Elasticity Reinforcement in Propylene–Ethylene Random Copolymer

Article pubs.acs.org/Macromolecules

Elasticity Reinforcement in Propylene−Ethylene Random Copolymer Stretched at Elevated Temperature in Large Deformation Regime Jiayi Zhao,† Yingying Sun,‡ and Yongfeng Men*,† †

State Key Laboratory of Polymer Physics and Chemistry, Changchun Institute of Applied Chemistry, Chinese Academy of Sciences, University of Chinese Academy of Sciences, Renmin Street 5625, 130022 Changchun, P.R. China ‡ ExxonMobil Asia Pacific Research & Development Co., Ltd, 1099 Zixing Road, Minhang district, 200241 Shanghai, P.R. China S Supporting Information *

ABSTRACT: Tensile deformation behavior of a random propylene− ethylene copolymer with 12 mol % ethylene counits at room temperature and 63 °C was investigated using the in situ small and wide-angle X-ray scattering techniques. Under both conditions, the deformation mechanism changed from slip of crystalline lamellar blocks to stretching induced melting and recrystallization process at a critical strain of about 0.9. This critical strain in tensile deformation of semicrystalline polymers normally marks the starting of plateau value of elastic strain. Further stretching leads to increase of plastic deformation only due to the fibrillation. However, a peculiar elasticity reinforcement was observed at strain larger than 1.3 when the sample was stretched at 63 °C. Wide angle X-ray scattering results indicate that at this strain of 1.3 fibrillation of the originally unoriented crystals finished so that further stretching leads to a deformation of a rather stable entangled network embedded by fibrils that possesses high elasticity.

■

INTRODUCTION The deformation mechanism in semicrystalline polymers that are composed of crystalline lamellae and entangled amorphous chain in between is complex as both phases affect the process.1 Generally, semicrystalline polymers show both plastic and elastic deformation after being stretched. A separation of the total strain into elastic part and plastic one can be achieved via a so-called “step-cycle” test.2−10 It was found that the recovery properties of a semicrystalline polymer change at four critical strains (A, B, C and D) in the order of increasing in strain reflecting characteristic microstructural evolution in the system. Macroscopically, the systems show onset of plastic strain at A, a sudden increase in plastic strain at B, a plateau for elastic strain at C, and finally a decrease in elasticity at D. Such results can be understood by regarding a semicrystalline polymer as interpenetrated networks of hard crystalline skeleton and amorphous entangled soft network. In this way, the plastic deformation at small strains can be understood as a result of isolated intralamellar crystalline block slips (A), a change into a collective block slips (B), crystalline block disaggregation (melting)−recrystallization-induced fibrillation (C), and disentanglements (D). Notably, critical strain at B corresponds to the conventional yield point. Clearly, during a tensile stretching plasticity in a semicrystalline polymer occurs at A before yielding. Fibrillation starts at C that introduces more plasticity in the system as the fibrillar structure generated by the stretching-induced crystalline black disaggregation (melting)− recrystallization process always oriented with its long axis along the stretching direction. However, before disentanglements set © XXXX American Chemical Society

in, the apparent plasticity shown in the system during tensile deformation can be fully recovered by heating the sample over its melting point. Therefore, the strain at point D marks the onset of true plasticity in the system. The critical strains have been also confirmed by true stress−strain relationship investigations and in situ wide-angle X-ray diffraction (WAXD) measurements.3,4,10 Specially, the position of critical strain C is determined by the interplay between the stability of crystalline blocks and the state of amorphous entangled network. Only if the stretched amorphous entangled network reaches certain strain can it generate a critical stress leading to a destruction of the crystalline blocks followed by recrystallization. This two-phase model explains mechanical response of semicrystalline polymers with sufficient crystallinity. A change of the position of critical strain at C toward larger value was also discovered in a low crystallinity ethylene/octane copolymer that led to a more general three-phase model taking into account the heterogeneity of the distribution of the crystalline phase.11 In such system, the disaggregation (melting)−recrystallization process of crystalline blocks was delayed due to the heterogeneous strain distribution in the system. Such heterogeneity in structure as well as in macroscopic mechanical behavior brings more complicity in understanding the mechanical properties of low crystalline systems. In the late stage of tensile deformation, fibril bundles Received: November 29, 2015 Revised: January 3, 2016

A

DOI: 10.1021/acs.macromol.5b02585 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules

was derived from inverse Fourier transformation of the experimental intensity distribution I(q) as follows

together with entangled network were highly stretched, mostly resulting in cavitation caused by breaking interfibrillar tie molecules for highly oriented isotactic polypropylene12 with a critical stress for the initiation of stress-whitening depending on molecular weight but not on stretching temperature. On the other hand, stress-hardening behavior13 as well as wellprotected α daughter lamellae14 at higher stretching temperatures instead of cavitation have been observed for low crystallinity polyolefins at larger strains. In this work, we investigated the structure evolution at different temperatures of a propylene−ethylene copolymer of low crystallinity during tensile deformation. A similar delay of the onset of fibrillation was observed in this system as previously found in the ethylene/octene copolymer.11 Moreover, a peculiar elasticity reinforcement at larger strain was observed when the sample was stretched at elevated temperature. As will be shown in the following sections, this peculiar behavior in change in elasticity after large deformation can be linked to the microstructural feature of the system at large deformation, which is composed of stretched amorphous network and loosely packed fibrils at elevated temperature.

■

∞

K (z) =

EXPERIMENTAL SECTION

b0 b

∞

∫0 I(q) dq

(2)

which was used for computing the thicknesses of crystalline lamellae (dc), amorphous layers (da), and the long period (dac) of the samples composed of two-phase layer-like systems.17 For systems with a structure of stacks of lamellae, the correlation function shows characteristic features that allow the long spacing defined as the average thickness of a lamella together with one interlamellar amorphous layer measured along the lamellar normal to be determined. Because the weight crystallinity measured by DSC was less than 50%,18 the smaller value obtained from the correlation function was assigned to dc. da was then obtained as da = dac − dc. To investigate the orientation of crystalline and amorphous phases, the “halo method” proposed by Hsiao et al.19,20 was applied in data analysis of two-dimensional WAXD patterns. Fraser method21 has been used to correct the distortion of the pattern caused by flat-plate detector. The two-dimensional patterns can be divided into two fractions by standard procedure: isotropic part and anisotropic part. The azimuthal scan was first drawn along the angular axes starting from the center of the WAXD patterns. At each pixel position, a minimum intensity value was obtained from the azimuthal scan, which eventually yielded the isotropic part. The anisotropic part was obtained by subtracting the isotropic part from whole WAXD pattern. The anisotropic part was composed of oriented crystals and amorphous phase. And the isotropic part is composed of unoriented amorphous and crystalline phases. The quantitative evaluation of mass fraction of crystalline and amorphous phases (oriented and unoriented) was determined with a peak fitting method developed by us. Besides the in situ SAXS and WAXD measurements during tensile stretching, step-cycle and shrinkage heat retraction2 tests have been carried out to further understand the recovery properties and deformation mechanism. In the step-cycle experiment, the sample was stretched step-by-step at a constant speed. After each step, the crosshead speed was inverted to contract the sample until a stress of zero was achieved. Thereupon, the sample was extended again at this given speed, until it reached another preset larger strain. This combines a stepwise stretching of the sample with loading−unloading cycles. The imposed strains are therefore decomposed into a quasielastic (recoverable) part and a quasi-plastic (irreversible) part. In addition, shrinkage heat retraction is used to measure the truly irreversible plastic part which can roughly describe the extent of disentanglement and chain fracture. By heating the recovered freestanding strip up to a temperature causing the melting of all crystals, the sample recovered further toward its initial undeformed state due to the recovery of the stretched entangled network. An incomplete recovery after shrinkage heat retraction indicates true plastic deformation due to disentanglements or chain fracture during tensile stretching. Differential scanning calorimeter (DSC) measurements were carried out with a DSC1 Star System (Mettler Toledo Swiss) which had been calibrated for temperature and melting enthalpy by using indium as a standard under a N2 atmosphere with a heating rate of 10 K/min. The samples were heated up at a constant rate of 10 K/min.

The random propylene−ethylene copolymer iPPcoE12 was produced using metallocene catalyst and supplied by ExxonMobil Asia Pacific Research & Development Co. with a trade name of Vistamaxx 3980fl. The melt flow rate (MFR) is 8.5 g/10 min (230 °C/2.16 kg). The number-average (Mn) and weight-average molecule weight (Mw) are 94 and 195 kg/mol by GPC. The weight concentration of ethylene is 8.3%. The material was first compression-molded into a film of 0.5 mm in thickness at 180 °C and held at this temperature for 5 min to remove thermal history. Crystallization occurred during cooling the film to room temperature in air. Some of the amorphous chain slowly crystallized after aging at room temperature. To compare the effect of aging time, samples used in mechanical tests have been stored at room temperature for 10 days and 180 days. The “dog bone” tensile bars with dimensions of 10 × 5 × 0.5 mm3 were obtained with the aid of a punch. The tensile tests were carried out using a portable tensile testing machine (TST 350, Linkam, U.K.). In order to measure the strain of the deformed area, optical photo images of the sample were employed to measure the strain of the sample during testing. The Hencky measure of strain εH was used as a basic quantity of the extension, which is defined as

εH = 2 ln

∫0 I(q) cos(qz) dq

(1)

where b0 and b are the widths of undeformed and deformed area, respectively. Synchrotron small-angle X-ray scattering (SAXS) and WAXD measurements were performed at the beamline 1W2A, BSRF, Beijing. The wavelength of X-ray radiation was 0.154 nm. The sample-todetector distance was 3070 mm for SAXS, and was 139 mm for WAXD. The effective range of the scattering vector q (q = 4π(sin θ)/λ, where 2θ is the scattering angle and λ is the wavelength) was 0.078− 1.050 nm−1 for SAXS. A 2D MAR CCD X-ray detector with a resolution of 2048 × 2048 (pixel size = 79 μm) was used to acquire 2D-SAXS and 2D-WAXD patterns. We used a stepwise tension at a constant cross-head speed of 20 μm/s at deformation temperatures of 30 and 63 °C. The primary X-ray beam was first positioned at the middle of the horizontally placed sample strip. The SAXS and WAXD patterns were collected after every step for 100 s. The SAXS patterns were background calibrated and normalized using the standard procedure. Then one-dimensional scattering intensity distributions were integrated within ±10°along the horizontal direction of SAXS patterns. For the type of data treatment (Lorentz correction versus no correction) only plays a minor role,15,16 the data was calculated using Lorentz correction. The correlation function K(z)

■

RESULTS AND DISCUSSION Figure 1 shows DSC melting curves of iPPcoE12 aged at room temperature for different time. The position and enthalpy of the endothermic peak located at around 78 °C in each melting curve were not affected by aging time, indicating it was the final melting peak. For the one at lower temperature, the melting peak shifted to higher temperature and its enthalpy increased for longer aging time, which was due to the slow crystallization of some amorphous chain segments which cannot crystallize all at once during cooling from the melt.22 B

DOI: 10.1021/acs.macromol.5b02585 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules Xα = Xcrys − X γ

(6)

The bottom of Figure 2 illustrated the total crystallinity, crystallinity of γ phase and α phase as a function of aging time at room temperature. The crystallinity of γ phase increased slightly while the crystallinity of α phase almost remained the same, indicating the slow epitaxial growth25 of γ phase occurred during all the aging time. Figure 3 shows true stress−strain curves of iPPcoE12 aged for 10 days and 180 days stretched at room temperature and 63

Figure 1. DSC melting curves of iPPcoE12 aging at room temperature for different time. Heating rate: 10 K/min.

The polymorph of iPPcoE12 consisted of α phase and γ phase. A peak fitting method was designed to fit the curve into several Gaussian peaks, and representative fitting performed on iPPcoE12 aging for 6 months was shown in Figure 2(top). The relative crystallinity (Xcrys) was calculated by Xcrys =

∑ Acrys ∑ Acrys + ∑ Aamor

(3)

Figure 3. True stress−strain curves of iPPcoE12 aging at room temperature for 10 days and 180 days stretched at room temperature and 63 °C.

where Acrys and Aamor are the fitted areas contributed by crystalline and amorphous phases, respectively. The fraction of γ phase was calculated by23 fγ = Iγ(117)/[Iγ(117) + Iα(130)]

°C. The choice of a stretching temperature at 63 °C was to remove the effect of crystallinity generated upon aging at room temperature. Clearly, when stretched at room temperature sample having been aged for longer time shows much higher modulus and stronger strain hardening behavior than that of aged for shorter time. This aging effect mostly disappeared when that samples were stretched at elevated temperature of 63

(4)

where Iγ(117) and Iα(130) are the areas of the diffraction peaks. The crystallinity of γ phase (Xγ) and α phase (Xα) was given by24

X γ = fγ Xcrys

(5)

Figure 2. Integrated 1D WAXD curve of iPPcoE12 aging for 6 months before deformation (top) and plots of crystallinity, amorphous, content of α and γ achieved from WAXD results as a function of aging time at room temperature (bottom). C

DOI: 10.1021/acs.macromol.5b02585 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules °C as most of crystals formed during aging had been melted before stretching. A close inspection of the stress−strain curves at 63 °C revealed that the sample aged at 10 days showed a slight stronger strain hardening at large strain side. The behavior suggests that the 180 days sample possessed a slightly smaller number of junction points (crystallites) therefore lower network modulus at 63 °C. We speculate that this can be a result of structural difference in the samples upon aging for different period of time. With a long time aging, further crystallization of the system can lead some previously isolated crystallites (individual junction points) to be connected forming effectively slightly less number of such points. To understand the deformation mechanism of iPPcoE12 stretched at different temperatures, in situ SAXS and WAXD measurements during stretching have been carried out at room temperature and 63 °C. Selected WAXD (top) and SAXS (bottom) patterns of iPPcoE12 aging for 180 days stretched at 63 °C were given in Figure 4. The transition from isotropic

Figure 5. Plots of da, dc, and dac achieved from integrated intensity curve as a function of strain εH.

structural perfection during aging. For samples aged for 180 days, it possessed a much better interconnected structure due to secondary crystallization that helps to give a better force transmission keeping the long spacing of the corresponding lamellar stacks lager than that of the 10 days samples within a small range of strain. As long spacing represents local microstructural features of stacked lamellae not the complete structural features over the sample, it is not surprising that the difference was not reflected in the macroscopic stress−strain curves present in Figure 3. It is noted that the amorphous thickness and long spacing decreased at larger strain for higher stretching temperature while the value seemed to be invariant when stretched at room temperature. Another difference is that the thickness of the stretch-induced recrystallized lamellae at higher temperature was slightly higher than that at room temperature, following the stretch-induced crystallization line.28 In order to further investigate the recovery property of the iPPcoE12 sample, “step-cycle” measurements have been carried out at room temperature and 63 °C for iPPcoE12 aging for 180 days decomposed plastic and elastic strains as a function of total strain obtained from such step-cycle tests were presented in Figure 6. At the beginning of the tensile tests, elastic part and plastic part both increased where the plasticity came from the crystalline lamellar slips. The deformation mechanism changed from crystalline block slip to disaggregation (melting)recrystallization process at critical point C located at around 0.9 at both temperatures. This critical strain of around 0.9 is much higher than that of pure isotactic polypropylene29 indicating a similar situation as the low crystallinity ethylene/ octene copolymer of three phase model.11 The plastic strain increased at a larger slop and elastic strain started to decrease slightly at strain larger than 0.85 due to fibrillation process at room temperature. Difference recovery behavior was observed at 63 °C that the elastic part remained constant between 0.9 and 1.3 and increased at strain larger than 1.3, presenting strong elasticity reinforcement occurred in the last stage of tensile stretching. Clearly, a change in deformation mechanism occurred at the strain of 1.3. In order to characterize the true plastic deformation of the system after stretching, heat shrinkage experiments can be carried out where the irreversible strain εH,i obtained after the stretched sample was slowly heated up to its molten state present true plastic deformation. The value can roughly

Figure 4. Selected WAXD (top) and SAXS (bottom) patterns of iPPcoE12 aging for 180 days stretched at 63 °C.

diffraction circle to highly oriented diffraction point shown in WAXD patterns and similar orientational behavior at crystalline lamellar level represented in SAXS patterns indicated that the inner structure of the sample was highly oriented at the end of stretching. Figure 5 depicts the variation of long spacing (dac), lamellae thickness (dc), and amorphous thickness (da) along stretching direction derived from SAXS pattern for iPPcoE12 aging for 10 days and 180 days stretched at different temperatures. Because of the negligible difference in crystal density of α phase and γ phase,26 the system was regarded as a two-phase structure where correlation function1 can be used. The original lamellar thickness remained the same after long time storage. In general, the long spacing increased at small strain which was due to the elongation of the amorphous phase and then decreased later caused by fibrillation via a melting−recrystallization process.16,27 However, evolution of long spacing for sample aged for 180 days stretched at 63 °C showed no obvious decrease after an initial increase giving a notable larger value than the sample aged for 10 day stretched at the same temperature. This phenomenon can be linked to the difference in extent of D

DOI: 10.1021/acs.macromol.5b02585 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules

Figure 6. Variation of plastic (εH,p) and elastic (εH,e) parts as a function of total strain stretched at room temperature (top) and 63 °C (bottom).

Figure 7. Evolution of irreversible part (εH,i) as a function of total strain stretched at room temperature (top) and 63 °C (bottom).

characterize degree of chain scission and disentanglement in the system. As shown in Figure 7, the system cannot recover entirely at strain of 1.0 and beyond stretched at room temperature. In addition, the recovery property of the system has been damaged to a large extent after a plateau region. It is obvious that amorphous entanglement network was easier to be weakened due to high chain mobility30 at 63 °C. εH,i became stable after strain of 1.6 which indicates further stretching had no destruction on the fibril structure generated during fibrillation process. Further evidence of deformation mechanism can be found in the variation of mass fraction for crystalline and amorphous phases obtained from WAXD patterns in tension as shown in Figure 8. In case of room temperature stretching, the mass fraction of the four parts (unoriented and oriented crystalline and amorphous phases) fluctuated slightly at early stage. Apparent decrease of both unoriented parts and increase of both oriented parts illustrated the system gradually became more ordered even though around 30% unoriented amorphous

part still existed at the end of stretching. The mass fraction of isotropic crystalline phase continued to decline during stretching and became 0 at strain of 1.4, indicating isotropic crystalline phase has been depleted during elongation marking the end of fibrillation process. Mesophase was observed at strain of 1.6, instantly after exhaustion of unoriented crystalline phase, which was attributed to crystal−mesophase transition. Oriented mesophase which is composed of oriented chains with only partial packing ordering is generated by pulling chains out of the crystal.31 The higher entropy32 of it due to randomness of packing and hands of helices helps to stabilize the system and reduce the recovery tendency. In addition, some neighboring chains may also be restricted by oriented mesophase with high stability.33 Thus, the elastic strain decreased in the late stage. The bottom of Figure 8 shows the situation of sample stretched at 63 °C. An apparent difference by comparison of the two cases is that mesophase was not produced at high stretching temperature. The mass fraction of unoriented crystalline phase vanished at strain of 1.3, which coincides E

DOI: 10.1021/acs.macromol.5b02585 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules

Figure 8. Change tendency of fraction of oriented crystal/amorphous, unoriented crystal/amorphous and oriented mesophase part with increasing strain of iPPcoE12 aging for 180 days stretched at room temperature (top) and 63 °C (bottom).

with macroscopic deformation behavior that fibrillation process stopped at the same strain. According to model proposed by Tang et al.34 the elongation of fibrils is regarded as stretching of the amorphous chain between the lamellae and slippage of microfibrils which are composed of stacks of lamellae with their normal parallel to the stretching direction. The coupling force between microfibrils increases during stretching due to interfibrillar segments being stretched taut thus further sliding of microfibrils leading finally to slippage of fibrils built by bundles of microfibrils. Immediately, one observes the activation of slippage of fibrils at strain of 1.3 due to enhanced cohesiveness of fibrils brought about by perfect microfibrillar packing at higher stretching temperature.35The recovery tendency of the system came from the tie chains connecting fibrils. The slight increase in εH,i, shown in Figure 7, and oriented crystalline phase, shown in Figure 8 between a strain of 1.3 and 1.6, implied that further stretching still resulted in crystallization of the system. Production of the γ phase by melt crystallization confirmed the short crystallizable sequence in metallocene-catalyzed propylene copolymer.36 Quite a large number of the chain sequences in the amorphous region were short because only the repeating units with a number that was larger than the critical value can crystallize. It is reasonable that the short chain sequences in interfibrillar tie chains were highly stretched at larger strain and worked as the nucleus of straininduced crystallization.37

was the slippage of microfibrils, which leads to the formation of fibrils and the polymorph transition. Unusual elasticity reinforcement observed at high stretching temperature was mainly due to the slippage of fibrils after depletion of unoriented crystal since the cohesion between the microfibrils was much enhanced.

■

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.5b02585. Raw data of “step-cycle” experiments at 30 and 63 °C and photographs of samples during “step-cycle” tests (PDF)

■

AUTHOR INFORMATION

Corresponding Author

*(Y.M.) E-mail: [email protected]. Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS This work is supported by NSFC (21134006) and ExxonMobil. We thank Dr. X. Chen and Dr. R. Wittenbrink at ExxonMobil Asia Pacific Research and Development Company, Ltd., for helpful discussions and Prof. Zhonghua Wu and Dr. Guang Mo for assistance during synchrotron X-ray scattering measurements.

■

CONCLUSIONS In summary, we observed a peculiar elasticity reinforcement at the late stage of tensile deformation for a low crystallinity propylene/ethylene random copolymer stretched at elevated temperature. The oriented system at the late stage of tensile stretching can be viewed as a highly stretched network embedded by fibrils composed of microfibrils. At room temperature, crystal-mesophase transition occurred at strain of 1.6, resulting in the increasing speed of destruction for the entangled network and further decline of recoverability. Apparently, the deformation mechanism after critical point

■

REFERENCES

(1) Strobl, G. The physics of polymers, 2nd ed.; Springer: Berlin, 1997. (2) Hiss, R.; Hobeika, S.; Lynn, C.; Strobl, G. Macromolecules 1999, 32, 4390−4403. (3) Men, Y.; Strobl, G. J. Macromol. Sci., Part B: Phys. 2001, 40B, 775−796. (4) Al-Hussein, M.; Strobl, G. Macromolecules 2002, 35, 8515−8520. (5) Men, Y.; Rieger, J.; Strobl, G. Phys. Rev. Lett. 2003, 91, 095502. (6) Men, Y.; Strobl, G. Macromolecules 2003, 36, 1889−1898. F

DOI: 10.1021/acs.macromol.5b02585 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules (7) Al-Hussein, M.; Strobl, G. J. Polym. Sci., Part B: Polym. Phys. 2004, 42, 2074−2080. (8) Hong, K.; Rastogi, A.; Strobl, G. Macromolecules 2004, 37, 10165−10173. (9) Men, Y.; Rieger, J.; Hong, K. J. Polym. Sci., Part B: Polym. Phys. 2005, 43, 87−96. (10) Hong, K.; Strobl, G. Macromolecules 2006, 39, 268−273. (11) Sun, Y.; Fu, L.; Wu, Z.; Men, Y. Macromolecules 2013, 46, 971− 976. (12) Lu, Y.; Wang, Y.; Chen, R.; Zhao, J.; Jiang, Z.; Men, Y. Macromolecules 2015, 48, 5799−5806. (13) Mao, Y.; Burger, C.; Li, X.; Hsiao, B. S.; Tsou, A. H.; Mehta, A. K. Macromolecules 2011, 45, 951−961. (14) Mao, Y.; Li, X.; Burger, C.; Hsiao, B. S.; Tsou, A. H. Polymer 2013, 54, 1432−1439. (15) Jiang, Z.; Tang, Y.; Rieger, J.; Enderle, H.-F.; Lilge, D.; Roth, S. V.; Gehrke, R.; Heckmann, W.; Men, Y. Macromolecules 2010, 43, 4727−4732. (16) Jiang, Z.; Tang, Y.; Men, Y.; Enderle, H.; Lilge, D.; Roth, S. V.; Gehrke, R.; Rieger, J. Macromolecules 2007, 40, 7263−7269. (17) Tanabe, Y.; Strobl, G. R.; Fischer, E. W. Polymer 1986, 27, 1147−1153. (18) Men, Y.; Rieger, J.; Lindner, P.; Enderle, H. F.; Lilge, D.; Kristen, M. O.; Mihan, S.; Jiang, S. J. Phys. Chem. B 2005, 109, 16650− 16657. (19) Ran, S.; Zong, X.; Fang, D.; Hsiao, B. S.; Chu, B.; Ross, R. J. Appl. Crystallogr. 2000, 33, 1031−1036. (20) Ran, S.; Zong, X.; Fang, D.; Hsiao, B. S.; Chu, B.; Phillips, R. a. Macromolecules 2001, 34, 2569−2578. (21) Fraser, R. D. B.; Macrae, T. P.; Miller, A.; Rowlands, R. J. J. Appl. Crystallogr. 1976, 9, 81−94. (22) De Rosa, C.; Auriemma, F.; Perretta, C. Macromolecules 2004, 37, 6843−6855. (23) Turner-Jones, A. Polymer 1971, 12, 487−508. (24) Wang, Y.; Chen, C.; Xu, J. Z.; Lei, J.; Mao, Y.; Li, Z. M.; Hsiao, B. S. J. Phys. Chem. B 2012, 116, 5056−5063. (25) Cao, Y.; Van Horn, R. M.; Tsai, C.-C.; Graham, M. J.; Jeong, K.U.; Wang, B.; Auriemma, F.; De Rosa, C.; Lotz, B.; Cheng, S. Z. D. Macromolecules 2009, 42, 4758−4768. (26) Brandrup, J.; Immergut, E. H.; Grulke, E. A. Polymer handbook, 4th ed.; John Wiley & Sons: New York, 2003. (27) Jiang, Z.; Tang, Y.; Rieger, J.; Enderle, H. F.; Lilge, D.; Roth, S. V.; Gehrke, R.; Wu, Z.; Li, Z.; Men, Y. Polymer 2009, 50, 4101−4111. (28) Wang, Y.; Jiang, Z.; Fu, L.; Lu, Y.; Men, Y. Macromolecules 2013, 46, 7874−7879. (29) Spieckermann, F. C.; Wilhem, H. R.; Alfantis, E. C.; Zehetbauer, M. J.; Schafler, E. J. Eng. Mater. Technol. 2009, 131, 011109. (30) Zuo, F.; Keum, J. K.; Chen, X.; Hsiao, B. S.; Chen, H.; Lai, S. Y.; Wevers, R.; Li, J. Polymer 2007, 48, 6867−6880. (31) Ran, S.; Zong, X.; Fang, D.; Hsiao, B. S.; Chu, B.; Cunniff, P. M.; Phillips, R. a. J. Mater. Sci. 2001, 36, 3071−3077. (32) Ma, Z.; Shao, C.; Wang, X.; Zhao, B.; Li, X.; An, H.; Yan, T.; Li, Z.; Li, L. Polymer 2009, 50, 2706−2715. (33) Qiu, J.; Wang, Z.; Yang, L.; Zhao, J.; Niu, Y.; Hsiao, B. S. Polymer 2007, 48, 6934−6947. (34) Tang, Y.; Jiang, Z.; Men, Y.; An, L.; Enderle, H. F.; Lilge, D.; Roth, S. V.; Gehrke, R.; Rieger, J. Polymer 2007, 48, 5125−5132. (35) Garton, A.; Carlsson, D. J.; Sturgeon, P. Z.; Wiles, D. M. J. Polym. Sci., Polym. Phys. Ed. 1977, 15, 2013−2026. (36) Alamo, R. G.; Kim, M.-H.; Galante, M. J.; Isasi, J. R.; Mandelkern, L. Macromolecules 1999, 32, 4050−4064. (37) Tosaka, M.; Kohjiya, S.; Murakami, S.; Poompradub, S.; Ikeda, Y.; Toki, S.; Sics, I.; Hsiao, B. S. Rubber Chem. Technol. 2004, 77, 711− 723.

G

DOI: 10.1021/acs.macromol.5b02585 Macromolecules XXXX, XXX, XXX−XXX