Formation and Confined Crystallization of Polyethylene Nanophases

Apr 1, 2013 - ... Mauricio Lepienski , Eduardo Lenz Cardoso , Luiz Antônio Coelho , Daniela Becker ... J. Morales , Silvia Goyanes , Norma B. D?Accor...
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Formation and Confined Crystallization of Polyethylene Nanophases in Epoxy Thermosets Chongyin Zhang, Lei Li, and Sixun Zheng* Department of Polymer Science and Engineering and the State Key Laboratory of Metal Matrix Composites, Shanghai Jiao Tong University, Shanghai 200240, P. R. China ABSTRACT: In this contribution, we reported the investigation of the formation and confined crystallization behavior of polyethylene nanophases in epoxy thermosets. The nanostructured epoxy thermosets were prepared by the use of a poly(ε-caprolactone)-block-polyethylene-block-poly(ε-caprolactone) (PCL-b-PE-b-PCL) triblock copolymer. The crystalline midblock (viz. PE) of the triblock copolymer was prepared from an α,ω-diacetoxy-terminated polycyclooctadiene with the molecular weight as high as Mn = 11,000, which was synthesized via the ring-opening metathesis polymerization (ROMP) of cyclooctadiene catalyzed by Grubbs second generation catalyst. The formation of PE nanophases in epoxy thermosets was evidenced with transmission electronic microscopy (TEM), small-angle X-ray scattering (SAXS) and dynamic mechanical thermal analysis (DMTA). It was found that in the nanostructured thermosets, the spherical nanophases of PE with the size of 20−30 nm in diameter were dispersed into the continuous epoxy matrices. Wide angle X-ray diffraction (XRD) showed that the formation of PE nanophases did not alter the structure of PE crystals. The investigations of isothermal and nonisothermal crystallization kinetics showed that the crystallization of PE in the nanostructured thermosets was in a confined manner and the confinement has been interpreted on the basis of nanoscaled space, interdomain connectivity, and the cross-linked structures of epoxy matrices.

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improved.1,2 In addition, the nanostructured thermosets could be a class of versatile precursors to other nanomaterials with new and functional properties.7,8 In the past years, a variety of block copolymer architectures have been employed to access ordered or disordered nanostructures in thermosets via self-assembly or RIMPS approaches.9−39 However, most of these nanophases formed in the thermosets are not crystallizable. Although some block copolymers containing crystalline subchains such as poly(ethylene oxide) and poly(ε-caprolactone) have been used to access the nanophases in thermosets, these crystallizable blocks are generally thermoset-philic. Upon incorporating these block copolymers into the thermosets such as epoxy,3−9,12,17−19,21,26,29−31,35,38,39 phenolic4,40 or unsaturated resins,27 they are no longer crystallizable owing to their miscibility with these thermosets. It is of interest and importance to investigate the formation and crystallization behavior of crystalline nanophases in thermosets. Polyethylene (PE) could be an ideal candidate of subchain for the block copolymers to access the crystalline nanophases in thermosets owing to its high crystallinity and thermoset-phobicity. However, such an investigation remains largely unexplored.

INTRODUCTION Incorporating amphiphilic block copolymers into thermosetting polymers is an efficient approach to access ordered or disordered nanostructures in thermosets.1,2 In 1997, Hillmyer et al.3,4 first reported the strategy of creating nanophases in epoxy thermosets via the mechanism of self-assembly by the use of amphiphilic block copolymers. In this protocol, the precursors of epoxy act as the selective solvent of the block copolymers; some self-assembled nanostructures such as spherical, cylindrical, bicontinuous, and lamellar nanophases were generated before curing reaction; these preformed nanophases were fixed with the subsequent curing reaction. More recently, it is found that by using block copolymers reaction-induced microphase separation (RIMPS) mechanism was also involved with the formation of ordered or disordered nanophases in thermosets.5,6 In the RIMPS approach, it is not required that block copolymers are self-assembled into nanophases prior to curing reaction; all the subchains of the block copolymers may be miscible with precursors of thermosets. The microphase-separated morphologies are not formed until the polymerization reaction (viz. curing) occurs with a sufficient conversion of monomers. It has been identified that the formation of nanophases in thermosets can significantly optimize the intercomponent interactions between thermosetting matrix and modifiers and thus mechanical properties of materials (e.g., fracture toughness) can be greatly © 2013 American Chemical Society

Received: January 10, 2013 Revised: March 6, 2013 Published: April 1, 2013 2740

dx.doi.org/10.1021/ma4000682 | Macromolecules 2013, 46, 2740−2753

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In this contribution, we first reported the synthesis of an α,ωdiacetoxy-terminated polycyclooctadiene (PCOD) via the ringopening metathesis polymerization (ROMP) of cyclooctadiene with a sufficiently high molecular weight (e.g., the molecular weight being as high as Mn =104). Thereafter, the α,ωdiacetoxy-terminated PCOD was used to obtain an α,ωdihydroxyl-terminated hydrogenated PCOD (viz. polyethylene, PE). The latter was then used as a macromolecular initiator for the ring-opening polymerization of ε-caprolactone to afford poly(ε-caprolactone)-block-polyethylene-block-poly(ε-caprolactone) (PCL-b-PE-b-PCL) triblock copolymer. The formation of PE nanophases in epoxy thermosets containing PCL-b-PE-bPCL triblock copolymer was investigated with transmission electronic microscopy (TEM), small-angle X-ray scattering (SAXS) and dynamic mechanical thermal analysis (DMTA). The crystallization behavior of PE in the nanostructured thermosets was investigated by means of differential scanning calorimetry (DSC).

Guo et al.41 first reported the formation of polyethylene (PE) nanophases in epoxy thermosets by using poly(ethylene oxide)block-polyethylene (PEO-b-PE) diblock copolymer. It was found that the crystallization of the PE nanophases in the thermosets was in a confined manner.41,42 It should be pointed out that the PEO-b-PE diblock copolymer used in these previous studies was nonetheless a low-molecular-weight nonionic surfactant43 and it had a quoted molecular weight of Mn = 1400 Da with the PE mass fraction of 50 wt %. To the best of our knowledge, there has been no previous report on the formation of PE nanophases in thermosets by the use of high-molecular-weight block copolymer with sufficiently long PE blocks (e.g., the length of PE subchain being as high as 104 Da). Such an investigation requires the synthesis of the welldefined block copolymers containing both thermoset-philic and PE subchains. Generally, sequential polymerization approaches can be employed to obtain PE-containing block copolymers; the PE blocks can be synthesized in direct or indirect way.44 For a direct preparation, the PE chains with chain-end functionality can be synthesized via coordination olefin polymerization with transition-metal catalysts.45,46 It is found that some group 4 transition-metal catalysts bearing two phenoxyimine ligands are capable of providing vinyl-terminated PE. Nonetheless, these coordination polymerizations require high pressure of olefin. Recently, Shea et al.47 reported the synthesis of polymethylene with chain-end functionality via a boron-catalyzed polymerization of dimethylsulfoxonium methylide. It was realized that the synthesis of polymethylene with high molecular weights requires the initiator with sufficiently high purity to achieve high stoichiometric ratios with analytical precision. For an indirect approach, PE blocks can be obtained via the hydrogenation of polybutadiene with end functionality. Generally, polybutadiene is synthesized via anionic polymerization of 1,3-butadine. Nonetheless, anionic polymerization of 1,3-butadiene would render 1,2- and 1,4-addition structures, the ratios of which are quite dependent on the polarity of the solvents used.48 Therefore, the resulting hydrogenated polybutadiene can be taken as the copolymers of ethylene with butane (or ethyl ethylene) since the hydrogenation of 1,4addition structures afford the structural units of butane (or ethyl ethylene). It is such a small number of copolymerization units that can cause the depression in melting point and crystallinity of PE chains. It is worth pointing out that Hillmyer and Grubbs et al.49 more recently reported the ring-opening metathesis polymerization (ROMP) of cyclooctadiene (COD) by the use of ruthenium-based metathesis catalyst. The ROMP of COD yields α,ω-difunctional polycyclooctadiene (PCOD), which can be exactly taken as a polybutadiene of 1,4-addition. Therefore, the hydrogenation of an α,ω-difunctional PCOD will definitely afford an α,ω-difunctional PE homopolymer, which is suitable for the synthesis of the PE-containing block copolymers via sequential copolymerization. In this work, the latter approach would be adopted to synthesize a triblock copolymer composed of polyethylene (PE) and poly(εcaprolactone) (PCL) subchains. The design of the block copolymer is based on the knowledge that PCL is miscible with epoxy thermosets after and before curing reaction.50,51 It is expected that by the use of PCL-b-PE-b-PCL triblock copolymer the nanostructured thermosets containing the crystalline PE nanophases can be accessed via a mechanism of self-assembly followed by curing reaction.

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EXPERIMENTAL SECTION

Materials. Diglycidyl ether of bisphenol A (DGEBA) with epoxide equivalent weight of 185−210 was purchased from Shanghai Resin Co., China. 4,4′-Methylenebis(2-chloroaniline) (MOCA) was used as the hardener, purchased from Shanghai Reagent Co., China. 1,5Cyclooctadiene (COD) (99%) was purchased from TCI Co., Shanghai, China, and it was distilled over calcium hydride (CaH2) prior to use. Grubbs second generation catalyst (97%) was purchased from Aldrich Co., USA, and used as received. cis-2-Butene-1,4-diol diacetate (>95%) was purchased from TCI Co., Shanghai, China and used as the chain transfer agent of ring-opening metathesis polymerization (ROMP). Before use, cis-2-butene-1,4-diol diacetate was distilled under reduced pressure. p-Toluenesulfonyl hydrazide was supplied by Bangcheng Chemical Co., Shanghai, China, and used as received. ε-Caprolactone (99%) (CL) were purchased from Acros Co., Shanghai, China. Before use, CL was distilled over CaH2 under reduced pressure. Stannous octanoate [Sn(Oct)2] was purchased from Aldrich Co. Shanghai, China. All other reagents and solvents used in this work were purchased from Shanghai Reagent Co., China. Before use, toluene, o-xylene and tetrahydrofuran were refluxed over sodium and then distilled. Synthesis of α,ω-Dihydroxyl-Terminated Polycyclooctadiene. To a flame-dried flask was charged 15 mL of anhydrous chloroform. The solvent was purged with argon for 30 min and then cis-2-butene-1,4-diol diacetate (0.054g, 0.344 mmol) was added. The Grubbs second generation catalyst (2.2 mg, 2.6 mmol) dissolved in 1.0 mL chloroform was added with a gastight syringe. The flask was immersed in a thermostated bath at 40 °C and then 5.0 mL COD (4.3 g, 39.8 mmol) was dropwise added within 1 h. The polymerization was performed at 40 °C with vigorous stirring for 10 h. The polymerization was terminated by dropping the reacted mixture into 100 mL methanol containing 15 mL 1.0 M hydrochloric acid aqueous solution and the precipitates were isolated with filtration. After dried in vacuo at 30 °C for 24 h, the product (4.14 g) was obtained with the yield of 95%. 1H NMR (ppm, CDCl3): 5.30−5.40 (m, 370H, −CH CHCH2−), 4.62, 4.51 (d, 2H, −CHCHCH2OAc), 2.00−2.40 (m, 740H, −CH2CHCHCH2−). The above α,ω-diacetyl-terminated polycyclooctadiene (4.10g) was dissolved in 50 mL tetrahydrofuran at 0 °C and 15 mL 0.7 M methanol solution of sodium methoxide (NaOMe) (0.0095 g, 1.75 mmol with respect of NaOMe) was added within 1 h with vigorous stirring. The hydrolysis reaction was performed at 0 °C for additional 5 h and then the solution was dropped into 200 mL cold methanol containing 1.5 mL 1.0 M HCl aqueous solution to afford the precipitates. After dried in vacuo at 30 °C for 24 h, the α,ω-dihydroxylterminated polycyclooctadiene (3.51 g) was obtained with the yield of 85%. 1H NMR (ppm, CDCl3): 5.30−5.50 (m, 370H, −CH CHCH2−), 4.09, 4.18 (t, 2H, −CHCHCH2OH) 1.95−2.15 (m, 2741

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742H, −CH2CHCHCH2−). GPC (PS standard, THF): Mn = 11 000 with Mw/Mn = 1.53 Hydrogenation of α,ω-Dihydroxyl-Terminated Polycyclooctadiene. The hydrogenation of the above α,ω-dihydroxyl-terminated polycyclooctadiene was carried out by following the literature method.52−55 Typically, the α,ω-dihydroxyl-terminated polycyclooctadiene (3.20g, 59.3 mmol with respect of carbon−carbon double bonds) was dissolved in anhydrous xylene (300 mL) and then ptoluenesulfonhydrazide (66.2204 g, 0.36 mol) and 3,5-di-tert-butyl-4hydroxytoluene (99.3 mg) (viz. antioxidant) were added. The mixture was refluxed at 120 °C in the atmosphere of highly pure nitrogen for 15 h. The degree of hydrogenation was traced by means of Fourier transform infrared spectroscopy (FTIR). Thereafter, the solution was cooled to room temperature and the solution was concentrated via rotary evaporation. The concentrated solution was dropped into a great amount of ice-cold methanol to afford the precipitates. After dried at 40 °C in vacuo for 36 h, the product (3.420g) was obtained with the yield of 98%. It was noted that the product was no longer soluble in the solvents of polycyclooctadiene such as tetrahydrofuran and chloroform but it can be dissolved in some nonpolar solvents such as benzene, toluene and xylene at elevated temperature. 1H NMR (ppm, benzene-d6, 80 °C): 3.60 (t, 2H, −CH2OH) and 1.40 (s, 1502H, −CH2−). Synthesis of PCL-b-PE-b-PCL Triblock Copolymer. Poly(εcaprolactone)-block-polyethylene-block-poly(ε-caprolactone) triblock copolymer (PCL-b-PE-b-PCL) was synthesized via the ring-opening polymerization of ε-caprolactone (CL) with the α,ω-dihydroxylterminated hydrogenated polycyclooctadiene (i.e., polyethylene) as the macromolecular initiator and with stannous(II) octanoate [Sn(Oct)2] as the catalyst. Typically, the macromolecular initiator (0.40g, 0.04 mmol), CL (1.00g, 0.932 mol) and anhydrous xylene (8 mL) were added to a predried flask equipped with a magnetic stirrer. The flask was connected to a Schlenk line to degas via three pump-freeze−thaw cycles and then Sn(Oct)2 dissolved in anhydrous toluene [1/1000 (wt) with respect to CL] was added using a syringe. The flask was immersed in a thermostated oil bath at 120 °C for 36 h to attain a complete polymerization. Cooled to room temperature, the crude product was dissolved in tetrahydrofuran (20 mL) and the solution was dropped into 200 mL cold methanol to afford the precipitates. This procedure was repeated three times to purify the product. After dried in vacuo at 30 °C for 36 h, the polymer (1.21 g) was obtained with the yield of 86.4%. 1H NMR (ppm, toluene-d8, 100 °C): 4.0−4.10 (m, 180H, −OCCH 2 CH 2 CH 2 CH 2 CH 2 O−), 3.44 (t, 2H −OCCH2CH2CH2CH2CH2OH), 2.15−2.25 (m, 180H, −OCCH2CH2CH2CH2CH2O−), 1.50−1.70 (m, 358H, −OCCH2CH2CH2CH2CH2O−), 1.35−1.50 (s, 750H, −CH2CH2−) and 1.25−1.35 (m, 180H, −OCCH2CH2CH2CH2CH2O−). Preparation of Epoxy Thermosets Containing PCL-b-PE-bPCL. Desired amount of PCL-b-PE-b-PCL triblock copolymer was added to DGEBA with continuous stirring at 100 °C until the mixtures became homogeneous and transparent and then 4,4′-methylenebis(2chloroaniline) was added with continuous stirring until the full dissolution of the curing agent. The mixtures were poured into Teflon molds and cured at 150 °C for 3 h plus 180 °C for 2 h. The thermosets containing the block copolymer up to 50 wt % were obtained. Measurement and Techniques. Fourier Transform Infrared Spectroscopy (FTIR). The FTIR measurements were conducted on a Perkin-Elmer Paragon 1000 Fourier transform spectrometer at room temperature (25 °C). The films of specimens were obtained via casting the chloroform (or toluene) solution of the samples (2 wt %) onto KBr windows. For the preparation of thermoset specimens, the thermosets were granulated and the powder was mixed with KBr pellets to press into the small flakes for measurements. All the specimens were sufficiently thin to be within a range where the Beer− Lambert law is obeyed. In all cases 64 scans at a resolution of 2 cm−1 were used to record the spectra. Nuclear Magnetic Resonance Spectroscopy (NMR). The 1H NMR measurement was carried out on a Varian Mercury Plus 400 MHz NMR spectrometer at 25 °C and elevated temperatures (e.g., 80 and 100 °C). The samples were dissolved with deuterated chloroform

(CDCl3) (or/and benzene-d6, toluene-d8) the solutions were measured with tetramethylsilane (TMS) as an internal reference. Gel Permeation Chromatography (GPC). The molecular weights were determined on a Waters 717 Plus autosampler gel permeation chromatography apparatus equipped with Waters RH columns and a Dawn Eos (Wyatt Technology) multiangle laser light scattering detector and the measurements were carried out at 25 °C with tetrahydrofuran (THF) as the eluent at the rate of 1.0 mL/min. Wide Angle X-ray Diffraction. The wide-angle X-ray diffraction (XRD) experiments were carried out on a Shimadzu XRD-6000 X-ray diffractometer with Cu Kα (λ = 0.154 nm) irradiation at 40 kV and 30 mA using a Ni filter. Data were recorded in the range of 2θ = 5−40° at the scanning rate and step size of 4.0 o/min and 0.02 o, respectively. Small-Angle X-ray Scattering (SAXS). The SAXS measurements were taken on a small-angle X-ray scattering station (BL16B1) with a long-slit collimation system in the Shanghai Synchrotron Radiation Facility (SSRF), Shanghai, China, in which the third generation of synchrotron radiation light source was employed. Two dimensional diffraction patterns were recorded using an image intensified CCD detector. The experiments were carried out with the radiation of X-ray with the wavelength of λ = 1.24 Å at room temperature (25 °C). For variable temperature experiments, a TH MS600 Linkam hotstage with the precision of 0.1 °C was mounted. The intensity profiles were output as the plot of scattering intensity (I) versus scattering vector, q = (4π/λ) sin(θ/2) (θ = scattering angle). Transmission Electron Microscopy (TEM). Transmission electron microscopy (TEM) was performed on a JEOL JEM-2010 highresolution transmission electron microscope at an acceleration voltage of 120 kV. The samples were trimmed using a microtome machine and the sectioned samples were stained with RuO4 to increase the contrast. The stained specimens (c.a., 70 nm in thickness) were placed in 200 mesh copper grids for observations. Dynamic Mechanical Thermal Analysis (DMTA). Dynamic mechanical tests were carried out on a TA Instruments DMA Q800 dynamic mechanical thermal analyzer (DMTA) equipped with a liquid nitrogen apparatus in a single cantilever mode. The frequency used was 1.0 Hz, and the heating rate of 3.0 °C/min was used. The specimen dimension was 25 × 5.0 × 1.75 mm3. The experiments were carried out from −100 to 200 °C. Differential Scanning Calorimetry (DSC). Thermal analysis was performed on a Perkin-Elmer Pyris-1 differential scanning calorimeter in a dry nitrogen atmosphere. The instrument was calibrated with a standard Indium. The samples (about 10.0 mg in weight) were first heated up to 150 °C and held at this temperature for 3 min to eliminate thermal history, followed by quenching to −60 °C. In all the cases, the heating rate of 20 °C/min was used to record the heating thermograms and the cooling rate of −10 °C/min to record the cooling thermograms. Glass transition temperature (Tg) was taken as the midpoint of heat capacity change, melting temperature (Tm) and crystallization temperature (Tc) were taken as the temperatures at the maxima of endothermic transitions and the minima of exothermic transitions, respectively. To investigate the isothermal crystallization kinetics of plain PE, PE blocks in PCL-b-PE-b-PCL triblock copolymer and in the nanostructured thermosets, all the samples were hold at 150 °C for 3 min to melt the crystals and then quenched to the desired temperatures for isothermal crystallization toward completion. The conversions of crystallization, X(t), as a function of crystallization time (t) were determined using the following equation: X(t ) =

∫0

t

⎛ dH ⎞ ⎜ ⎟ dt / ⎝ dt ⎠

∫0

∞ ⎛ dH ⎞ ⎜

⎟

⎝ dt ⎠

dt

(1)

where the integral in the numerator is the enthalpy generated at time t and the integral in the denominator is the total enthalpy of crystallization at t = ∞.

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RESULTS AND DISCUSSION Synthesis of PCL-b-PE-b-PCL Triblock Copolymer. The route of synthesis for poly(ε-caprolactone)-block- polyethylene2742

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Scheme 1. Synthesis of PCL-b-PE-b-PCL Triblock Copolymer

block-poly(ε-caprolactone) (PCL-b-PE-b-PCL) triblock copolymer was depicted in Scheme 1. The ring-opening polymerization of ε-caprolactone (CL) with α,ω-dihydroxyl-terminated polyethylene as the macromolecular initiator was carried out to obtain the triblock copolymer. In order to synthesize the α,ωdihydroxyl-terminated polyethylene, the ring-opening metathesis polymerization (ROMP) of cyclooctadiene (COD) with cis-2-butene-1,4-diol diacetate as the chain transfer agent was performed, which was catalyzed by Grubbs second generation catalyst56−62 and the ROMP of COD afforded an α,ωdiacetoxy-terminated polycyclooctadiene (PCOD). With the hydrolysis reaction, the terminal groups of the α,ω-diacetoxyterminated PCOD was readily converted into primary hydroxyl groups, i.e., an α,ω-hydroxy-terminated PCOD was obtained. With p-toluenesulfonyl hydrazide as a reducing agent, the latter was further reduced into an α,ω-hydroxy-terminated hydrogenated PCOD, i.e., α,ω-hydroxy-terminated polyethylene (PE). The reaction of reduction was carried out in the solution of xylene at 120 °C for 15 h and the degree of hydrogenation was traced by means of Fourier transform infrared spectroscopy (FTIR). Shown in Figure 1 are the 1H NMR spectra of α,ωdiacetoxy-, α,ω-dihydroxyl-terminated PCOD and α,ω-dihydroxyl-terminated polyethylene. For α,ω-diacetoxy-terminated PCOD, the signals of resonance at 1.90−2.30 ppm are assignable to the protons of main-chain methylene and methyl in terminal acetoxyl groups; the resonance of methine protons of the main chain was detected at 5.41 ppm. The signal of resonance assignable to the protons of the methylene adjacent to the terminal acetoxyl groups appeared at 4.51 ppm. With the hydrolysis of the terminal acetoxyl groups of α,ω-diacetoxyterminated PCOD, the signal of resonance at 4.51 ppm was observed completely to shift to 4.10 ppm whereas other signals of resonance remained unchanged, indicating that the terminal acetoxyl groups were fully hydrolyzed into the terminal hydroxyl groups. The α,ω-dihydroxyl-terminated PCOD was subjected to gel permeation chromatography (GPC) and the GPC curve was presented in Figure 2. The GPC measurement gave the molecular weight of Mn = 11 000 with Mw/Mn = 1.53.

Figure 1. 1H NMR spectra of α,ω-diacetoxy-, α,ω-dihydroxylterminated PCOD, α,ω-dihydroxyl-terminated hydrogenated PCOD (viz. PE) and PCL-b-PE-b-PCL triblock copolymer.

The polydispersity of the α,ω-dihydroxyl-terminated PCOD was comparable to the results of literature.59,61,62 The numberaverage molecular weight by GPC was in good agreement with that according to end group analysis by means of 1H NMR spectroscopy. The α,ω-hydroxy-terminated PCOD was hydrogenated to afford an α,ω-hydroxy-terminated PE. After it was hydrogenated, only two signals of resonance were displayed in 2743

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Figure 2. GPC curve of α,ω-dihydroxyl-terminated PCOD with Mn = 11 000 Mw/Mn = 1.53.

the 1H NMR spectrum of α,ω-hydroxy-terminated PE, which were detected at 1.41 and 3.62 ppm, respectively. The former is assignable to the resonance of the methylene protons in the main chains of polyethylene whereas the latter to that of the methylene connected to the terminal hydroxyl groups. According to the ratio of the integral intensity of the resonance at 3.62 ppm to that at 1.41 ppm, the molecular weight of the α,ω-dihydroxyl-terminated polyethylene was estimated to be Mn = 10,500, which is comparable to the molecular weight of the α,ω-dihydroxyl-terminated PCOD (i.e., Mn = 11,000 with Mw/Mn = 1.53). The 1H NMR spectroscopy indicate that the α,ω-diacetoxy-, α,ω-dihydroxyl-terminated PCOD and α,ωdihydroxyl-terminated hydrogenated PCOD (viz. polyethylene) were successfully obtained. The α,ω-dihydroxyl-terminated polyethylene was used as the macromolecular initiator to synthesize poly(ε-caprolactone)block-polyethylene-block-poly(ε-caprolactone) (PCL-b-PE-bPCL) triblock copolymer via the ring-opening polymerization of ε-caprolactone. The resulting product was subjected to 1H NMR spectroscopy and the 1H NMR spectrum is also incorporated in Figure 1. Compared to the 1H NMR spectrum of α,ω-dihydroxyl-terminated polyethylene, there appeared several new signals of resonance at 1.38, 1.62, 2.20, and 4.06 ppm except for the intense signals of resonance at 1.42 ppm assignable to the resonance of methylene protons in the backbone of polyethylene chain. These peaks are attributable to the resonance of methylene protons of PCL chains as indicated in Figure 1. It should be pointed out that the minor peak at 3.45 ppm is assignable to the resonance of methylene protons of terminal hydroxymethyl groups of the PCL block. According to the ratio of integral intensity for the protons connected to the terminal hydroxyl groups at 3.45 ppm to that of other methylene in the PCL chains, the length of PCL block in the triblock copolymer was estimated to Mn = 10 300 Da; i.e., the molecular weight of PCL-b-PE-b-PCL triblock copolymer is calculated to be c.a. Mn = 31 100 Da. According to the lengths of PE midblock and PCL end blocks, the mass fractions of PCL and PE blocks in the triblock copolymer were calculated to be c.a. 2:1. The results of NMR and GPC show that the PCL-bPE-b-PCL triblock copolymer was successfully synthesized. The PCL-b-PE-b-PCL triblock copolymer was subjected to differential scanning calorimetry (DSC) and the DSC curves are presented in Figure 3. In the heating scan, dual endothermic peaks were displayed at 58 and 126 °C, respectively. The former is assignable to the melting transition of PCL block

Figure 3. DSC curves of PCL-b-PE-b-PCL triblock copolymer. Up: the heating scan at the rate of 20 °C/min after quenching from the melt at 150 °C. Down: the cooling scan at the rate of 10 °C/min.

whereas the latter to that of PE block. In the cooling scan, two endothermic peaks were detected at 47 and 112 °C, which correspond to the crystallization transitions of PCL and PE blocks, respectively. The results of DSC indicate that the triblock copolymer was composed of two crystallizable subchains, i.e., PCL and PE blocks. The PCL-b-PE-b-PCL triblock copolymer was subjected to small-angle X-ray scattering (SAXS) and the SAXS profiles are shown in Figure 4. At room temperature, the triblock copolymer barely displayed discernible peaks, which could be due to the superposition of the scattering peaks of crystals and microdomains of PCL and PE blocks. To observe the scattering of PCL and PE microdomains, the block copolymer was heated up to 150 °C to melt all the crystals. It is seen that the multiple scattering peaks were exhibited; the scattering peaks were situated at q values of 1, 30.5 and 70.5 relative to the first-order scattering peak positions (qm). The SAXS results indicate that in the melt the PCL-b-PE-b-PCL triblock copolymer was microphase-separated and that the cylindrical PE nanophases could be arranged into a hexagonal lattice. Formation of PE Nanophases in Epoxy Thermosets. Self-Assembly Behavior of PCL-b-PE-b-PCL in Epoxy Precursors. By the use of block copolymers, the formation of nanophases in thermosets can follow either self-assembly3,4 or reaction-induced microphase separation mechanisms (RIMPS),5,6 depending on the miscibility of the copolymer blocks with the thermosets. For the protocol of self-assembly, precursors of thermosets act as selective solvents of block copolymers and some self-organized nanophases (i.e., micelle) are formed prior to curing. With subsequent curing reaction, these disordered and/or ordered nanostructures are locked in. For the formation of nanostructures via reaction-induced microphase separation mechanism, it is not required that all the subchains of the block copolymer are miscible with precursors of thermosets before curing. Part of the copolymer 2744

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Figure 4. SAXS profiles of PCL-b-PE-b-PCL triblock copolymer at 25 and 150 °C.

Figure 5. SAXS profiles of the mixtures composed of DGEBA, MOCA, and PCL-b-PE-b-PCL triblock copolymer at 150 °C. The red, heavier, and smooth lines are the simulated form factor scattering with the Percus−Yevick model for disordered packing of spheres.

blocks are demixed out of the matrix of thermosets with the occurrence of the curing reaction. In the present work, we designed and synthesized the PCL-bPE-b-PCL triblock copolymer by knowing that: (i) PCL is miscible with epoxy before curing and also miscible with 4,4′methylenebis(2-chloroaniline)(MOCA)-cured epoxy50,51 and (ii) PE is immiscible with epoxy after and before the curing reaction. Therefore, the PCL-b-PE-b-PCL triblock copolymer could be self-assembled into nanophases in epoxy before curing reaction. The self-assembly behavior of PCL-b-PE-b-PCL in the precursors of epoxy (viz. DGEBA + MOCA) was investigated by means of small-angle X-ray scattering (SAXS). To melt the crystals of PE, the ternary mixtures composed of DGEBA, MOCA and PCL-b-PE-b-PCL triblock copolymer were rapidly heated up to 150 °C, and the SAXS measurements were carried out at this temperature. Shown in Figure 5 are the SAXS profiles. In all the cases, the well-defined scattering peaks were displayed, indicating that there existed the microdomains in the ternary mixtures composed of DGEBA, MOCA and PCL-b-PEb-PCL triblock copolymer at 150 °C, i.e., the PCL-b-PE-b-PCL triblock copolymer was indeed self-organized into the nanophases in the precursors of epoxy at 150 °C. It is proposed that the self-organized nanophases would be fixed if the ternary mixtures were cured at 150 °C. Nanostructures of Thermosets. The ternary mixtures composed of DGEBA, MOCA, and PCL-b-PE-b-PCL triblock copolymer were cured at 150 °C to obtain the thermosets. All the as-prepared thermosets were homogeneous and transparent, indicating that no macroscopic phase separation occurred in the process of curing reactions. The morphologies of the epoxy thermosets containing PCL-b-PE-b-PCL triblock copolymer were investigated by means of transmission electron microscopy (TEM) and small-angle X-ray scattering (SAXS). Shown in Figure 6 are the TEM micrographs of the thermosets containing PCL-b-PE-b-PCL triblock copolymer up to 50 wt %.

Prior to the morphological observation, the specimens of ultrathin sections were stained with RuO4 to increase the contrast. In this case, epoxy matrix was oxidized by RuO4 whereas the PE microdomains remained less affected. It is seen that in all the cases, the epoxy thermosets containing PCL-bPE-b-PCL triblock copolymer were microphase-separated. For the thermosets with the PCL-b-PE-b-PCL triblock copolymer up to 40 wt %, the spherical microdomains with the diameter of 10−20 nm were dispersed into the continuous matrices (Figure 6A through 6D). With increasing the content of PCL-b-PE-bPCL triblock copolymer, the sizes and quantity of the spherical microdomains slightly increased. The spherical nanophases are attributable to PE blocks whereas the continuous matrix to epoxy which was mixed with PCL blocks. For the thermoset containing 50 wt % of PCL-b-PE-b-PCL triblock copolymer, some PE nanodomains became worm-like; i.e., the short cylinders with the diameter of c.a. 10 nm began to appear at this composition (See Figure 6E). The formation of PE nanophases was further investigated by means of small-angle X-ray scattering (SAXS) and the SAXS profiles were presented in Figure 7. To melt the crystals of PE microdomains, all the nanostructured thermosets were heated up to 150 °C, at which the SAXS measurements were carried out. It is seen that the thermosets clearly displayed the scattering peaks, the intensity of which increased with increasing the content of PCL-b-PE-b-PCL triblock copolymer. The results of SAXS indicate that the thermosets were indeed microphase-separated. It is noted that the SAXS profiles presented in Figure 7 were quite close to those of the ternary mixtures composed of DGEBA, MOCA, and PCL-b-PE-b-PCL triblock copolymer at 150 °C (see Figure 5). This observation suggests that the microphase-separated morphologies in the thermosets were formed from the fixation of the preformed 2745

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Figure 6. TEM micrographs of the epoxy thermosets containing PCL-b-PE-b-PCL triblock copolymer: (A) 10, (B) 20, (C) 30, (D) 40, and (E) 50 wt %.

measured SAXS data by the use of a model of hard sphere with polydispersity. The structure factor scattering for hard spheres was fitted according to Percus−Yevick closure relation.63 The only interaction effect taken into account in this model is the excluded volume present in a dispersion of hard spheres. The polydispersity is taken into account by simply averaging the partial structure factor of the single components.64 The

microphase-separated morphologies before the curing reaction. In other words, the formation of PE nanophases in the thermosets followed the mechanism of self-assembly. In Figures 5 and 7, the broad and round scattering peaks were exhibited. It is proposed that these scattering peaks resulted from the form factor scattering of the spherical or worm-like nanoobjects in the thermosets. Referring to the TEM results, we fitted the 2746

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Table 1. Characterization of Microdomain Structures for the Mixtures Composed of DGEBA, MOCA, and PCL-b-PE-bPCL Triblock Copolymers after and before the Curing Reaction fitting parameters microdomains of PE

10

sphere

20

sphere

30

sphere

40

sphere

50

sphere + short cylinder

before curing

after curing

R = 11.4 nm, σ = 0.12 R = 17.5 nm, σ = 0.12 R = 27.4 nm, σ = 0.12 R = 26.8 nm, σ = 0.12 R = 20.5 nm, σ = 0.32

R = 19.0 nm, σ = 0.33 R = 23.1 nm, σ = 0.14 R = 32.9 nm, σ = 0.20 R = 33.9 nm, σ = 0.26 R = 30.1 nm, σ = 0.31

utilized since the results of TEM showed that this thermoset displayed a combined morphology including spherical and cylindrical nanophases (Figure 6E). For comparison, we also fitted the measured SAXS data of the mixtures composed of DGEBA, MOCA and PCL-b-PE-b-PCL triblock copolymer assuming that the mixtures possessed the identical morphologies with those of the nanostructured thermosets; the fitting results are also summarized in Table 1. It is worth noticing that before the curing reaction, the sizes of PE nanophases in the ternary mixtures composed of DGEBA, MOCA, and PCL-bPE-b-PCL triblock copolymer were significantly smaller than those of PE nanophases in the cured blends. In other words, the sizes of PE nanophases were increased with the occurrence of curing reaction. It is proposed that the increased sizes of PE nanophases measured by means of SAXS resulted from the demixing of the miscible blocks (viz. PCL) out of epoxy matrix with the occurrence of curing reaction. In the formation of nanophases in thermosets via self-assembly mechanism, the demixing behavior has been interpreted as a transition from equilibrium morphology to a chemically pinned metastable state as the cross-linking reaction progresses through the gel point.3,4 The demixing behavior of miscible blocks was recently found in the nanostructured epoxy thermosets via reactioninduced microphase separation mechanism.32,39 The demixed PCL chains would surround the micelles of PE. As a consequence, the sizes of PE nanophases measured by means of SAXS were bigger than those of the PE cores. In fact, the demixing behavior of PCL out of epoxy matrix was further demonstrated by plotting experimental and calculated Tg’s according to Gordon−Taylor equation as functions of content of PCL by the use of the model PCL homopolymer with the molecular weight identical with the length of PCL block in the triblock (viz. Mn = 11 000) (the results in detail not presented here for brevity).30,32,35 In addition to the demixing of PCL blocks out of epoxy matrix, the possibility that the increased sizes of the PE micelles resulted from the coalescence of the initial micelles owing to curing cannot be excluded. The formation of the PE nanophases in the epoxy thermosets was further evidenced by dynamic mechanical thermal analysis (DMTA) and the DMTA curves of the nanostructured thermosets in the range of −100 to 200 °C are shown in Figure 8. The control epoxy exhibited a well-defined α transition centered at ca. 159 °C, which corresponded to the glass−rubber transition of the cross-linked polymer. Apart from the α transition related to glass−rubber transition, the

Figure 7. SAXS profiles of the nanostructured thermosets PCL-b-PEb-PCL triblock copolymer at 150 °C. The red, heavier, and smooth lines are the simulated form factor scattering with the Percus−Yevick model for disordered packing of spheres.

intensity of scattering at a given scattering vector [I(q)], depends on the square of the contrast difference (Δρ2), the number of scattering particles (N), the shape and size of the scattering particle described by the form factor [P(q)], and interdomain correlations accounted for by a structure factor [S(q)]: I(q) = Δρ2 × N × P(q) × S(q)

PCL-b-PE-bPCL (wt %)

(2)

Assuming the PCL blocks (viz. corona) have essentially the same scattering density as the epoxy matrix (i.e., they are contrast-matched), the scattering arises almost from the contrast between PE microdomains and the matrix composed of epoxy and PCL. The scattering of PE microdomains is represented by the spherical form factor convoluted with a Gaussian distribution with a standard deviation (σ) to account for polydispersity in the radii of the spherical microdomains. The interparticle interaction was modeled as hard-sphere potentials between micelles with disordered (liquid-like) packing in a matrix as described by Ornstein and Zernike65 and Percus and Yevick.63 Also incorporated in Figure 7 are the form factor scattering [P(q)] (i.e., the red, heavier and smooth lines). The curving fitting yielded both the radii of micelles and standard error, which are summarized in Table 1. The fitting results showed that the radii of the spherical PE nanodomains were in the range of 19 to 34 nm with standard error of 0.14− 0.33. The small values of polydispersity imply that the sizes of the spherical PE nanophases were quite uniform. It is seen that the polydispersity slightly increased with increasing the content of the triblock copolymer in the thermosets. It should be pointed out that the size for the PE microdomains for the thermoset containing 50 wt % of PCL-b-PE-b-PCL triblock copolymer could be subject to some underestimation if the above model of hard sphere with polydispersity was only 2747

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Figure 8. DMTA curves of the epoxy thermosets containing PCL-bPE-b-PCL triblock copolymer.

Figure 9. XRD profiles of plain PE, PCL-b-PE-b-PCL and the nanostructured thermosets containing PCL-b-PE-b-PCL triblock copolymer.

thermoset exhibited the secondary transitions (viz. β relaxation) at ∼−56 °C.66−68 For the nanostructured thermosets, the α transitions assignable to epoxy matrix were observed to shift to the lower temperatures with increasing the content of PCL-bPE-b-PCL triblock copolymer. The decreased Tg ’s are attributed to the plasticization of PCL that possessed the Tg as low as −65 °C on epoxy matrices; i.e., the PCL blocks remained miscible with epoxy matrix. It should be pointed out that the glass transition of PE nanophases cannot be detected due possibly to their high crystallinity. For the nanostructured thermosets, the intensity of the β relaxations was reduced with inclusion of the triblock copolymer. The motion of hydroxyether structural units in the amine-cross-linked epoxy could be to some extent suppressed owing to its mixing with PCL chains, indicating that the PCL blocks remained miscible with epoxy matrix. It should be pointed out that in addition to the α transitions assignable to the epoxy matrices, the transitions related to the melting of PE microdomains were discernible at c.a. 125 °C for the thermosets containing 10 and 50 wt % of PCL-b-PE-b-PCL triblock copolymer. For other nanostructured thermosets, the superposition of the melting transitions of PE nanophases with the glass-to-rubber transition of epoxy matrix gave rise to the broadening of relaxation in the spectra of DMTA. In addition, it is of interest to note that at the glassy states all the nanostructured thermosets containing PCL-b-PE-b-PCL triblock copolymer possessed the storage moduli much higher than control epoxy due possible to the incorporation of the crystalline PE nanophases. Crystallization Behavior of PE Nanophases. Structure of PE Nanocrystals. Wide angle X-ray diffraction (XRD) was utilized to examine the structure of crystals of the PE in the nanophases and the diffraction patterns of PE, PCL-b-PE-bPCL, and the PE nanophases in the thermosets are shown in Figure 9. For the plain PE, two intense diffraction peaks were exhibited at 2θ = 21.42 and 23.80°, which were in accordance with those reported in the literature with unit cell parameters

(orthorhombic, a = 7.40 Å, b = 4.93 Å, c = 2.53 Å) and the diffraction peaks are assignable to (110) and (200) reflections of PE.69 For the PCL-b-PE-b-PCL triblock copolymer, the diffraction peaks at 2θ = 21.21, 21.81, and 23.52° were also detected since the (110), (111), and (200) reflections of PCL crystals were quite close to those of plain PE.70 Upon incorporating the triblock copolymer into epoxy thermosets, the diffraction peaks at 2θ = 21.42 and 23.80° were preserved and the intensity increased with increasing the content of PCLb-PE-b-PCL triblock copolymer. This result indicates that the crystalline lattice of PE in the nanophases remained almost unchanged. Compared to plain PE, however, the diffraction peaks of the PE nanophases were slightly shifted to the higher values of 2θ with decreasing the content of PCL-b-PE-b-PCL triblock copolymer in the thermosets, suggesting that the packing of PE chains in the unit cell were to some extent altered. Crystallinity of Nanophases. The nanostructured thermosets were subjected to differential scanning calorimetry (DSC) to investigate the crystallinity of PE nanophases in the thermosets. Before the measurements, all the samples were heated up to 150 °C to melt the crystals and then quenched to −60 °C for the second DSC scans. The DSC thermograms of plain PE and the nanostructured thermosets are shown in Figure 10. The plain PE displayed the melting transition at 132 °C, which is quite close to that of commercial high density polyethylene. For the PCL-b-PE-b-PCL triblock copolymer, the melting temperature of the PE was decreased to 125 °C. In addition, there was another exothermic peak at 58 °C assignable to the melting transition of PCL blocks. In contrast to PCL-b-PE-b-PCL triblock copolymer, no melting transition of PCL blocks was detected for the nanostructured thermosets (see Figure 3), indicating that the PCL chain in the thermosets were no longer crystallizable owing to its miscibility with the 2748

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Figure 11. Plot of PE crystallinity as a function of the content of PCLb-PE-b-PCL triblock copolymer for plain PE and PE in the triblock copolymer and in the nanostructured thermosets. Figure 10. DSC curves of the nanostructured thermosets containing PCL-b-PE-b-PCL triblock copolymer at the heating rate of 20 °C/min.

terms of the cooling DSC scans. Shown in Figure 12 are the cooling DSC curves of plain PE, the triblock copolymer and the

epoxy matrix.50,51 The melting temperatures (Tm’s) of PE in the nanostructured thermosets remained almost invariant but were slightly lower than that of PE in the PCL-b-PE-b-PCL triblock copolymer. With the heating DSC thermograms, the crystallinity of PE in the nanostructured thermosets can be calculated in terms of the following equation: Xc =

ΔHf /ϕPE ΔH 0f

× 100% (3)

where ΔHf is the fusion enthalpy of the rescanned samples after quenching from the melts; ΔHfo is the fusion enthalpy of the perfectly crystallized PE and equals 293 J/g;71 ϕPCL is the mass fraction of PE block in the thermosets. The plot of crystallinity (Xc) as a function of the content of PCL-b-PE-b-PCL is shown in Figure 11. It is seen that the crystallinity of PE in the thermosets was much lower than those of the plain PE and in the PCL-b-PE-b-PCL triblock copolymer. The crystallinity increased with increasing the content of PCL-b-PE-b-PCL triblock copolymer. This observation suggests that the PE crystals developed in the nanophases were less perfect than those in plain PE and in the triblock copolymer. It is proposed that the following factors could affect the crystallinity of PE in the nanostructured thermosets. First, there exists the restriction of glassy epoxy matrices on PE crystallization in the nanophases. Owing to the chemical linkage between PCL and PE block, two ends of one PE block in the nanophases were anchored onto the interface between the PE microdomain and epoxy matrix and this portion of PE chain at the intimate interface could fail to crystallize. Second, the density change that accompanies the crystallization process was hindered and the crystallization of PE in the nanostructured thermosets could be in a confined manner (vide inf ra). Crystallization Behavior of PE Nanophases. Nonisothermal Crystallization. The nonisothermal crystallization of PE in the nanostructured thermosets was investigated in

Figure 12. DSC curves of the nanostructured thermosets containing PCL-b-PE-b-PCL triblock copolymer at the cooling rate of 10 °C/min.

nanostructured thermosets. It is seen that in all these samples, the PE were capable of crystallization during the cooling scans. For plain PE and PE in the triblock copolymer, the crystallizations occurred at Tc = 118 and 112 °C, respectively. In the nanostructured thermosets, the crystallization of PE occurred at ca. Tc = 75 °C, which was much lower than those of plain PE and the PE in the triblock copolymer. If we defined 2749

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the undercooling degree of ΔT = Tm − Tc to represent the driving force required for the crystallization, the undercooling degree of both plain PE and the PE in the triblock copolymer were obtained to be ca. ΔT =14 °C. In contrast, the undercooling degrees of PE in the nanostructured thermosets were measured to be in the range of ΔT = 49−52 °C. The increased undercooling degrees indicate that the crystallization of PE in the nanostructured thermosets became much more difficult than that of plain PE or/and the PE in the triblock copolymer. This result is attributable to the formation of PE nanophases in epoxy thermosets. The microdomains (or nanophases) constituted the confined spaces, which restricted the crystallization. It is known that the confined crystallization of polymers in the nanometer scale could be affected by several factors such as size of crystalline microdomains, interdomain connectivity through grain boundary, edge and screw dislocations and rigidity of continuous matrix around the crystalline microdomains. If the crystalline microdomains are isolated, homogeneous nucleation will dominate the crystallization process and thus no crystal growth process was needed to complete the crystallization.72,73 In the present case, the spherical PE nanophases with the diameter of 10−20 nm were dispersed in continuous epoxy matrices with the content of PCL-b-PE-b-PCL triblock copolymer up to 50 wt % and the crystallization of PE was confined into the nanometer spaces, i.e., the confined crystallization occurred. Isothermal Crystallization. The isothermal crystallizations at different temperatures (i.e., Tc’s) of plain PE, PE in the triblock copolymer and in the nanostructured thermosets were investigated by means of DSC. Representatively shown in Figure 13 are the plots for the investigation of isothermal crystallization kinetics of PE in the nanostructured thermosets containing 50 wt % of PCL-b-PE-b-PCL triblock copolymer. The DSC curves of isothermal crystallization at various temperatures of crystallization (Tc ’s) showed that the completion of isothermal crystallization at Tc = 87 °C required about 50 min (see Figure 13A). However, a slight increase in Tc gave rise to a greatly prolonged time to complete the crystallization. For instance, it took at least 200 min to complete the isothermal crystallization in the nanostructured thermoset while the temperature was enhanced to 89 °C. The conversions of crystallization [viz. X(t)] as functions of crystallization time (t) is shown in Figure 13B and the characteristic sigmoid X(t) ∼ t curves were exhibited at all the Tc’s. The curves were significantly shifted right side along the time axis with increasing the temperature of crystallization, suggesting that crystallization processes became progressively slow. This observation implies that the nucleation was a ratecontrolling step under the present conditions. The kinetics of isothermal crystallization of plain PE, PE in the triblock copolymer and in the nanostructured thermosets was analyzed by using Avrami equation:74 X t = 1 − exp( −K nt n)

(4)

where Kn is the overall kinetic rate constant and n is Avrami exponent that depends on the type of nucleation and on the geometry of growing crystals. The values of Kn and n were obtained from the intercept and the slope of the linear plots of ln[−ln(1 − X(t))] as a function of ln t (Figure 13C). It is seen that the experimental results were well coincident with the linear relation for the early part of transformation. From the sigmoid X(t) ∼ t curves, the times of half-crystallization (t1/2), defined as the time required for half of the crystallinity to

Figure 13. Plots of isothermal crystallization kinetics for the nanostructured thermoset containing 50 wt % of PCL-b-PE-b-PCL. (A) DSC curves of isothermal crystallization at various Tc’s; (B) plot of the conversions of crystallization [X(t)] as functions of crystallization time (t); (C) plots of ln[−ln(1 − X(t))] versus ln(t).

develop, were determined. The values of n, Kn, t1/2 and Tc’s are summarized in Table 2. The values of Kn (or t1/2) decreased 2750

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within the nanodomains and the crystals of PE would cross through the PCL microdomains. As a consequence, the crystallization of PE blocks was in an unconfined (or breakout) manner. In fact, we indeed noted that the undercooling degree of PE in PCL-b-PE-b-PCL was quite close to that of control PE and that the Avrami exponents of the isothermal crystallization for PE in the triblock copolymer were as high as 3.0. In marked contrast to the case of PE crystallization in the triblock copolymer, the crystallization of PE in the nanostructured thermosets was in a confined manner. In the nanostructured thermosets, the spherical PE microdomains were isolated and dispersed in the continuous epoxy matrix. As in polystyreneblock-polyethylene (PS-b-PE) diblock copolymer,75 the spherical PE nanophases were surrounded by epoxy matrix which possessed high rigidity and thus was capable of confining the crystallization within the nanoscaled space. This judgment was readily confirmed by the measurements of Tg’s by means of dynamic mechanical thermal analysis (DMTA) (see Figure 8). For the nanostructured thermosets containing up to 40 wt % PCL-b-PE-b-PCL triblock copolymer, the Tg’s of the epoxy matrices were not lower than the temperatures of crystallization for PE in the nanostructured thermosets, suggesting that before the crystallization of PE microdomains, the epoxy matrices have already undergone the rubber-to-glass transitions, i.e., the crystallizations of PE nanophases were confined with the nanoscaled spaces by the rigid epoxy matrices. For the nanostructured thermoset containing 50 wt % PCL-b-PE-bPCL triblock copolymer, the confinement was still preserved although the Tg of epoxy matrix at this composition was 79.5 °C and quite close to the Tc of PE (viz. 73 °C; see Figure 12). Nonetheless, the continuous matrix was a highly cross-linked network, which can significantly hinder the breaking out of PE crystallization. This case was in marked contrast to that of PCLb-PE-b-PCL triblock copolymer. In addition to the effect of epoxy matrix on the crystallization of PE in the nanostructured thermosets, the chemical linkage between PCL and PE blocks of triblock copolymer could also constitute the restriction of PE crystallization in the nanostructured thermosets. In the PE nanophases, two ends of each PE midblock were anchored onto epoxy matrix via the chemical linkages between PE and PCL blocks. As a consequence, the mobility of PE chains near the interfaces between PE microdomains and epoxy matrices were greatly restricted and thus crystallization was hindered. For the same reason, the PE nanophases displayed the much low crystallinity compared to the plain PE and the PE in the triblock copolymer (see Table 2).

Table 2. Crystallization Temperatures (Tc), Avrami Exponent (n), Constant of Crystallization Rate (Kn) and Half-Crystallization Time (t1/2) samples PE

PCL-b-PE-b-PCL (100 wt %)

PCL-b-PE-b-PCL (50 wt %)

PCL-b-PE-b-PCL (40 wt %)

Tc (°C)

n

122 122.5 123 123.5 124 118 118.5 119 119.5 120 87 87.5 88 88.5 89 86 86.5 87 87.5 88

3.76 3.67 3.85 3.62 3.86 2.89 2.91 3.13 3.01 3.17 1.73 1.78 1.80 1.74 1.82 1.33 1.59 1.62 1.71 1.85

K

t1/2 (min)

0.346 0.143 7.17 × 1.78 × 2.38 × 0.345 0.282 0.151 91.2 × 38.4 × 13.6 × 10.3 × 5.48 × 2.22 × 1.55 × 62.0 × 19.6 × 11.3 × 8.75 × 4.13 ×

1.21 1.54 3.28 5.20 14.3 1.27 1.36 1.63 1.96 2.49 9.71 10.6 14.7 27.1 28.5 6.15 9.42 11.8 13.0 16.0

10−3 10−3 10−5

10−3 10−3 10−3 10−3 10−3 10−3 10−3 10−3 10−3 10−3 10−3 10−3

with increasing Tc, indicating that the isothermal crystallization of PE in these samples became progressively slow. Theoretically, the value of n = 4.0 can result from spherulitic growth from sporadic nuclei whereas n = 3.0 can be attached to the spherulitic growth of instantaneous nuclei. For plain PE, the n values obtained were close to 4.0, suggesting a threedimensional spherulitic growth of sporadic nuclei. For PE in the triblock copolymer, the Avrami exponents were decreased to about 3.0, suggesting that the crystallization process may be the spherulitic growth of instantaneous nuclei. Notably, the isothermal crystallizations of PE in the nanostructured thermosets were quite different from plain PE and PE in the PCL-b-PE-b-PCL triblock copolymer. The n values for PE in the nanostructured thermosets were obtained to be in the range of 1.3−1.8 for the isothermal crystallization at 86 and 89 °C. The fact that the n values were smaller than 2.0 indicates that the growth of PE crystals in the thermosets was approximately in a one-dimensional manner, i.e., the confined crystallization occurred. Interpretation of Confined Crystallization. The crystallization of crystalline subchains in block copolymer is quite dependent on intermicrodomain connectivity, mechanical states and glass transition temperatures of the continuous matrices.75−78 For the PCL-b-PE-b-PCL triblock copolymer, the results of SAXS (See Figure 4) showed that in the melt state this copolymer was microphase-separated, in which the PE cylindrical phases were dispersed in the continuous PCL matrix with hexagonal lattice. However, the unconnectivity among the adjacent PE microdomains did not result in the confined crystallization of PE. The cooling DSC scan (See Figure 11) showed that the crystallization for PE in the triblock copolymer occurred at ∼112 °C, which was much higher than the melting temperature of PCL (∼60 °C). This observation suggests that PCL microdomains are in the molten state while PE blocks undergo the crystallization at 112 °C. It is proposed that the PCL melts were unable to confine the crystallization of PE

■

CONCLUSIONS In this work, we first synthesized polycyclooctadiene (PCOD) via the ring-opening metathesis polymerization (ROMP) of cyclooctadiene (COD) with cis-2-butene-1,4-diol diacetate as the chain transfer agent (CTA), which was catalyzed by Grubbs second generation catalyst. The terminal acetoxy groups of the PCOD were hydrolyzed into hydroxyl groups. The α,ωdihydroxyl-terminated PCOD with the molecular weight of Mn = 11,000 was then hydrogenated into a α,ω-hydroxylterminated polyethylene. With the α,ω-hydroxyl-terminated polyethylene as the macromolecular initiator, the ring-opening polymerization of ε-caprolactone was carried out and the poly(ε-caprolactone)-block-polyethylene-block-poly(ε-caprolactone) (PCL-b-PE-b-PCL) triblock copolymer was obtained. The PCL-b-PE-b-PCL triblock copolymer was incorporated into epoxy and the nanostructured thermosets containing PE 2751

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nanophases were obtained via self-assembly approach followed by curing reaction. The formation of PE nanophases was evidenced with small-angle X-ray scattering (SAXS) and transmission electronic microscopy (TEM). Depending on the content of PCL-b-PE-b-PCL triblock copolymer, the PE blocks in the thermosets formed the spherical, worm-like nanophases. Wide angle X-ray diffraction (XRD) showed that the formation of PE nanophases did not alter the structure of crystals. It is identified that the crystallization of PE blocks in the nanometer scale was in a confined manner. The confinement has been interpreted on the basis of the interdomain connectivity and the cross-linked structures of epoxy matrices.

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AUTHOR INFORMATION

Corresponding Author

*(S.Z.) Telephone: 86-21-54743278. Fax: 86-21-54741297. Email: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The financial supports from Natural Science Foundation of China (No. 51133003 and 21274091) and National Basic Research Program of China (No. 2009CB930400) were gratefully acknowledged. The authors thank the Shanghai Synchrotron Radiation Facility for the support under Project Nos. 10sr0260 and 10sr0126.

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