High-Strength Triple Shape Memory Elastomers from Radiation

Article Cite This: ACS Appl. Polym. Mater. 2019, 1, 1735−1748

pubs.acs.org/acsapm

High-Strength Triple Shape Memory Elastomers from RadiationVulcanized Polyolefin Elastomer/Polypropylene Blends Aizezi Maimaitiming,†,‡ Maojiang Zhang,†,‡ Hairong Tan,†,‡ Minglei Wang,†,‡ Mingxing Zhang,†,‡ Jiangtao Hu,† Zhe Xing,† and Guozhong Wu*,† †

Shanghai Institute of Applied Physics, Chinese Academy of Sciences, Shanghai 201800, China University of Chinese Academy of Sciences, Beijing 100049, P. R. China

‡

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S Supporting Information *

ABSTRACT: High-strength and high-temperature triple shape memory elastomers from multiphase polyolefin elastomer (POE)/polypropylene (PP) blends are prepared via radiation-induced vulcanization. Irradiation generates cross-linking mainly in the amorphous sequences of POE chains, long chain branches in the amorphous regions of PP crystals, and grafting between PP and POE chains at the interface of the POE and PP phases. The discoid-like PP phase with an average size of 1.6 × 11 μm2, dispersed in the POE matrix, is aligned parallel to the plane of the recompressed films. The PP lamellar crystals is partially orientated, and the PP lamellar crystals with (040)α and (110)α planes are arranged parallel and perpendicular to the surface of the highly cross-linked films, respectively. The integration of heterogeneous macromolecular structure, partially orientated crystals, and separated multiphase morphology contributes to excellent triple shape memory effects and high tensile strength up to 38 MPa. Consequently, two temporary shapes with fixity ratios of 82 and 97% are programmed by the reversible aggregation and solidification of the segregated PP lamellar crystals and POE bundle-like crystals. Under thermal triggering at the switching temperatures of 90 and 180 °C, the entropy-driven elasticity of the discrete PP and POE switching segments provides a good shape recovery of 88 and 94% to the memorized complex permanent shape. These results indicate that designing and tailoring the unconventional macromolecular architecture and condensed-state structure provide a novel strategy for developing low cost, high performance, and flexible intelligent polyolefins. KEYWORDS: polypropylene, polyolefin elastomer, radiation, orientation, multiphase morphology, confined crystallization, triple shape memory polymer segregation.8 The phase separation behavior not only heavily influences the crystallization behavior of each component and final properties of the blends9−15 but also plays a critical role in the triple shape memory effect (T-SME).4,16−19 The highdensity polyethylene (HDPE)/polypropylene (PP) blend with a cocontinuous architecture, stabilized with a 2,5-dimethyl-2,5di(tert-butylperoxy)hexane, shows more pronounced T-SMEs than that with a sea−island morphology.17 Excellent T-SME is characterized by the in situ compatibilized thermoplastic polyurethane (TPU)/olefin block copolymer (OBC)-g-glycidyl methacrylate/poly(ε-caprolactone) (PCL) blends, in the presence of dicumyl peroxide, with a fine cocontinuous morphology rather than the unmodified counterpart and that containing dispersed elongated particles and cavities.18 For the radiation cross-linked polyolefin blends, soft phase-rich blends exhibit excellent dual-SME or T-SME rather than both crosslinked rigid phase-rich and non-cross-linked blends.19−21

1. INTRODUCTION Driven by potential high-technological applications, the shape memory polymers (SMPs) have attracted considerable fundamental research interest in both academia and industry. As an important class of intelligent materials, SMPs are capable of changing their shapes in a predefined way with response to specific triggering (e.g., heat, light, or water). SMPs consist of stable networks and switching segments. The chemically or physically cross-linked networks can store the “memorized” permanent shape, while the switching segments can fix one or more temporary shapes by reversible aggregation and solidification of switching domains, including crystals and liquid crystals. The unique properties of SMPs have extensive applications in heat shrinkable tubes and films, intelligent medical devices, controlled drug delivery, membranes, selfhealing materials, sensors, actuators, 4D printing, aerospace, and the fabric industry.1−5 As a low-cost, efficient, and easy scalable method, polymer blending is widely used in preparation of triple SMPs.6,7 For most of the polymer blends, the thermodynamic incompatibility of each component contributes to phase separation and © 2019 American Chemical Society

Received: March 29, 2019 Accepted: May 30, 2019 Published: May 30, 2019 1735

DOI: 10.1021/acsapm.9b00289 ACS Appl. Polym. Mater. 2019, 1, 1735−1748

Article

ACS Applied Polymer Materials Moreover, a new “predictive material” from lightly cross-linked syndiotactic PP is tailored via dynamic coupling of crystal phase transition between mesomorphic trans-planar crystals and helical modifications with relaxation process.22 During the past decades, controlling the macromolecular architecture and hierarchical structure has been used to design and prepare next-generation high-value polymers.23−34 The physically cross-linked polyurethanes containing 15−40 wt % isolated hard segment domains show better dual shape recovery, and the shape fixity was tightly related to the crystallinity and strain-induced crystallization of the soft segment.4,35 The polymer networks based on covalent-bondlinked immiscible PCL (50−60 wt %) and poly(cyclohexyl methacrylate) segments show a favorable T-SME during the one-step dual shape programming process.36 Bilayer polymers consisting of two epoxy dual SMPs are tailored for T-SMP; well-separated thermal transition, a strong interface, and a relative ratio between the microphases are key factors for the T-SME.37 Reversible T-SME is achieved by designing and preparing polymer networks containing star-shaped precursors of polypentadecalactone and PCL segments with phase segregation; the two segments occur chain orientation when external stress is applied, and the cooling rate and applied stress during cyclic thermomechanical test markedly influence the T-SME.38 The development of metallocene and nonmetallocene catalysts contributes to manufacturing highperformance thermoplastic elastomers, such as polyolefin elastomers (POEs) and OBCs.39 As compared with other rubbers (e.g., natural rubber and ethylene−propylene−diene rubber, EPDM), the thermal, mechanical, and rheological properties of the POEs can be tailored by controlling the 1octene content, in situ generating the long chain branches (LCBs), and regulating the molecular weight, respectively.13,39 The unique characteristics of the POEs further provide a facile and flexible strategy to design and prepare high-performance POE/PP-based thermoplastic polyolefins (TPOs) and thermoplastic vulcanizates (TPVs).10,13,39−48 The final properties of TPOs and TPVs depend on their multiphase morphology, which is influenced by the molecular structure and composition ratio of raw materials, viscosity ratio of components, shear force, and manufacturing process.10,13−15,41−48 As light-weight and recyclable counterparts, these materials are widely used as compatibilizer, interior and exterior parts of automobile, household appliances, wire and cable, footwear, adhesives, and photovoltaic packaging.13,14,39,42,47 From the larger commercial distribution point of view, the POEs are supplied as pellets for easy handling, fast mixing, and compounding, which allows for batch and continuous production of TPO, TPVs, and SMPs.42 Consequently, significant research efforts are devoted to developing dual-, triple-, and multiresponsive and multifunctional SMPs from the inexpensive blends and composites of the POEs via chemical and radiation cross-linking, supermacromolecular hydrogen bonding and controlling the cocontinuous and elongated particle morphology.16,21,49−51 The radiation crosslinked microstructure also significantly influences the SMEs. As absorbed dose increases, the cross-linked density increases and, correspondingly, the shape fixity decreases, whereas the shape recovery increases.19−21,49 Moreover, thermo-, electro-, and solvent-stimuli-responsive SMPs with flake-like crystal morphology are prepared by melt-mixing of POE, carbon black, and lauric acid.52 It is a challenging task to prepare SMPs from a typical TPVs with a sea−island morphology. In situ

compatibilization of EPDM and PP phases by zinc dimethacrylate improves the rubber/plastic interface significantly, which further enhances the efficiency of stress delivery between the two phases. As a result, the dual-SME of TPVs is fulfilled.53,54 However, these SMPs have a narrow range of shape transition temperatures, lower elasticity and mechanical strength, and lack of high-temperature resistance and melt strength, which limit their commercial developments. High-energy radiation has been widely used to manufacture heat-shrinkable tubes and films and sterilize medical devices.3,55 Recently, radiation is applied for improving the interfacial adhesion of carbon fiber and matrix resin via grafting,56 preparing TPVs via dynamic cross-linking,47,48 synthesizing intelligent hydrogels for drug delivery, and curing composites and coatings via polymerization.57,58 Electron beams (EB) from high-current, high-voltage (80 kV−10 MeV) electron accelerators and γ-rays from cobalt-60 sources are primely used in the industrial radiation process. The dominant radiation chemistry employed in industrial applications proceeds via free radical mechanisms, which involves cyclic free radical reactions of initiation, propagation, and termination. EB and γ-rays present huge differences in dose rates (EB: 5−100 kGy/s; γ-rays: 1−10 kGy/h), maximum penetration depth (EB: 38 mm from 10 MeV; γ-rays: 300 mm), and distribution in space.58−60 Despite this, EB and γrays can penetrate into the amorphous phase, interphase, and the crystal of HDPE and PP, and all the bonds are cleavable, homogeneously producing macroradicals.59 Consequently, the two irradiations are considered equivalent for formation of secondary alkyl macroradicals and simultaneous H2 emission and creation of macromolecular defects (trans-vinylene, vinyl, and trans,trans-diene) in the initiation stage at low doses.61 However, the formation rate of macroradicals is a function of the concentration of the groups present in the polymeric chain (CH3, CH2, and CH), the energy of the bond that is broken, and steric hindrance. After irradiation, the macroradicals in polymers with higher crystallinity migrate from the crystalline phase and interphase to the amorphous phase. The stability and migration of macroradicals depend on chemical structure and morphology of polymer and temperature and solubility and migration of oxygen and additives. The interactions of macroradicals with chain defects, chain ends, and branches as well as oxygen and additives are diffusion controlled (kinetic reactivity), which influences the lifetimes of the macroradicals.59 Consequently, the radiation effects depend on chain structure and morphology of the polymer,62,63 radiation source (EB, γ-rays, X-rays, and heavy ion),58,61,64 radiation condition (beam energy,47 absorbed dose,65 dose rate,66 atmosphere,63 and temperature67), and post-treatment.64 For example, at 10−60 kGy, EB generate a star-like long-chainbranching topology in PP, whereas γ-rays give a tree-like one, which possesses decreased and increased melt-strain-hardening behavior with increasing Hencky strain rates.64 In this work, we designed fine segregated morphology of extrusion moltencompounded 70POE/30PP blends via controlling the viscosity ratio of POE and PP components and subsequently stabilized the morphology using EB and γ-ray cross-linking at doses of 50−200 kGy. After recrystallization of PP and POE crystals and rearrangement of cross-linked networks, high-strength, high-temperature resistance and flexible triple-SMPs are prepared. The structure−property correlation is established by using corresponding characterization and measurement methods. 1736


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ACS Applied Polymer Materials

time was 3 frames × 10 min. All the X-ray images were corrected for background scattering, air scattering, and beam fluctuations. Differential Scanning Calorimetry (DSC). DSC measurements were performed on a TA Q200 differential scanning calorimeter calibrated with indium. Pieces of samples of 8−10 mg were used for testing. The melting and crystallization temperatures (Tm and Tc) were considered as the peak temperature of the first heating and cooling curves of DSC, respectively. The characterization was performed at a standard heating and cooling rate of 10 °C min−1 under a flowing N2 atmosphere. All the samples were isothermally kept for 3 min at the end of each step to avoid thermal history. The crystallinity values of PP and POE were calculated via eq 1

2. EXPERIMENTAL SECTION Materials. Yarn grade isotactic polypropylene (YP) was purchased from Sinopec Shanghai Company (grade name: PPT30S; MFR = 3.0 g (10 min)−1 (230 °C (2.16 g)−1); Mw = 387 kg mol−1, Mw/Mn = 3.3; density = 0.901 g cm−3). High-melt-strength grade isotactic polypropylene (HP) was purchased from LCY Chemical Corp. (grade name: PT182; MFR = 0.4 g (10 min)−1 (230 °C (2.16 g)−1); Mw = 505 kg mol−1, Mw/Mn = 5.4; density = 0.903 g cm−3). POE was purchased from Dow Chemical Company (Mw = 161 kg mol−1; Mw/ Mn = 1.84; 1-octene content = 8.7 mol %, determined by 13C NMR). An amorphous ethylene−propylene−diene monomer (EPDM) was used for comparison. These materials were used without further purification. Sample Preparation and Irradiation. The POE/YP blend in the composition of 70/30 (wt %) was prepared in a corotating twinscrew extruder at 220 °C, with a screw diameter of 16 mm and a L/D ratio of 40 (HAAKE Polylab OS Rheomex PTW16/40, Thermo Electron GmbH, Karlsruhe, Germany). POE/HP and EPDM/HP blends in the composition of 70/30 (wt %) were prepared using a laboratory roll mixer at 180 °C. The pellets of POE/YP and thick sheets of POE/HP and EPDM/HP were preheated on a plate vulcanization press at 190 °C under 0.3 MPa for 10 min and then compressed into sheets of 1.8 mm thickness at 190 °C under 10 MPa for 5 min, followed by cooling the melt to room temperature. For γ irradiation, the molded sheets of 1.8 mm thickness were packed in a vacuum package and irradiated at an ambient temperature with different doses of 50−200 kGy in a 60Co source. For electron beam irradiation, the sheets of 1.8 mm thickness were packed in a vacuum package and irradiated at an ambient temperature with 1.25 MeV electrons and with total doses of 50−200 kGy by using an accelerator of the type DD-1.5 (Shanghai Institute of Applied Physics, China). The total doses were applied on the two surfaces in n steps of 10 kGy to avoid the penetration effect and to minimize the temperature rise resulting from the irradiation. After irradiation, the sheets were annealed at 80 °C for 4 h, followed by annealing at 100 °C for 4 h. After annealing, the sheets were recompressed under the same condition into films of 0.7−0.8 mm thickness. The irradiated samples were labeled as POE/(HP/YP)-X (X = dose, kGy); e.g., POE/HP100 indicates a sample prepared by radiation vulcanization of POE/ HP blend at a dose of 100 kGy. Gel Content Analysis. Sol−gel analysis was performed to determine the gel content of the compression molded sheets, according to ASTM D2765-01. Approximately 0.3 g of the sample was wrapped in a 120-mesh stainless-steel cage and subjected to extraction for 8 h in refluxing xylene. Samples were then dried under vacuum at 60 °C to a constant mass and weighed. Gel content was then calculated as a ratio of final weight to initial weight of the sample multiplied by 100. Morphology. The phase morphology of the films was performed in a Zeiss cross beam 540 field-emission scanning electron microscope. The films were cryo-fractured in liquid nitrogen, and the fractured surfaces were emerged in xylene at 80 °C for 4 h to etch the POE matrix. Then, the samples were dried at 60 °C for 24 h. Finally, the fractured surfaces were coated with thin layers of platinum using a sputter coater prior to SEM measurement. The average size of discoid PP phase was calculated by using at least 200 PP droplets in two or three SEM images. X-ray Diffraction and Scattering Testing. X-ray diffraction (XRD) was performed with Cu Kα (λ = 0.1542 nm) radiation. The powder profiles were obtained with an autonomous Bruker D8 Advance diffractometer performing a continuous scan at the diffraction angles 2θ of 5°−40° at a scanning rate of 2θ of 0.02° (2θ) s−1 and a total scanning step of 1711. Two-dimensional X-ray scattering measurements (2D-WAXD and 2D-SAXS) were performed ̀ at the laboratory beamline Anton Paar SAXSpoint 2.0 compacts (Xian Jiaotong University). The wavelength λ (Cu Kα) of the radiation source was 0.1542 nm, and the power was 50 W (50 mV/1 mA). Scattering patterns were collected by a EIGER R 1M hybrid photon counting detector with a pixel size of 79 μm2. The image acquisition

Xc =

ΔHm × 100% ΔHm°

(1)

where Xc is the crystallinity (%), ΔHm is the melting enthalpy of PP or POE (J g−1), and ΔH°m is the extrapolated value of the melting enthalpy of a 100% crystalline sample (209 J g−1 for PP, 290 J g−1 for POE).47 The apparent crystallinity of PP in blends was calculated by dividing Xc by the weight fraction of PP. Tensile Stress−Strain Testing. Tensile properties were measured by using an Instron 5943 at 23 °C and a testing speed of 50 mm min−1. Compression remolded films were cut into dumbbell-shaped specimens of 35 mm length, 2 ± 0.1 mm width, and 0.8 ± 0.1 mm thickness. At least five samples were used for measurement tensile properties of each sample, and the average values of the stress and corresponding strain of at least three samples were plotted as stress− strain curves. Dual and Triple Shape Memory Effects (D-SMEs and TSMEs). SMEs were measured by cyclic heating−deformation− cooling−heating process using strain recovery process, visual observation, and dynamic mechanical thermal analysis (DMTA) as reported in the literature.50,51 The thermal transition temperatures Ttran1 of 90 °C and Ttran2 of 180 °C were selected according to DSC measurement. The detailed experimental determination of dual and triple shape memory effects by strain recovery process and visual observation is provided in the Supporting Information or corresponding figures. The shape fixed ratio (Rf) and shape recovery ratio (Rr), and the influence of the programming and recovery conditions on the TSMEs, were determined by DMTA. For the programming process, rectangle samples of 15 × 2.5 × 1.65 mm3 size underwent the isothermal process at 180 °C for 3 min before they were stretched to a εS1 under a stress of 0.01−0.05 MPa within 2 min. The temporary shape B was fixed by cooling the samples to 90 °C at a cooling rate of 10 °C min−1 under constant stress, followed by an isothermal process for 5−10 min, and then the samples were further maintained isothermally for 5 min to determine the fixed εf1 and Rf(A→B); a stress of 0.15−0.5 MPa was employed to deform the samples to a εS2 within 2 min. The temporary shape C was fixed by cooling the samples to 0 °C at a cooling rate of 10 °C min−1 under constant stress, followed by isothermal for 5−10 min, and then the samples were further maintained isothermally for 5 min to determine the fixed εf2 and Rf(B→C). For the recovery process, the samples were heated to 90 °C at a heating rate of 10 °C min−1 and maintained isothermally for 10 min to determine the recovered εf1 and Rr(C→B); then, the samples were heated to 180 °C at a heating rate of 10 °C min−1 and maintained isothermally for 10 min to determine the recovered εf0 and Rr(B→A). The Rf(x→y), Rr(y→x), and total recovery Rr(total) were calculated via eqs 2−4. εfy − εfx R f(x → y) = εsy − εfx (2) R r(y → x) =

R r(total) = 1737

εfy − εrx εfy − εfx

(3)

εs2 εs2 − εf0

(4) DOI: 10.1021/acsapm.9b00289 ACS Appl. Polym. Mater. 2019, 1, 1735−1748

Article


vulcanized POE/HP-200 film only swells in hot xylene, instead of being etched. This is in consistent with the correlation between gel content and absorbed dose. The interface of dispersed PP phase and POE matrix gradually disappears, and the size of the PP phase slightly decreases with an increase in the dose and gel content (Table 1) because radiation results in in situ grafting between POE and PP phases, which increases the compatibility of POE with PP. The POE-g-PP chains were evidenced by extraction of the POE/YP-0 blend and γvulcanized POE/YP-50 film in toluene at 90 and 140 °C using a 120-mesh stainless-steel cage. At 90 °C, all macromolecules in the POE/YP-0 blend and only soluble POE macromolecules in the POE/YP-50 film are extracted. All macromolecules in the POE/YP-50 film are extracted in the boiling toluene at 140 °C. This result indicates that POE-g-PP chains are in situ generated by radiation at the interlamellar stacks or interspherulitic boundaries of the PP phase.13−15 To investigate the reproducibility of the aligned morphology in different mixing process and subsequent compression, the extrusion molten compounding of the POE/YP-0 and roll mixing of the POE/HP-0 are used, considering the difference of ηPOE/ηPP. As shown in Figure 1a,b, the mixing processes do not significantly influence the multiphase morphology of vulcanized elastomers. This is because the viscosity of YP and HP is much higher than that of POE, and no significant phase coalescence or transition occurs. However, a thicker and larger elongated discoid-like morphology of YP phase is observed in the core layer of the POE/YP-0, as shown in Figure S1. The reason for generating this heterogeneous morphology is unclear, yet it may due to the viscosity mismatch of the POE and YP components and the difference of cooling and shear rate in the core and skin layers of the sheet during compression. Gisbergen et al. found that the 80PP/20EPDM blend experienced a phase transition from sea−island to cocontinuous morphology during injection molding.69 Ono et al. reported that split, cylindrical, and ellipsoid-like morphologies were observed for the 70PP/30POE blends with different ηPOE/ηPP.12 Svoboda et al. studied the morphology and elastic property of the PP/POE blends in the whole composition. For the POE/PP blends in the composition of 80/20 and 70/30 wt %, the PP phase with sphere (size: 0.1−1 μm) and elongated shape (smaller dimension: 0.5−1 μm; larger dimension: 2−10 μm) were dispersed in the POE matrix. For the 60POE/40PP blend, the PP phase was more elongated, and the morphology was a cocontinuous one.13 Svoboda et al. also systematically studied the multiphase morphology and crystallization behavior of the 50PP/50POE blends, containing POE with different 1-octene contents and molecular weights. A 50PP/ 50POE blend containing moderate-molecular-weight POE shows elongated stripe-like morphology.41 Crystal Structure. Figure 2 and Figure S2 show the X-ray powder diffraction profiles of the pristine POE/PP and EPDM/HP blends and their corresponding vulcanized elastomers. The common α form of PP is obtained in the pristine blends and vulcanized elastomers, as characterized by the presence of the (110)α, (040)α, (130)α, (111)α, (131)α, and (060)α reflections at 2θ of 14.13°, 16.93°, 18.58°, 21.31°, 21.95°, and 25.51°, respectively. The (111)α and (131)α reflections overlap with the (200)POE reflection of POE crystals. The (111) α , (131) α , and (110) α reflections disappeared in the γ-vulcanized EPDM/HP at doses of 100, 150, and 200 kGy and γ-vulcanized POE/YP blends at doses of 150 and 200 kGy are shown in Figure S2c and Figure 2b,

3. RESULTS AND DISCUSSION Multiphase Morphology. The thermodynamical immiscibility of POE and PP chains contributes to phase separation in the melt state before crystallization. It has been reported that multiphase morphology of immiscible blends depends on the composition of components, viscosity ratio (ηPOE/ηPP), interfacial adhesion (compatibility), and process conditions.15,42 The average size of dispersed rubber domain in the hard TPOs and soft TPVs significantly influences their thermal and mechanical properties.10,13,15,42 Fine dispersed rubber particles act as uniform heterogeneous nucleating sites that increase the crystallization rate of PP and generate more uniform and smaller crystals of PP.13,68 Smaller rubber particles also increase the processability and elastic properties of TPV.42−46 In this case, the melt viscosity of both PP components is higher than that of POE, and it is expected that PP dispersed in continuous POE matrix in an irregular two-phase morphology rather than sea−island morphology.12,13,15,69 Aiming to prepare high-strength T-SM vulcanizates with higher switching temperature without phase inversion, the phase morphology is stabilized using radiation cross-linking. The irradiated sheets are recompressed into films to orientate the PP phase and recrystallize the POE and PP components. The blends were cryo-fractured in liquid nitrogen followed by etching POE matrix in xylene at 80 °C, and surface morphology was observed with SEM. Figure 1 and Figure S1

Figure 1. SEM images for γ-vulcanized films: (a) POE/HP-100, (b) POE/YP-100, (c) POE/HP-150, and (d) POE/HP-200 (the bright bulged part: PP phase; the dark recessed part: cross-linked POE phase; the vertical direction is parallel to the plane of films).

show the SEM micrographs of the γ-vulcanized POE/HP films at doses of 100, 150, and 200 kGy and of the γ-vulcanized POE/YP-100 film. Interestingly, the discoid-like PP phase, dispersed in the cross-linked POE matrix, is highly oriented parallel to the plane of films. The morphology of discoid-like PP phase is confirmed by Figure 1c. The isotropic rod-like PP phase is observed in the fractured sides, perpendicular to the plane of γ-irradiated POE/HP films, in contrast to anisotropy PET microfibrillar highly orientated parallel to the draw direction in the PP matrix.70 The dispersed PP phases of 1.8 × 13.4, 1.5 × 11.4, 1.6 × 10.2, and 1.6 × 9.6 μm2 width and length, respectively, are observed for the corresponding samples. The moderately cross-linked POE phase in the γvulcanized POE/HP-100 film is significantly etched in hot xylene, whereas the highly cross-linked POE matrix in the γ1738


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Table 1. Crystallization and Thermal Behaviors of Pristine POE/(HP/YP) Blends and Their Electron Beam (E)- and γ-Ray (G)-Vulcanized Elastomers iPP

POE

sample

Tm,p (°C)

Tm,o (°C)

Xc (°C)

Tc,p (°C)

Tc,o (°C)

HP YP POE POE/HP-0 POE/HP-E50 POE/HP-E100 POE/HP-E200 POE/HP-G50 POE/HP-G100 POE/HP-G150 POE/HP-G200 POE/YP-0 POE/YP-E50 POE/YP-E100 POE/YP-E200 POE/YP-G50 POE/YP-G100 POE/YP-G150 POE/YP-G200

161.6 164.9

154.4 156.9

40.7 53.8

117 119.6

120.3 122.8

160.9

153.3

31.7

115.5

156.1 156 157.3 154.8 153.4 147.3 161.3

146.3 141.5 150.5 141.2 139.7 138.9 153.9

32 26 33.7 29 27.3 30 40.3

155.7 153.7 158 156.2 153.9 146.6

148.3 144.5 150.3 146.1 141.2 139.7

40.7 34.6 40.7 32.3 33.7 35.7

Tm,p (°C)

Tc,p (°C)

Tc,o (°C)

117.8

79 74.5

63.6 70.6

66.9 76.5

113 110.3 114.4 114.2 113.9 113.9 112.2

119.6 120.1 118.7 120.1 119.9 120.3 118

76.2 75.5 77.8 74.8 74.1 75.3 77

66.1 63.9 67.4 65.1 63.6 63.1 66.3

71.2 64.1 72.3 70.9 69.5 68.9 70.9

111.4 109.5 112.6 106.5 107.6 96

117.4 121.6 118.1 118.7 121.9 112.6

77.6 77.9 78.4 76.8 76.7 76.5

61.9 59.1 62.8 60.9 59.9 58.8

66.9 64.1 68.3 67 65.9 65

[gel] (%)

0 0 68.7 83.8 0.8 82.6 87.7 89.7 0 0 68 83.8 0 80.6 89.7 91.5

Figure 2. X-ray powder diffraction profiles of electron beam (a) and γ-vulcanized (b) blends of polyolefin elastomer and yarn grade polypropylene (YP).

Figure 3. 2D-WAXD (left) and 2D-SAXS (right) patterns of pristine and γ-vulcanized POE/YP films (absorbed doses: a, 0 kGy; b, 50 kGy; c, 100 kGy; d, 200 kGy). 1739


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Figure 4. 1D-WAXD (a) and 1D-SAXS (b) curves of the pristine POE/YP blend and γ-vulcanized POE/YP films at doses of 50, 100, and 200 kGy.

respectively. The crystal structure of POE is characterized by the (110)POE and (200)POE reflections at 2θ of 21.88° and 23.5°. Electron beam, γ irradiation, and absorbed dose strongly influence the intensity of all crystallographic plane reflections, especially for the intensity of the (040)α and (110)α plane reflections. The increase in the absorbed dose results in a significant increase in the intensity of the (040)α reflection and a decrease in the intensity of the (110)α reflection, when measured from the parallel direction to the plane of films. The apparent crystal size (D) of the PP phase in the samples in the direction perpendicular to the (110)α and (040)α planes is calculated via eq 5:71 Dhkl =

Kλ B0 cos(θhkl)

Crystal structure and morphology were changed in the ultrathin, highly constrained polymer layers of microscale or nanoscale thickness.24,74 The flattened spherulites or discoids of PP were generated in the coextruded microlayer films with alternated PP and polystyrene layers. The size of flattened spherulites increased with the decrease in the layer thickness of PP, from about 18 μm for the PP control film to 34, 38, and 42 μm for 460, 108, and 65 nm films, respectively. The edge-on lamellae with the (040)α plane was lying flat on the interface of films.74 The folded surfaces of the PE lamellae in the polyolefin block copolymers, composed of glassy poly(cyclohexylethylene), elastomeric poly(ethylene-alt-propylene), and semicrystalline PE, were perpendicular to the draw direction or microphase-separated lamellar interfaces (domain spacing: 20 nm).75 In this case, the average thickness of 1.6 μm for discoid-like PP phase, dispersed in the highly cross-linked POE matrix, is much lower than the dimension of 18−30 μm of common PP spherulites.74 The crystallization kinetics and temperature of PP in the 50POE/50PP blends with stripe-like morphology are much lower those that of the same blends with cocontinuous morphology because striped structure generated larger obstacles to growing PP lamellae.13 Zeng et al. observed that the crystallization of the LCB-PP under the fast cooling condition, or at a lower crystallization temperature, leads to the formation of predominant edge-on lamellar structure. At a high crystallization temperature of 145 °C, flat-on lamellae with chain axis aligned perpendicular to the film plane were observed.76 In the radiation-vulcanized elastomers, the LCBs are in situ generated in the PP phase, and POE-g-PP chains are generated at the interface of POE and PP phases. It is deduced that the chain entanglement of LCBs leads to chain orientation parallel to the film plane in the PP phase,77 which may result in the orientation of lamellar crystals with (040)α plane parallel to the plane of films. The cross-linked networks and POE-g-PP chains lead to extension of PP chain in the longer interface of PP and POE phase, perpendicular to the film plane, which may contribute to the orientation of the lamellar crystals with (110)α plane. Ono et al. found that the PP lamellar crystals are orientated normal to the slab-like POE domains in the 70PP/30POE blend, while the PP lamellar crystals penetrates randomly into the circularlike POE domains.12 Moreover, considering the characteristic of the lower dose rate of γ radiation, the probability of migration of trapped macroradicals and subsequent branching between the macroradicals and double-bond-terminated macromonomers at the interface of PP and POE phases may

(5)

where B0 is the half-width in radians of the reflection corrected for instrumental broadening and λ is the wavelength of the Cu Kα radiation (λ = 0.1542 nm). The shape factor K used here is equal to 0.89. The calculated values of D[(110)α] for the POE/YP-0 and γ-vulcanized POE/YP-50, POE/YP-100, and POE/YP-200 samples are 21.59, 18.87, 16.51, and 17.79 nm, respectively. The calculated values of D[(040)α] for the corresponding samples are 20.59, 21.49, 19.21, 20.13, and 20.44 nm, respectively. The results indicate that the crystal growth in the directions perpendicular to the (110)α (c-axis) become smaller, as the absorbed dose increases. Figure 3 shows the 2D-WAXD and 2D-SAXS patterns for the POE/YP0 and γ-vulcanized POE/YP-50, POE/YP-100, and POE/YP200 films. The YP phases in the POE/YP-0 and POE/YP-50 show a monocyclic 2D-WAXD pattern and isotropic 2D-SAXS scattering behavior, and no orientation is found according to the azimuthal angle distribution, as shown in Figure S3. However, the γ-vulcanized POE/YP-200 film shows arcs in the 2D-WAXD pattern and elliptical anisotropic shape in the 2DSAXS pattern, indicating that the YP lamellar crystals are partially orientated. In contrast to the XRD, the 1D-WAXD curve of the POE/YP-200 film in Figure 4 shows that the (040)α reflection disappears. The intensity of the (110)α reflection in all samples does not significantly change. The XRD and 1D-WAXD results indicate that the PP lamellar crystals with the (040)α and (110)α planes are orientated parallel and perpendicular to the plane of the POE/YP-200 film. This result can be interpreted by the crystallization under the condition of spatial confinement and topological constraints of the LCBs and POE-g-PP chains.72−74 1740


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Figure 5. DSC nonisothermal crystallization (a, c) and melting (b, d) curves of pristine blends (POE/HP-0 and POE/YP-0) and electron beam (a, b) and γ-vulcanized (c, d) elastomers from polyolefin elastomer with high-melt-strength (HP) and yarn (YP) grade polypropylenes.

understanding of crystallization behavior and thermal property is necessary to determine the programming and recovery conditions. It is well accepted that a well-separated thermal transition is one of the key factors for T-SME.16,36−38 In this context, DSC characterization as a vital tool, showing the crystallization and melting behaviors, can indirectly reveal the macromolecular microstructure. Figure 5 shows the nonisothermal crystallization and melting behaviors of the compression remolded, pristine POE/(YP/HP) blends and their vulcanized elastomers. Table 1 summarizes the Xc of PP, crystallization, and melting temperatures of each component. As expected, the Xc of the PP component in the blends is lower than that of pure PP due to the influence of phase separation on the subsequent crystallization. The melt of the continuous POE matrix decreases both the concentration of crystallizable PP chains and the number of the diluted PP chains diffusing onto the crystallization growth front.79 As a result, the concentration dilution and obstruction effects suppress the crystallization of some PP chains located at the interface of POE and PP phases. However, the isothermal crystallinity of the PP phase in the 50PP/50OBC blend with cocontinuous morphology is higher than that of pure PP, and the isothermal crystallization rate of PP phase in the 70PP/30OBC blend with a spinodal morphology is higher than that of the pure PP phase.79 The irradiation source and applied dose do not significantly influence the melting behavior of the POE, which indicates that three-dimensional cross-linked network is generated mainly in the amorphous random ethylene/1-octene sequences

be much higher than that of EB. This is proved by the higher gel contents (Table 1) and shifted crystallization and melting curves in the γ-irradiated samples, as mentioned in the next section. As a result, the topological constraint efficiency in the γ-irradiated samples is much higher than that in electron-beamirradiated samples, which yields higher orientation of (040)α planes, as shown in Figure 2 and Figure S2. The apparent long space (L) values, calculated by using Bragg’s law, of the pristine POE/YP-0 and γ-vulcanized POE/YP-50, POE/YP100, and POE/YP-200 are 15.33, 14.45, 16.01, and 16.86 nm, respectively. The value of the crystalline lamellar thickness (Lc) is calculated via eq 6:71 Lc =

Xc L (ρc /ρa )(1 − Xc) + Xc

(6)

where the Xc is the apparent crystallinity of PP phase, measured by using DSC. ρc of 0.936 g cm−3 and ρa of 0.852 g cm−3 are the densities of the crystalline and amorphous PP phase.78 The values of Lc for corresponding samples are 5.83, 5.56, 4.85, and 5.66 nm. The calculated values of amorphous thickness La (= L − Lc) are 9.5, 8.89, 11.16, and 11.20 nm. The results indicate that the values of L and La increase with the absorbed dose. One can note that the presence of LCBs, POEg-PP, and cross-linking chains partially inhibited the crystal growth of the PP phase, which results in an increase of the thickness of the amorphous region of PP phase. This is in consistent with the DSC measurement. Nonisothermal Crystallization and Thermal Properties. For the thermally triggered SMPs, a fundamental 1741


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Tensile Properties. The orientation of phase and crystals in composites often gives significant improvement in mechanical and other properties.24,70,80 To meet the application-derived specific demands, complex requirements for properties (e.g., elasticity and thermal property) and functions (e.g., strength, biodegradability, and electrical conductivity) in addition to SME should be fulfilled.1 Table 2 summarizes the strength at 100% strain, elongation at break

of the POE chains, which agrees well with that of electron beam cross-linked POE.62 In the chemically cross-linked vulcanizates, the peroxide induces cross-linking in both crystalline and amorphous sequences of POE chains and reduces the concentration of larger crystals.16,42,46,62 However, the crystallization peak temperature decreased from 66.3 °C for the crystals of POE in pristine POE/YP-0 blend to 58.8 °C for that in the γ-vulcanized POE/YP-200 sample. This is because the cross-linking affects the folding and rearrangement of the crystallizable ethylene sequences in the POE chains. The broader exothermic peak at the lower temperatures is the characteristic of melting behavior of bundle-like crystals in the POE matrix. The crystals of POE gradually change from thick and thin lamellar crystals to mixed small lamellae and bundlelike crystals and then to fringed micellar or bundle-like crystals with a decrease of 1-octene content. The corresponding crystal morphology was characterized by a decreasingly well-defined spherulite morphology, unbanded spherulite, and the granular, nonlamellar morphology.40 Irradiation markedly decreases the crystallization rate and melting temperature of PP. It is wellknown that irradiation causes degradation and branching in the PP macromolecules. However, the changes in the molecular weight, molecular weight distribution, and branching topologies of PP depend on the employed irradiation source and absorbed dose.64 At a dose lower than 60 kGy, electron beam and γ irradiation yield a decrease in the molecular weight of PP but result in different star- and tree-like LCB topologies in the amorphous regions of PP crystals. Because of the different absorbed dose rate, the two irradiations generate distinct grafting between POE and PP chains. The decrease in the molecular weight of PP leads to a higher crystallinity of PP, and in situ grafting between POE and PP chains enhances the compatibility of POE with PP phase. These two factors result in a uniform crystallization of PP phase (crystallization curves of POE/HP-50 and POE/YP-50, Figure 5c) and the higher crystallinity of the PP phase, as listed in Table 1. In contrast, at a dose higher than 100 kGy, the crystallization peak and melting temperatures of PP phase shift to the lower temperature. The crystallization peak temperature decreased from 112.2 °C for the PP crystals in the pristine POE/YP-0 to 96 °C for that in the γ-vulcanized POE/YP-200 sample. The melting temperature decreased from 161.3 °C for the PP crystals in the pristine POE/YP-0 to 158, 156.2, 153.9, and 146.6 °C for that in the γ-vulcanized POE/YP elastomers at doses of 50, 100, 150, and 200 kGy, respectively. Interestingly, all the samples irradiated at 200 kGy show a double melting peak of PP. The high dose results in higher concentrations of LCBs and stereodefects in PP phase, which result in decrease in the lamellar size. The topological constraint in the PP phase and cross-links in the POE matrix, surrounding the moderately branched PP phase, also prevent the rearrangement and diffusion of the PP chains. As shown in Figure 5a,c, the electron-beam-vulcanized POE/YP elastomers show more discrete crystallization and melting temperature than the γvulcanized counterparts, which is beneficial to pronounced shape memory effect. Moreover, the baseline above the melting temperature of PP phase in the γ-vulcanized POE/YP-200 is higher than that of other samples. This result indicates that the recompressed film may have shape memory effects that recover from the temporary shape of the film to the permanent shape of the sheet and stores more energy during recompression than the other sample.

Table 2. Tensile Properties of Pristine POE/(HP/YP) Blends and Their Electron Beam (E)- and γ-Ray (G)Vulcanized Elastomers sample POE/HP-0 POE/HP-E50 POE/HP-E100 POE/HP-E200 POE/HP-G50 POE/HP-G100 POE/HP-G150 POE/HP-G200 POE/YP-0 POE/YP-E50 POE/YP-E100 POE/YP-E200 POE/YP-G50 POE/YP-G100 POE/YP-G150 POE/YP-G200

strength at 100% strain (MPa) 7.3 9.8 10.9 10.7 10.0 8.5 9.3 9.5 6.7 10.2 11.0 10.5 12.2 10.5 11.0 11.6

± ± ± ± ± ± ± ± ± ± ± ± ± ± ± ±

0.1 0.2 0.2 0.4 0.6 0.5 0.4 0.4 0.2 0.4 0.3 0.4 0.8 0.3 0.1 0.5

elongation at break (%) 1101 994 1084 811 1031 933 709 721 1183 950 1045 885 957 878 723 648

± ± ± ± ± ± ± ± ± ± ± ± ± ± ± ±

17 35 21 29 71 70 35 23 9 25 34 45 22 28 14 42

strength at break (MPa) 12.5 22.4 33.3 32.6 25.2 30.6 36.6 31.7 12.1 20.2 33.2 33.3 25.6 38.1 37.4 35.0

± ± ± ± ± ± ± ± ± ± ± ± ± ± ± ±

0.2 1.2 1.0 1.0 2.9 2.0 3.2 2.5 0.2 0.9 1.9 1.3 1.2 1.5 1.9 1.5

and tensile strength of pristine blends and their vulcanized elastomers. Figure 6 and Figure S4 show the stress−strain curves as functions of the irradiation type and absorbed dose. Figure S5 shows the stress−strain curves as functions of the irradiation type and absorbed dose, after shape recovering from the temporary shape of the film to the permanent shape of the sheet. As shown in Figure S4, the nonrecompressed POE/YP-0 and electron beam irradiated POE/YP sheets at doses of 50, 100, 150, and 200 kGy show a nonuniform strain response under stretching, lower tensile strength, and necking behavior as compared with that of films, as seen in Figure 6c. This indicates that the recrystallization of POE and PP and orientation of PP lamellae during recompression are essential for improving tensile properties. As shown in Figure 6a,c, the pristine POE/HP-0 film behaves as a hard elastomer, whereas the POE/YP-0 exhibits a strain softening zone, necking zone, and weaker strain hardening zone, which are characteristics of a plastomer. Moreover, the tensile strength and elongation at break of vulcanized elastomers nonmonotonically change with the absorbed dose. The strength at 100% strain slightly increases with the dose. The high strength of the vulcanized elastomers can be interpreted by the unique macromolecular architecture, orientated morphology, and partially orientated crystals. From the morphology of the elastomers shown in Figure 1, the elastomers can be seen as a highly orientated rigid PP discs reinforced POE composite. The increase in the absorbed dose corresponds to the increase in the cross-linking and orientated degree of lamellar crystals with the (040)α plane (Table 1 and Figure 2). Consequently, the covalent bonding of the partially orientated lamellae of PP and bundle-like crystals of POE with the amorphous POE matrix provides the high 1742


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Figure 6. Tensile stress−strain curves of the electron beam (a, c) and γ-vulcanized (b, d) elastomers from POE/HP and POE/YP blends.

results demonstrate that heterogeneous macromolecular structure, partially orientated crystals, and highly orientated multiphase morphology during compression comprehensively contribute to the high tensile strength of the vulcanized elastomers. Dual Shape Memory Effect. The D-SME was determined by the strain recovery process. The elastomers exhibit distinct dual shape effects at the moderate and high switching temperatures due to the presence of segregated POE bundlelike crystals and PP lamellar crystals. The temporary shapes C and B were programmed at 90 and 180 °C by stretching the dumbbell-shaped or rectangle specimens to a strain of 100− 500% under stress, followed by cooling to 20 °C for fixing the temporary shapes. The permanent shape A was recovered from the temporary shapes B and C by heating at the switching temperatures of 90 °C (Ttran1) and 180 °C (Ttran2). Figures S6−S9 show the dual shape memory effects of the pristine POE/YP blend and selected vulcanized samples. The pristine POE/YP blend shows much higher average strain fixed ratio (Rf) of 95%, but lower strain recovery ratio (Rr) of 75% when recovered from a strain of 350% at Ttran1. In this pristine blend, the permanent shape is only stabilized by PP crystals, and the POE and PP phases do not link with chemical bonds. As seen in Figure S6, the POE/YP-0 blend exhibits lower hightemperature resistance and shows macroscopic shape deformation. This is because the low melt strength cannot prevent partial irreversible slip or flow of POE chains when stretched at 90 °C. Interestingly, the electron-beam-vulcanized POE/HP50 sample exhibits a balance between Rf of 95% and Rr of 91% at Ttran1. This result demonstrates that a dual shape memory polymer can be made via blending POE with PP without crosslinking, but in situ grafting between POE matrix and dispersed PP phase is necessary to enhance their compatibility and melt

mechanical strength of the elastomers, which is impossible to realize in common TPVs.42,46 An exception is that the values of strength at 100% strain of the γ-vulcanized POE/YP-50 and POE/HP-50 films are higher than that of other samples. As sheen in Figure 6b,d, the orientation of the lamellae with the (040)α plane, lying on flat in the PP discoid of POE/YP-50 and POE/HP-50 films, is stronger than that of the γ-vulcanized POE/YP-100 and POE/HP-100 samples. The compatibility of POE with PP phase is enhanced by radiation-induced in situ grafting between POE and PP phases. In addition, the chain entanglements of the LCBs, generated at the low dose of 50 kGy, and higher crystallinity of the low-molecular-weight LCBPP dispersed in the POE matrix may also enhance the strength at 100% strain of the samples. To confirm the influence of the bundle-like crystals of the POE matrix on the unconventional strain hardening property, an amorphous EPDM elastomer is used for comparison. As expected, the γ-vulcanized EPDM/ HP-150 film shows a similar strength at 100% strain to that observed for POE/HP-150 film, but it shows a lower tensile strength and no strain hardening property, as shown in Figure 6b. Although the PP phase in the POE/HP-150 shows higher crystallinity, its lower interfacial adhesion with the noncrystalline EPDM phase contributes to a lower tensile strength. The partial crystal orientation significantly influences the tensile property and strain hardening behavior of highly cross-linked samples. As shown in Figure S5, the shape recovered and γvulcanized POE/YP-200 sheets show higher elongation at break and lower tensile strength. From 2D-SAXS testing, the shape recovery sheets show no crystal orientation (not shown). In contrast, the γ-vulcanized POE/YP-200 film shows lower elongation at break, higher tensile strength, and sharp strain hardening property, which correspond to the orientation and fragmentation of PP lamellar crystals during stretching. These 1743


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Figure 7. Triple shape memory effects of the γ-vulcanized POE/HP-100 elastomers (A and A′: the original permanent shape; B and B′: the first temporary shape; C: the second temporary shape; D and E: the transition shapes during C−B′ and B′−A′ recovery; the A−B−C−B′−A′ cycle includes flower fully blooming-half closed-fully blooming-fully closed steps).

Scheme 1. Schematic Description in Real Time of the Triple Shape Memory Effects of the γ-Vulcanized POE/HP-100 Elastomersa

A and A′: original permanent shape in fully closed flower model; B and B′: the first temporary shape in fully blooming flower model; C−C′: The second temporary shape in half-blooming flower model; D: the transition shape in half-closed flower model during B′−A′ recovery. a

indicates that the shape recovery is three-dimensional that can be used as a special sealant material. Triple Shape Memory Effects. In general, T-SMEs were observed in the one-component physical cross-linked block copolymers and covalent or noncovalent bonded multiphase blends. As mentioned in the Introduction, the T-SMPs contain two segregated crystallizable phases, glass transition phases or liquid crystalline transition phases, by which two temporary shapes can be programmed. Consequently, the thermal transition temperatures Ttran can be the melting temperature, glass temperature, or liquid crystalline transition temperature. When heated above Ttran, stepwise recovery to the permanent shape can be achieved.1,3,51 For the block copolymers, the phase with highest Ttran can stabilize the permanent shape, which is beneficial for resetting the permanent shape.1,2,35 For the vulcanized elastomers, the permanent shape A is stabilized by the macromolecular cross-linked networks, in situ generated by radiation, at the amorphous sequences of POE and the interface of POE matrix and dispersed PP phase. The

strength. The samples irradiated at a dose higher than 100 kGy possess much higher Rr of 97−100% at Ttran1 because the permanent shape is stabilized by both covalent bonded networks in the POE matrix and physical cross-links of PP lamellar crystals. The Rf of 95% and 86% are exhibited by the electron beam vulcanized POE/YP-200 and POE/HP-200 samples at the strains of 100% and 200%, respectively. The lower Rf of POE/HP-200 under a higher strain is due to the better elasticity and higher cross-linked density. The pristine POE/YP-0 blend and electron beam vulcanized POE/HP-50 sample are not proper to measure their dual shape memory effects at Ttran2 because of their much lower stability at this temperature. The samples irradiated at a dose higher than 100 kGy show a higher Rf of 97−100% and a Rr of 100% at Ttran2. The high Rf and Rr can be attributed to the reversible thermal transition of both POE and PP crystals. It should be noted that the recompressed films of 0.8 mm thickness recover to the sheets of 1.8 mm thickness, when heated at Ttran2, which 1744


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Figure 8. Time-dependent triple shape memory effects of the γ-vulcanized POE/YP-100 (a) and POE/YP-200 (b) elastomers determined by DMTA.

temporary shape B was programmed at Ttran2 of 180 °C and fixed by the confined lamellar crystals of PP at Ttran1 of 90 °C. The temporary shape C was programmed at Ttran1 and fixed by the bundle-like crystals of POE at 0 or 20 °C. The T-SMEs were semiquantitatively determined by the strain recovery process as displayed in Figure S10. The total recovery ratios Rf(C→B) of γ-vulcanized POE/HP-100 and POE/YP-100 and electron-beam-vulcanized POE/HP-100 and POE/HP-200 samples are 98%, 100%, 94%, and 97%, respectively. The T-SMEs in blooming/closed flower for the γ-vulcanized POE/HP-100 and POE/HP-200 samples, determined by visual observation, are shown in Figure 7 and Scheme 1. The temporary shapes B and C were programmed by rotating the petals 180° and 270° from the permanent shape A. As heated at Ttran1, the half-bloomed flower of POE/HP-100 opened and recovered from the temporary shape C to fully open flower with temporary shape B. As heated at Ttran2, the fully opened flower with temporary shape B closed and recovered completely to the permanent shape A. More importantly, the T-SMEs of the elastomers are accompanied by changes in the light transparency. The SMPs are relatively transparent in the permanent shape A. When heated at T > Ttran1, the samples become opaque due to the melting of fringed-micellar crystals of POE. Upon further heating to T > Ttran2, the lamella crystals of PP are totally molten and the materials become completely transparent. The shape recovery in SMP is driven by the entropy elasticity of switching segments. In the amorphous, rubberelastic state above Ttran, the polymer chains exhibit a completely random coil conformation. As a mechanical deformation is applied, the polymer chains align along deformation direction and state in the low entropy. If the orientated chains are cooled under the constant stress to a temperature below Ttran, the temporary shapes are fixed by aggregation and solidification of switching segments, which prevents the immediate recovery of the polymer chains. The entropy-driven recovery of permanent shape takes place when the switching segments are softened again at Ttran.1,2 Consequently, the T-SMEs of the elastomer are not only influenced by the internal macromolecular architecture and segregated crystal structures of POE and PP phases but also by external mechanical deformation.38 To study these effects, the time-dependent T-SMEs of the γ-irradiated POE/YP-100 and POE/YP-200 were determined by DMTA and are plotted in Figure 8 and Figure S11. The Rf(A→B) and Rf(B→C) of POE/YP100 and POE/YP-200 are 82.3, 40% and 97.2, 96.6%, respectively. The Rr(C→B) and Rr(B→A) of POE/YP-100 and POE/YP-200 are 94.2, 87.6% and 87.6, 76.9%, respectively.

The much lower Rf(A→B) of 40% for POE/YP-200 can be attributed to shorter deformation time and higher cross-linked density. As stretched, lower entropy loss corresponds to the transformation of random coil conformation of the POE/YP200 sample to a relatively lower orientated conformation and lower crystallinity. As shown in Figure 8a,b, the deformation stress, applied for programming the highly cross-linked chains into temporary shapes B and C, is much higher for POE/YP200 than that for POE/YP-100. Moreover, the longer crystallization time gives higher shape fixity for POE/YP-100, as shown in Figure S11. As a result, the entropy loss during deformation at a given strain is lower. Moreover, lower strain during deformation leads to lower chain orientation, lower storage energy, and lower strain-induced crystallization, which results in lower shape recovery ratio. Faster heating and longer isothermal time at Ttran2 mean faster strain recovery. The LCBs and grafting between POE matrix and dispersed PP phase in the POE/YP-200 is higher than that in the POE/YP-100, and the crystallization of PP takes more time to be completed, as shown in Figure 5c. Consequently, a higher Rf(A→B) of 70% for the POE/YP-200 is achieved under a lower strain at longer isothermal crystallization condition (not shown). It should be noted that R′r(C→B) of 66.5%, calculated by using ε′r1, for the POE/YP-200 sample is higher than that of 54.8% for the POE/ YP-200 sample. The total shape recovery ratios of POE/YP100 and POE/YP-200 are 94.9 and 98.6%, respectively. The results indicate that higher cross-linking gives better shape recovery, as reported in the literature.19,20 The presence of Stype recovery (εf2−ε′r1−εr1) of temporary shape C to its permanent shape demonstrates that additional crystals are formed under melt stretching and subsequent cooling. The crystals with higher thickness are molted at a higher temperature. In addition, the entropy-driven shape recovery by melting the fringed-micellar crystals of POE requires longer time to be completed.

4. CONCLUSION The preparation of high-strength triple shape memory elastomers with segregated and unconventional orientated morphology is reported. The 70POE/30PP blends are molten compounded by using the extrusion and roll mixing process. Subsequent radiation of the compression-molded sheets in situ generates unique macromolecular architectures: covalent bond linking of the cross-linked amorphous region of POE matrix with segregated crystallizable ethylene sequence (randomly distributed in soft POE domains) and LCB-PP chains. Radiation cross-linking leads to a significant decrease in the 1745


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ACS Applied Polymer Materials crystallization peak temperature of POE in blends, whereas it does not influence melting temperature. As the absorbed dose increases, crystallization and melting temperatures of PP in the γ-vulcanized elastomers markedly decrease. The electronbeam-vulcanized POE/YP elastomers show more discrete crystallization and melting temperature than the γ-vulcanized counterparts, which is beneficial to pronounced triple shape memory effect. At low shape transition temperature, the vulcanized elastomers without gels exhibit pronounced dual shape property with shape fixity and recovery ratios of 91−95% because the temporary shape is fixed by bundle-like crystals in soft POE domains and interface strength of the POE and PP phases is enhanced; at high shape transition temperature, the vulcanized elastomers with concentration of 68−92% gels exhibit excellent dual shape memory with shape fixity and recovery ratios of 97−100% due to the fact that temporary shape is fixed by both the bundle-like POE crystals and PP lamellar crystals. The triple shape memory effect heavily depends on the cross-linked macromolecular structure and programming and recovering conditions. The elastomers with concentration of 68−80% gels show enhanced T-SME but with lower recovery stress. As the absorbed dose increases, the shape fixity and recovery ratio decrease for first temporary shape fixed by PP lamellar crystals; the shape fixity and recovery ratio are high for the second temporary shape fixed by bundle-like POE crystals due to the strain-induced crystallization. The overall shape recovery of the elastomers is 94− 100%. The integration of elastomeric macromolecular networks with segregated bundle-like crystals of the POE phase and partially orientated lamellar crystals of the PP phase is attributed to better elasticity, higher tensile strength, lower elongation at break, and strain hardening behavior. Moreover, the elastomers show high-temperature resistance that can be deformed to a strain of 500% at the molten state. The threedimensional complex permanent shape can be stored by γ cross-linking, which may overcome the drawback of the processability of chemical cross-linked polyolefin blends as well as the unpleasant smelling and color of traditional TPVs. The whole process is expected to be efficient, eco-friendly, and easy to scale up and might be easily extended to multistimuli and multifunctional SMPs via finely dispersing the electronic conductive graphene and carbon nanotubes and magnetic nanoparticles in the pristine elastomers.

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determined by strain recovery (Figures S6−S9); triple shape memory effects of the γ-vulcanized POE/HP-100 and POE/HP-200 samples determined by strain recovery (Figure S10); DMTA observation of T-SMEs for γ-vulcanized POE/YP-100 (Figure S11) (PDF)

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel: (+86)-21-39194531. Fax: (+86)-21-39195118. ORCID

Guozhong Wu: 0000-0003-3814-2074 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the National Natural Science Foundation of China (Grants 11605275 and 11675247), the National Key R&D Plan of China (Grant 2016YFB0303004), and the Science Challenge Project (Grant TZ2018004).

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REFERENCES

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsapm.9b00289. Experimental section for strain recovery and visual observation; SEM images for the core layer of the γvulcanized POE/YP-100 film (Figure S1); XRD profiles for (POE/EPDM)/HP elastomers (Figure S2); azimuthal angle distribution of the pristine POE/YP-0 and γvulcanized POE/YP elastomers at doses of 50, 100, and 200 kGy (Figure S3); tensile stress−strain curves of unrecompressed pristine POE/YP-0 and electron-beamirradiated POE/YP sheets at doses of 50, 100, and 200 kGy (Figure S4) and γ-vulcanized POE/YP sheets, shape recovered from corresponding films (Figure S5); dual shape memory effects of POE/YP-0, electronbeam-vulcanized POE/HP-50 and γ-vulcanized POE/ HP elastomers at doses of 50, 100, 150, and 200 kGy, 1746


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High-Strength Triple Shape Memory Elastomers from Radiation

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