembled, microphase-separated morphology. The microstructure is described in more detail in the next section of this paper. Water absorption isotherms at 23 °C for the four block copolymers and the random copolymer, FDr, are shown in Figure 3. In general, equilibrium swelling was achieved in

Figure 2. TGA (solid lines) and DTG (dashed lines) curves of (a) PF, (b) PF/PD blend 1:1 (w/w), (c) F16D38F19, and (d) PD.

derivative for the first mass loss process, Td, for all the polymers used in this study are summarized in Table 1. The least stable of the polymers shown in Figure 2 was PF, which had a degradation process characterized by three distinct mass loss processes, at about 200, 320, and 370 °C, which is similar to the degradation behavior of poly(methyl methacrylate), PMMA.16 Based on the degradation assignments for PMMA given by Manring and co-workers,17−20 the degradation processes for PF were assigned to β-scission at the vinyl end-groups due to radical transfer produced by termination and/or by disproportionation during polymerization, hemolytic cleavage due to an inductive effect of neighboring ester groups at head-to-head linkages, and random chain scission, respectively. The PD was stable to about 425 °C, and a blend of the two homopolymers exhibited four distinct mass loss processes that correspond to those observed for PF and PD. The mass loss below 100 °C in the PD sample (and in the block copolymer) was due to water as a result of either incomplete drying of the samples or absorption of water during the transfer of the sample to the TGA. The F16D38F19 block copolymer was stable to higher temperature than either the PF or the PF/PD blend. Only a single, though broad mass loss process began at ∼250 °C, with the maximum mass loss rate occurring at ∼380 °C. The main degradation temperatures for all the block copolymers were similar (see Table 1), though the amount of degradation products decreased as the more stable DMA concentration of the block copolymer increased. No attempt was made to identify the degradation products of the copolymers. However, one potential byproduct is perfluorooctane derivative, which is currently considered an emerging contaminant,21 or a potential threat to human health and the environment.22 As such, the polymers described herein may not be suitable for a commercial product. However, the FOSM and FOSA monomers have been previously used in our research and by other groups interested in hydrophobic modification of hydrogels prior to their being considered as contaminants by the EPA. In the current paper, we used FOSM

Figure 3. Swelling ratio as a function of time for triblock copolymers (a) F16D20F25, (b) F16D38F19, (c) F16D50F19, (d) F16D71F15, and (e) a random copolymer, FDr. The numbers to the right of the data are the equilibrium swelling ratios.

about 1 day, and the value of the equilibrium swelling ratio increased exponentially with the mass fraction of DMA (xDMA) in the block copolymer (Figure 4). The swelling ratio for the

Figure 4. Equilibrium swelling ratio vs DMA mass fraction for block and random copolymers. The dashed curve is a least-squares exponential fit of the block copolymer data, and the dotted line is an exponential fit of the random copolymer data.

random copolymer with the same DMA concentration as the F16D50F19 block copolymer was about 30% lower than for the block copolymer (see Figure 4). The swelling ratio data for DMA/FOSA random copolymers from Tian et al.8 are also plotted in Figure 4. Those results indicate that the swelling of the block copolymer and random copolymer are similar at the lower DMA mass fractions, but the S values for the xDMA > ∼0.4 are lower for the random copolymers than for the block copolymers. An exponential function (S = a exp(bxDMA)) fit the equilibrium data well (see Figure 4), but the coefficient b in the exponential was greater than for the random copolymer, which 648

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Macromolecules corresponds to a greater increase in swelling per mass fraction DMA. Note that there is no physical significance to the exponential fit used in Figure 4. Other nonlinear fits could have been used and would have produced the same conclusions with regard to the effect of the copolymer architecture on the swelling behavior. An exponential fit was used merely for convenience, since it only required one fitting parameter. The differences in S for the random and block copolymers were much greater than the experimental error (cf. curves c and e in Figure 3) and indicate that the hydrogen bonding sites in the DMA, which are responsible for the water absorption, are more available to water in the block copolymer than in the random copolymer. That might be expected because of the differences in the microstructure for the block and random copolymers. As is discussed in the next section, the block copolymers exhibited cylindrical or lamellar nanodomains ∼20 nm thick as opposed to the approximately 6 nm diameter core−shell nanodomains microstructure of the random copolymers.9 SANS studies of random DMA/FOSA copolymers indicated that FOSA constituted the core of about 3 nm diameter core−shell nanodomains, which were surrounded by a water-depleted shell of DMA. The swelling data in Figure 4 are consistent with the conclusion that the DMA layer in contact with the hydrophobic nanodomains cannot effectively hydrogen bond with water. The nonlinearity of the swelling data in Figure 4 for the random copolymers may be explained by a nonlinear decrease in the relative ratio of water-depleted DMA at the interface of the hydrophobic nanodomains and the waterswollen DMA in the bulk as DMA concentration increases. A similar argument could be made for the nonlinear swelling data for the block copolymers if the DMA adjacent to the PF nanodomains was similarly water depleted, but that was not assessed in this study. Another explanation for the nonlinearity is that there are more than one type of water present in these materials, e.g., bonded water and free water, and the relative amounts change in a nonlinear fashion with composition. For spherical domains, the ratio of the interfacial area, Asp, to the lamellae volume, Vsp, is proportional to the inverse of the diameter of the spheres, Dsp−1. For lamellae or cylindrical domains, A/V scales with inverse lamellae thickness, tlam−1, or the cylinder diameter, dcyl−1, respectively. Therefore, the ratio of the interfacial area of the FA nanophases of the block and random copolymers compared in Figure 4 is ∼1/10; i.e., the nanodomains in the block copolymer have an order of magnitude less interfacial area between the FA and PD phases than in the random copolymer. That explains the lower swelling ratio for the random copolymers; i.e., more interfacial area means more water depleted DMA. Photographs of the water-swollen triblock copolymer hydrogels are shown in Figure 5a. The swelling was isotropic in the plane shown in this figure, and the network nature of the hydrogels is apparent, even though there are no covalent crosslinks in those materials. As will be discussed below, the PD blocks are each connected to two PF end-blocks that reside in different unswollen nanodomains, which prevents solubility and suppresses mobility of the PD chains. As such, the DMA nanodomains swell, but the chain segments cannot flow; i.e., there was no permanent deformation of the block copolymer as a result of swelling and deswelling. In contrast, the FxDy diblock precursors did not behave as cross-linked networks (Figure 5b). When immersed in water, the shape of the film was lost and the film fragmented into highly water-swollen pieces. In that case, the DMA block was connected to only one DF chain and one

Figure 5. Photographs of (a) triblock copolymer films and hydrogels and (b) diblock copolymer films, dry and in immersed in water.

DF nanodomain, so there was no restraint on the mobility of the chains as there was for the triblock copolymer where each DMA block was restrained by two unswollen DF nanodomains. Structure and Morphology of Triblock Copolymers and Hydrogels. Figure 6 shows the WAXD data for the dry homopolymers and triblock copolymers annealed at 170 °C for 2 days and the water-swollen hydrogels prepared from the triblock copolymers. The PD homopolymer was amorphous and exhibited two broad diffraction peaks at q = 8.4 and 17 nm−1, which correspond to average spacings of d = 0.74 and 0.37 nm, respectively. The origin of those spacings is most likely due to interchain correlations involving the PD backbone and intrachain correlations between the amide side groups. The PF homopolymer showed two relatively sharp peaks at q = 2.2 and 4.1 nm−1 and a broad peak at q = 12 nm−1. The two sharper peaks correspond to Bragg spacings of d = 2.9 and 1.5 nm, respectively, which are the first- and second-order reflections of the lamellar organization of a smectic phase of the perfluoroalkyl groups that are oriented normal to the chain backbone (Figure 7). That interpretation is consistent with the smectic structure reported for poly(fluoroalkyl acrylates) by Wang et al.,15 Copart et al.,23 and Honda et al.24,25 The broad, amorphous peak at q = 12 nm−1 (d ∼ 0.53 nm) represents the average separation of the fluoroalkyl chains. The identification of a smectic structure is also consistent with the DSC results discussed earlier in the paper, which indicated a very low heat of transition that was attributed to the liquid crystal clearing point. A weak amorphous halo at q ∼ 26 nm−1 (d = 0.24 nm) is probably due to the average separation of the PF backbone chains. The WAXD for the triblock copolymers showed the same two relatively sharp peaks as were observed in PF, corresponding to d ∼ 2.9 and 1.5 nm, though the intensities were weaker due to the lower concentration of PF in the copolymers. Similarly, the peak for the fluoroalkyl chain packing corresponding to d ∼ 0.53 nm was broader than in the PF homopolymer due to overlap with the PD diffraction peaks at q = 8.4 and 12 nm−1. A new diffraction peak for the dry triblock copolymers is seen at q < 1.0 nm−1 in Figure 6a, which is due to the block copolymer microstructure. That peak is better resolved by small-angle X-ray scattering (SAXS) and is described in more detail in the SAXS discussion below. The first reflection in the WAXD of the perfluoralkyl smectic phase persisted for the swollen triblock copolymer hydrogels 649

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Figure 6. Room temperature WAXD curves for triblock copolymers: (a) annealed dry copolymer and (b) hydrogel swollen with water to equilibrium.

that may be the additional stress on the PF chains due to the osmotic swelling of the connected PD chains. The SAXS curves of compression molded and annealed triblock copolymer samples (Figure 8) show an organized

Figure 7. Schematic of smectic phase structure of the fluoroalkyl chains.

(Figure 6b), though the intensity was greatly diminished due to the dilution of the polymer by water. The second reflection of the mesophase also remained, though it was very weak and not clearly resolved for the triblock copolymer with the lowest PF concentration. The 0.53 nm spacing of the fluoroacrylate side chains (peak at q = 12 nm−1) also persisted for the hydrogels. Those results were not unexpected, since the hydrophobic PF nanodomains were not swollen by water. The equilibrium solubility of water in PF was insignificant, 0.15 ± 0.21%, as determined gravimetrically by immersing PF films in DI water for 9 days and measuring the mass increase. The broad peaks at 8.4 and 17 nm−1 observed by WAXD for the PD homopolymer and the dry block copolymers were not observed for the hydrogels. Instead, a single, but very broad amorphous halo at q = 20 nm−1 indicated a distance in real space of d ∼ 0.31 nm. The origin of that correlation is not clear. Water only swells the PD phase, and the concentration of water relative to polymer in the swollen triblock copolymers was ∼200−800% depending on the composition of the block copolymer. One would expect that in such a highly swollen sample there would be no clear correlations between the polymer chains, and even if there were, the intensity of the scattering would be too small to resolve. The only plausible explanations are that the d ∼ 0.31 nm correlation corresponds to an increase in the average separation of the backbone chains in the water-resistant PF phase or a decrease in the separation of the amide side groups in the water-swollen PD phase. The latter explanation seems counterintuitive in that it is not clear why the separation of the amide groups would become closer when the PD phase was swollen with water. The alternative assignment, i.e., the correlation, corresponds to the separation of the PF chains would indicate a ∼30% decrease in the packing of the PF chains when the hydrogel is swollen. One reason for

Figure 8. SAXS curves at room temperature for the triblock copolymers.

microstructure with a characteristic size of ∼30 nm. Four distinct low-q reflections are evident in the data for F16D20F25, which had the highest volume fraction of FOSM, ϕPF = 0.579. The ratio of q-values (q/q*, where q* is the value for the first peak) for the peaks in the SAXS curves of F16D20F25 scaled as ∼1:2:4:5, which is characteristic of an alternating lamellae microstructure. The third-order peak is missing, which is due to cancelation of the peak in that structure factor by interference from the form factor.26 The first SAXS peak, q* = 0.193 nm−1, corresponds to a lamellae periodicity dq* = 32 nm. The broad peak at q = 2.1 nm−1 corresponds to the 3.0 nm smectic phase also observed in the WAXD results. A transmission electron micrograph of the dry F16D20F25 triblock copolymer (Figure 9) is consistent with an alternating lamellar microstructure. The darker phase in the micrograph is the PF. Actually, the morphology of only one plane of the block copolymer shown in Figure 9 cannot distinguish between a LAM or a HEX microstructure, but the SAXS data in Figure 8 clearly indicate a LAM morphology for this polymer. The PF lamellae (dark) thickness measured from Figure 9 is ∼22 nm thick, and the PD lamellae (light) are ∼14 nm. The period, ∼36 nm, is ∼10% larger than that indicated by SAXS, but the interface between the two nanophases in Figure 9 is rough, which makes it difficult to get an accurate measurement of the lamellae periodicity. The volume fraction of PF (ϕPF) 650

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LAM. Therefore, the experimental SAXS value of dq* = 36 nm corresponds to either a lamellar microstructure d* = 36 nm or hexagonally packed PF cylindrical nanodomains with an interdomain spacing of a = (4/3)1/2d* = 41.6 nm and a cylinder radius of r0 = d*(√3ϕPF/2π)1/2 = 11.4 nm. For F16D50F19 (ϕPF = 0.299), the phase diagram reported by Matsen27 indicates the PF should form either a disordered, spherical, or HEX phase. Since neither spherical nor disordered microstructures correspond to the SAXS curve in Figure 8, F16D50F19 most likely forms a HEX phase with d* = 32 nm, a = 37 nm, and r0 = 9.2 nm. The SAXS results for F16D71F15 (ϕPF = 0.215) were similar to those for F16D50F19. Two reflections were observed, but it was difficult to resolve the peak position of the second-order reflection. As with the data for F16D50F19, the wavenumber (q) where q/q* = 2 is marked in red type in Figure 8, and that appears to be the best estimate for the two peaks in the structure factor. The SAXS data for F16D71F15 could be for a HEX or LAM microstructure, but the predictions of the Matsen27 phase diagram favor an assignment of HEX. Forming the hydrogel by swelling F16D20F25 with water changed the microstructure from LAM to HEX (Figure 10).

Figure 9. Transmission electron micrograph of the lamellar microstructure of dry F16D20F25.

calculated from the photo in Figure 9 is 0.611, which is about 5% higher than the value of ϕPF calculated from the composition (Table 1). On the whole, the SAXS and TEM results are pretty much in agreement. The assignments of the microphase texture for the other three block copolymers shown in Figure 8 are uncertain. “The following discussion provides some general remarks with regard to the possible microstructures of these block copolymers and how they compare with mean field calculations by Matsen.27 However, the conclusions should be considered very tentative at this time. As was pointed out by one of the reviewers of the paper, these block copolymers were, in general, strongly segregated, polydisperse, and asymmetric. Those properties can produce microstructure textures more complex than the LAM and HEX structures that were considered in the following discussion.” F16D38F19 and F16D50F19, which had FOSM volume fractions (ϕPF) of 0.364 and 0.299, respectively, only exhibit two clear low angle reflections, but they were broad and the second scattering peak overlapped with the first, which prevented definitively assigning the peak positions. Of those two block copolymers, the second peak for F16D38F19 was cleaner, and it appeared that the ratios of q/q* were 1:2, which is consistent with the structure factor for alternating lamellar (LAM) or hexagonally packed cylinders (HPC) if the structure factor were missing the q/q* = √3. F16D50F19 also exhibited two reflections, but in this case there was too much overlap of the two to clearly identify the value of q/q*. A red arrow is shown in Figure 8 for the peak position if q/q* = 2, and no other ratio in the region where the peak looks like it would appear would correspond to a typical block copolymer morphology. As a result, the mesophase textures of F16D38F19 and F16D50F19 are most likely LAM or HEX. Because of the fluorinated and hydrophilic components of the block copolymer, one would expect these materials to fall into the strong segregation regime of phase separation, i.e., χN > 10,28 where χ is the interaction parameter for the two homopolymers and N is the number of repeat units. Mean field theory (MFT) calculations by Matsen27 predict the phase diagram in the strong segregation regime for ABA triblock copolymers with molecular weight asymmetry of the two A blocks. For F16D38F19 and F16D50F19, the asymmetry factor is τ = 0.45, where τ = NA1/NA, where NA1 is the number of segments in the shorter A block and NA is the total number of A segments.27 Allowing for some error in the phase boundaries proposed by Matsen27 for high χN F16D38F19 (ϕPF = 0.364) would lie close to the HEX ↔ LAM phase boundary, but there are not enough reflections to distinguish between HEX and

Figure 10. SAXS curves for F16D20F25: (a) dry, annealed sample, (b) fully hydrated sample, and (c) rehydrated sample. The arrows denote the SAXS reflections for the hydrogel (b).

The LAM texture for the dry block copolymer had a first-order SAXS peak corresponding to dq* = 31 nm. Note that the values of dq* for the dry F16D20F25 in Figures 8 and 10 differ by about 3%, which is probably due to some absorption of water by the hygroscopic sample and differences in the amount of absorbed water in the two experiments shown in Figures 8 and 10. As with the scattering data in Figure 8, the values for q/q* for the F16D20F25 SAXS peaks in Figure 10 were ∼1:2:4:5, which are consistent with a LAM microstructure. The hydrogel formed from F16D20F25, however, exhibited SAXS reflections at q = 0.158, 0.245, 0.332, and 0.410 (see arrows in Figure 10), which give values of q/q* of approximately 1:√3:2:√7, which are characteristic of a HEX microstructure. The second reflection deviated from √3q* by about 12%, the third reflection deviated from 2q* by 5%, and the fourth reflection deviated from √7q* by about 2%, which may be a consequence of polydispersity of the molecular weights of the three blocks. Because of the absence of a suitable solvent for GPC measurements, the polydispersity of each sequential RAFT polymerization was not determined. However, one would normally expect that even if each separate polymerization step gave a narrow molecular weight distribution (MWD), the additive broadening of the MWD of the triblock copolymers should be broader than one would expect from the RAFT polymerization of a homopolymer or even a triblock copolymer 651

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for the random copolymer hydrogel and the oriented open cell pore structure for the block copolymer hydrogel in Figure 11 to the microstructure of the two polymers. That is, the FDr gel microstructure consists of spherical nanodomains rich in FOSM9, while the block copolymer hydrogel probably was probably an anisotropic HEX texture. The characteristic size of the nanodomains in the two polymers differed by an order of magnitude: ∼3 nm for the FOSM nanodomains in the FDr hydrogel and ∼30 nm for the F16D50F19 hydrogel. In both cases, PD surrounds the hydrophobic nanodomains and is the only component that swells. The pore sizes shown in Figure 11 are 2−3 orders of magnitude larger than the nanodomains in the two hydrogels, so it is not obvious how the microstructure would affect the macrostructure of the hydrogels. Nor is it clear whether the structure of the hydrogel was maintained intact during the freeze-drying operation, which brings into question whether the differences shown in the two photos in Figure 11 accurately represent the gel structures. The ability to develop different types of gel structures, however, could be useful in membrane applications of these hydrogels, so this is a subject worthy of further studyusing small-angle neutron scattering to directly image the pore structure. Rheology. The frequency dependence of the dynamic storage (G′) and loss (G″) shear moduli of the block copolymer hydrogels at room temperature are shown in Figure 12. For each hydrogel, G′ > G″ over the entire frequency range, which is consistent with the formation of network in these materials. The elasticity arises from the physical network formed by the microphase separation of the block copolymers. Water only swells the PD nanodomains, and each PF block bridges two PD chains such that the PF nanodomains provide the physical cross-links that form the hydrogel network. The weak frequency dependence of G′ and G″ in Figure 12 is due to the viscoelastic behavior of the block copolymer hydrogels that arises from the nonpermanent nature of the network. For the three hydrogels with the highest concentrations of PD (F16D38F19, F16D50F19, and F16D71F15), tan δ was ∼0.1, which indicates that those gels were nearly elastic in their mechanical behavior. In contrast, F16D20F25, which had the lowest PD concentration (and lowest swelling ratio), exhibited a much higher viscous component in the mechanical response (tan δ ∼ 0.1−0.3). In general, G′ decreased with increasing PD content, which is expected since increasing the PD concentration increased the water absorption. The effective cross-link density, nc, of the hydrogels was calculated from eq 129

prepared with a difunctional RAFT catalyst. For a HEX microstructure, the SAXS data in Figure 10 for the F16D20F25 hydrogel (ϕPF = 0.313) indicate an interdomain spacing cylinder of a = 45 nm and a PF cylinder diameter of 7.0 nm. The SAXS peak for the smectic ordering of the fluorocarbon side chain of PF persisted in water-swollen F16D20F25. That is not surprising, since the water only swells the PD phase. The bilayer structure, however of the smectic phase expanded by 0.15 nm in the hydrogel. Although a definitive explanation of the expansion of the smectic spacing is not obvious, the result is consistent with the WAXD data discussed earlier in this paper of a 30% decrease in the packing of the chains in the PF nanodomains upon swelling. Since the PF phase was not swollen by water, the origin of that decrease in packing density of the PF nanodomains must be due to the LAM → HEX transition. The porous structures of the random copolymer hydrogel and the F16D50F19 triblock copolymer hydrogel are compared in the SEM photos of freeze-dried hydrogels shown in Figure 11.

Figure 11. SEM images of hydrogel pore structure of (a) FDr hydrogel and (b) F16D50F19 hydrogel (inset to each photo is 3× the magnification of the larger photo).

Although the two copolymers had nearly identical compositions, ∼70 vol % DMA, the triblock copolymer absorbed nearly 60% more water than the random copolymer (see Figure 4). The FDr hydrogel had a spherical, closed pore structure with ∼2−5 μm pores surrounded by dense walls of ∼2−5 μm (Figure 11a). In contrast, the block copolymer hydrogel had an open cell, sponge-like morphology with a relatively large macroporous structure of ∼20 μm oriented and interconnected channels surrounded by ∼2−5 μm thick walls (Figure 11b). It is tempting to attribute the spherical closed cell pore structure

Figure 12. Linear viscoelastic properties of the triblock copolymer hydrogels as a function of frequency: (a) storage modulus (G′) and (b) loss modulus (G″). The strains used for the experiments were 0.03% for F16D20F25, 0.1% for F16D38F19, and 0.5% for F16D50F19 and F16D71F15. 652

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Macromolecules ⎛ 2⎞ G = ⎜1 − ⎟ncRTφ2 2/3 f⎠ ⎝

formed from random copolymers of DMA and FOSA. For example, G′ at room temperature and ω = 100 rad/s for FDr and F16D50F19 (each is ∼90 mol % PD) were 200 kPa11 and 22 kPa, respectivelya difference of an order of magnitude. Using eq 2 to account for the differences in swelling ratio of the two hydrogels, the effective cross-link density of the random copolymer is about 5 times that of the block copolymer (calculated from eq 1) of the same composition. The higher apparent cross-link density for the random copolymer hydrogel may be a consequence of the different morphology and size of the nanodomains cross-links in each system. As was discussed earlier in this paper, the water-swollen block copolymers were viscoelastic solids. As such, their mechanical response is expected to be dependent upon the frequency and amplitude of the oscillatory shear. Figure 13 shows the results for three cycles where the F16D50F19 hydrogel was first sheared for about 200 s at ω = 1 rad/s and a strain amplitude, γ, of 1%, which was within the linear response region, and then γ was increased stepwise from 0.01 to 100%. When γ was within the linear region, G′ and G″ were insensitive to the strain amplitude. Under these conditions, the swollen block copolymer was nearly an elastic solid with G′ = 19 kPa and tan δ = 0.05. When γ > 1%, G′ decreased and a crossover where G″ became greater than G′ occurred at about γ = 3%. The crossover where G″ > G′ was due to a strain-induced solid to liquid transition of the hydrogel. That is, at high strain amplitude the hydrogel behaved as a viscoelastic liquid. The transition from elastic solid to viscoelastic behavior indicates a transformation of the material structure from a cross-linked network to a solution, which was possible because of the physical nature of the network. When γ was returned to a value within the linear response region, however, the material became an elastic solid again and G′ and G″ returned to their original values within 15−20 s after removing the nonlinear strain. This behavior was repeated three times in Figure 13, and the response of the hydrogel and the solid to liquid transition were reversible.

(1)

where f is the functionality of the cross-links, which was assumed to be 4, R and T are the gas constant and absolute temperature, respectively, and ϕ2 is the volume fraction of the cross-linked polymer in the hydrogel. ϕ2 was calculated from eq 2 −1 ⎡ SρP ⎤ ⎥ φ2 = ⎢1 + ⎢⎣ ρS ⎥⎦

(2)

where S is the swelling ratio, ρP is the polymer mass density, and ρS is the solvent mass density, in this case water (1.00 g/ mL). The nc values calculated from eqs 1 and 2 are summarized in Table 2. Note that there is no justification for assuming f = 4, Table 2. Cross-Link Density and Viscoelastic Properties of Physical Gels at Room Temperature

a

gels

ρP (g/cm3)

ϕ2

nca (mol/m3)

G′a (kPa)

tan δa

F16D20F25 F16D38F19 F16D50F19 F16D71F15

1.30 1.29 1.24 1.25

0.50 0.32 0.23 0.13

30−55 50−65 34−41 11−13

19−43 27−38 16−22 3.2−4.3

0.19−0.31 0.069−0.10 0.028−0.14 0.038−0.093

Range of values from ω = 1 to 100 rad/s.

since the functionality of the cross-link in a microphaseseparated system is not clear. However, since the microstructures of the four block copolymers were similar, choosing any number for f provides a convenient way to at least provide a qualitative comparison between those materials. Since these were physical hydrogels and G′ = G′(T,ω), nc is also a function of frequency and temperature. As a consequence, a range of nc calculated from the data in Figure 12 are reported in Table 2. In general, nc increased with increasing PF concentration. The value of the storage modulus G′ for the four block copolymer hydrogels at room temperature ranged from 3.2 to 43 kPa and generally increased with increasing PF concentration and frequency. These values appear to be on the high side compared with other conventional covalent hydrogels and other neutral block copolymer hydrogels reported in the literature, where G′ is ∼1−10 kPa. However, the modulus values are much lower than those of the physical hydrogels

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CONCLUSIONS Triblock ABA copolymers (bcp) where the A blocks were a hydrophobic polymer (PF) and the B blocks were a watersoluble polymer (PD) were synthesized by three sequential RAFT polymerizations. As a result of the separate polymerizations of the two A blocks, their molecular weights were

Figure 13. Dependence of G′ and G″ of the F16D50F19 hydrogel for a time sweep and then three consecutive strain sweep/time sweep cycles at room temperature. The frequency was held constant at ω = 1 rad/s. The time sweeps were run with a fixed strain of 0.5%, which was within the linear response region. The strain sweeps were run between 0.01 and 100%. For each discrete strain step the stress was allowed to reach steady state before the strain was increased. The small changes that occurred at the beginning of each strain sweep were due to a transient due to changing the strain amplitude. 653

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(8) Tian, J.; Seery, T. A. P.; Weiss, R. A. Macromolecules 2004, 37, 9994. (9) Tian, J.; Seery, T. A. P.; Ho, D. L.; Weiss, R. A. Macromolecules 2004, 37, 10001. (10) Hao, J.; Weiss, R. A. Macromolecules 2011, 44, 9390−9398. (11) Hao, J.; Weiss, R. A. Polymer 2013, 54, 2174−2182. (12) Hao, J.; Weiss, R. A. Macro Lett. 2013, 2, 86−89. (13) Renault, B.; Cloutet, E.; Lacroix-Desmazes, P.; Cramail, H. Macromol. Chem. Phys. 2008, 209, 535. (14) Ezrin, M.; Lavigne, G. Proc. Annu. Technol. Conf., Soc. Plast. Eng. 1991, 49, 2230−2233. (15) Wang, Q.; Zhang, Q.; Zhan, X.; Chen, F. J. Polym. Sci., Part A: Polym. Chem. 2010, 48, 2584. (16) Ferriol, M.; Gentilhomme, A.; Cochez, M.; Oget, N.; Mieloszynski, J. L. Polym. Degrad. Stab. 2003, 79, 271. (17) Manring, L. E. Macromolecules 1988, 21, 528. (18) Manring, L. E. Macromolecules 1989, 22, 2673−7. (19) Manring, L. E.; Sogah, D. Y.; Cohen, G. M. Macromolecules 1989, 22, 4652−4. (20) Manring, L. E. Macromolecules 1991, 24, 3304−9. (21) Emerging ContaminantsPerfluorooctane Sulfonate and Perfluorooctanoic Acid, Emerging Contaminants Fact Sheet; U.S. Envronmental Protection Agency: EPA 505-F-13-002, March, 2013. (22) Lau, C.; Anitole, K.; Hodes, C.; Lai, D.; Pfahles-Hutchens, A.; Seed, J. Toxicol. Sci. 2007, 99, 366. (23) Copart, J.-M.; Girault, S.; Juhue, D. Macromolecules 2001, 17, 7237−7244. (24) Honda, K.; Morita, M.; Otsuka, H.; Takahara, A. Macromolecules 2005, 38, 5699. (25) Honda, K.; Morita, M.; Sakata, O.; Sasaki, S.; Takahara, A. Macromolecules 2010, 43, 454. (26) Roe, R. J. Methods of X-ray and Neutron Scattering in Polymer Science; Oxford University Press. New York, 2000. (27) Matsen, M. W. J. Chem. Phys. 2000, 113, 5539−5544. (28) Semenov, A. M. Macromolecules 1993, 26, 6617−6621. (29) Mark, J. E.; Erman, B. Rubberlike Elasticity: A Molecular Primer; John Wiley & Sons, Inc.: New York, 1988.

asymmetric. The bcps exhibited a complex microstructure with two length scales, one being the bcp texture of ∼20 nm and the other being a smectic phase of ∼3 nm. The phase behavior of the bcps corresponded fairly well with the predictions of Matsen27 for triblock copolymers with molecular weight asymmetry. Diblock copolymers were inefficient at forming hydrogels, but the triblock copolymers did form highly elastic hydrogels with water swelling ratios as high as S = 500, depending on the block copolymer composition. Block copolymer hydrogels based on DMA and FA were more high swollen than random copolymers with the same composition, which is believed to be a consequence of a water depleted PD layer that surrounded the PF nanodomains. For a given composition the ratio of interfacial area to volume of the PF nanodomains was an order of magnitude greater in the random copolymer hydrogel than in a block copolymer hydrogel, which led to less available hydrogen bonding sites for absorbing water in the random copolymer. Swelling also induced a block microstructure change from LAM to HEX, which was reversible when the hydrogel was dried. The dynamic shear modulus of the block copolymer hydrogels ranged from 103 to 4 × 104 Pa, depending on the composition, and exhibited a weak dependence on frequency due to the viscoelastic nature of the physical crosslinks, i.e., the hydrophobic nanodomains that anchored a network of the hydrophilic PD chains. These block copolymer hydrogels also exhibited a smectic liquid crystalline phase within the hydrophobic nanodomains as a consequence of organization of the perfluoroalkyl side chains. The smectic phase exhibited an order−disorder transition at ∼100 °C.

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ASSOCIATED CONTENT

S Supporting Information *

Figure S1. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected] (R.A.W.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by a grant from the Civil, Mechanical and Manufacturing Innovation Division (Engineering Directorate) of the National Science Foundation, Grant CMMI1300212. We thank Prof. Greg Sotzing at the University of Connecticut for the DTD-GC/MS analysis.

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REFERENCES

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