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Deformation mechanism of hard elastic polyethylene film during uniaxial stretching: Effect of stretching speed
Polymer 178 (2019) 121579
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Deformation mechanism of hard elastic polyethylene film during uniaxial stretching: Effect of stretching speed
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Yuanfei Lina,b, Xueyu Lib, Xiaowei Chenb, Minfang Ana, Qianlei Zhangb, Daoliang Wangb, Wei Chenb, Panchao Yina, Lingpu Mengb,**, Liangbin Lib,* a
South China Advanced Institute for Soft Matter Science and Technology, South China University of Technology, Guangzhou, 510640, China National Synchrotron Radiation Laboratory, Anhui Provincial Engineering Laboratory of Advanced Functional Polymer Film, CAS Key Laboratory of Soft Matter Chemistry, University of Science and Technology of China, Hefei, 230026, China
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H I GH L IG H T S
of stretching speed on structures and mechanical property was studied. • Effect speed is of importance in how microphase separation occurs and develops. • Stretching deformation is inhomogeneous at low stretching speed region I. • The • Plastic deformation occurs earlier at low stretching speed.
A R T I C LE I N FO
A B S T R A C T
Keywords: Hard-elastic polyethylene films Deformation mechanism Stretching speed
The effects of stretching speed on the structural evolutions and mechanical behaviors of hard-elastic polyethylene films are studied with in-situ and ex-situ small-angle X-ray scattering (SAXS), scanning electronic microscope (SEM) and tensile tests in a wide stretching speed range (0.04–4 mm/s). Based on the evolutions of structural parameters extracted from SAXS results and the surface morphologies from SEM experiments, the stretching speed space can be divided into two regions with the boundary of 0.8 mm/s. Stress induced microphase separation of amorphous phase triggers the yielding behavior, which distributes more homogeneously with the increase of stretching speed. In region I, microphase separation tends to develop into cavities at smaller strain due to the thorough relaxation process of molecular chains in amorphous phases, which results in the inhomogeneous deformation during further stretching. In region II, the relaxation of molecular chains is not enough to response to the variation of external tensile field, thus inducing the uniform distribution of the occurrence of microphase separation.
1. Introduction Polypropylene (PP) and polyethylene (PE) microporous membranes produced by dry process serve as one of the most key materials in Lithium battery production, namely battery separator. With the quick development of electrical vehicles, high-performance Lithium battery separator in the market is required to possess great mechanical property and high porosity at the same time, which brings large challenge for industry due to the multi-step and multi-parameter properties of film processing. Generally, the production of the microporous membranes is based on the cold and hot stretching of hard-elastic precursor films, during which the relationship between structural evolutions and non-
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linear mechanical properties is still controversial. For decades, the deformation mechanisms of hard-elastic materials have attracted much attention from polymeric researchers for their perfect hyperelasticity and the importance in the post-stretching processing of microporous membranes. The microporous structures supported by fibrillar bridges are indeed originated from the coupled effects of deformation temperature, strain and strain rate. Note that most related studies were based on ex-situ experimental techniques, including tensile tests, small- and wide-angle scattering (SAXS/WAXS), scanning electronic microscope (SEM) and atomic force microscope (AFM) and so on [1–11]. With the combination of SAXS and SEM results, the hard-elastic materials (like PE or PP) have been proved to be
Corresponding author. Corresponding author. E-mail addresses: [email protected] (L. Meng), [email protected] (L. Li).
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https://doi.org/10.1016/j.polymer.2019.121579 Received 23 April 2019; Received in revised form 8 June 2019; Accepted 13 June 2019 Available online 18 June 2019 0032-3861/ © 2019 Published by Elsevier Ltd.
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composed of oriented lamellar stacks were supplied by Wuhan Bosheng Technology Company, Ltd. Based on differential scanning calorimetry measurement, the melting temperature is 134.5 °C. And the initial crystallinity is around 64.8% according to wide-angle X-ray scattering result. The thickness and initial length of the films are 25 μm and 30 mm, respectively. The other detailed structural information can be found in our previous studies [37–39].
composed of row-nucleated lamellae, i.e., highly oriented lamellae [1,12–15]. Sprague and Clark tried to use leaf spring model (energy model) combined with interlamellar separation to explain roughly the elastic deformation of crystals at room temperature [1,5,16,17], which was also partially verified by the theoretical calculation results based on a microbuckling model for idealized hard-soft (lamella-amorphous) laminated composite structures recently [13]. Besides, Goritz and coworkers proposed another hypothesis on entropic effects of interlamellar layers to account for the hard elasticity, based on the analysis of thermal effects in the amorphous part during tensile deformation [2,18]. And void or cavitation observed from SEM pictures of samples after going through further hot stretching was believed to form from interlamellar amorphous region [10,11,14,15,19,20]. Since the synergetic effects of temperature, strain and strain rate play important roles on the intrinsic deformation of oriented lamellar stacks, it remains as a large challenge to figure out what exactly happens during tensile deformation at a specific external field. Generally, molecular chains in crystal or amorphous phases possess different chain relaxation dynamics at a typical temperature, while the variation of strain rate might accelerate or decelerate the dynamics processes, to some extent, resulting in different deformation mechanisms [21–23]. Benefiting from the development of ultrafast in-situ SAXS/WAXS techniques, it has been found that at temperatures below alpha relaxation temperature (Tα) stress-induced amorphization or microphase separation of interlamellar amorphous phase might dominate the initial nonlinear mechanical behavior [24,25], while crystals slipping and melting-recrystallization phenomenon can be observed at higher deformation temperatures [25–29]. However, few studies focused on the effects of strain rate or stretching speed on the structural evolutions during tensile deformation of hard-elastic materials, which require more efforts. In fact, strain rate also serves as one of the important roles during processing of other semi-crystalline polymers, which will affect the structure and mechanical property of plastic production [30–34]. Based on the previous studies on compression-molded or injected samples, crystals intends to deform plastically before the occurrence of cavitation in amorphous phase at low strain rates, while yielding stress increases with the increase of strain rate [30,35,36]. Due to the complicated and hierarchic property from the existence of spherules, the intrinsic mechanical behavior of lamellar stacks is hard to be obtained. And although the oriented lamellar stacks cannot fully represent that of lamellar stacks in spherulites due to their different microscopic chain conformation, but they after all exclude the interference of orientation factor or larger-scale spherulite are excluded, to some extent. Thus, considering that hard-elastic materials are composed of highly oriented lamellar stacks, they can be regarded as good but probably not perfect sample for the study on the deformation mechanisms of semi-crystalline polymers. In this work, with the combination of in-situ synchrotron radiation SAXS techniques and ex-situ SEM measurement, the structural evolutions of hard-elastic PE precursor films with highly oriented lamellar stacks during uniaxial tensile deformation is systematically studied at room temperature with different stretching speeds ranging from 0.04 to 4 mm/s (over two orders of magnitude). The corresponding initial strain rate is ranging from 0.003 to 0.3 s−1. The effects of stretching speed (strain rate) on the deformation mechanisms and structural evolutions during stretching are discussed, which play an important role to understand deeply the relationship between structure and mechanical property of hard-elastic materials and might guide the processing of microporous membranes by dry process in industry.
2.2. Experimental procedures Uniaxial tensile deformation of HDPE precursor films was performed with a home-made tensile apparatus [40]. The deformation direction was along the machine direction, i.e., the normal direction of lamellar stacks. The deformation temperatures were set at room temperature and the employed stretching speeds of each clamp were ranging from 0.04 to 4 mm/s. Note that the corresponding initial strain rate was calculated as 0.003 to 0.3 s−1. Due to the limitation of the apparatus, larger initial strain rate cannot be achieved. Ten different stretching speeds were chose in the experiment. In-situ two-dimensional (2D) SAXS measurements with a wavelength of 0.103 nm were carried out during the whole stretching at the beamline BL19U2 in Shanghai Synchrotron Radiation Facility (SSRF) with the aid of Pilatus 1M detector [41]. The sample-to-detector distances for SAXS and WAXS tests were calibrated as 5937 mm and 260 mm, respectively. The time resolution of the detector was set based on the strain resolution of around 0.4%, which was different for specific stretching speed to obtain high scattering intensity as possible as we can. Fit2D was employed to analyze the SAXS data, which were corrected by subtracting the contributions of background scattering from air and sample holder. The 2D SAXS data were integrated to obtain 1D scattering profile as a function of q = 4πsinθ/λ. And long period of the lamellar structure was obtained through Bragg Law L = 2π/qmax, where qmax was the position of the scattering peak. The ex-situ SAXS experiments of typical stretched samples were performed using an in-house setup in Hefei [42], providing parallel beam with a wavelength of 0.154 nm. And a Pilatus 300K detector was used to collect 2D SAXS patterns. The images of surface morphology of microporous membranes were taken using a Field Emission Scanning Electron Microscope (FE-SEMZEISS ΣIGMA) in Institute of Nuclear Energy Safety Technology (Chinese Academy of Sciences). The voltage was 1 kV in this study to minimize the electronic damage, and samples were not sputter-coated with a gold ion beam before tests. 3. Results Fig. 1(a) shows the engineering stress (σ)-strain (ε) curves of HDPE films at room temperatures with different stretching speeds ranging from 0.04 to 4 mm/s. And the corresponding enlarged image in small strain range is provided in Fig. 1(b) to observe the characteristics more clearly. Fig. S1 in the Supporting Information presents the corresponding true stress-strain curves as a reference. With v ≤ 0.8 mm/s, the stress-strain curves behave with similar variation trend: stress σ increases linearly at the very beginning and then goes through a short plateau or weak softening, followed by a strain hardening, while further increasing stretching speed introduces a significant double yielding at large strain and the stress-strain curves seem no longer sensitive to stretching speed with 0.8 mm/s < v < 4 mm/s. Besides, Fig. 1(c) extracts the evolutions of yield strain (εy), the onset strain for strain hardening (εh) and Young modulus (E) during stretching with different stretching speed. Both εy and εh decreases with v increasing, especially with v smaller than 0.4 mm/s, while E increases almost linearly from 1.2 GPa to 2.0 GPa with the increase of v (in logarithm coordinate). With the nonlinear mechanical properties at different stretching speeds bearing in mind, we will turn to the corresponding structural evolutions in the following. Fig. 2 shows several representative two-
2. Experimental 2.1. Material and characterization The annealed high-density polyethylene (HDPE) precursor films 2
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Fig. 1. (a) The engineering stress-strain curves of hard-elastic polyethylene films at room temperature with different stretching speeds. (b) The zoomed-in image of stress-strain curves. (c) The evolutions of yield strain (εy), the onset strain for strain hardening (εh) and Young modulus (E) as a function of stretching speed v.
broad peak (at qmax ≈ 0.2 nm−1) and a shoulder peak (at qmax ≈ 0.4 nm−1) can be easily observed before stretching and maintain the ratio of 1:2 during stretching in a small strain range. Combined with the unchanged structure of lamellar structure before and after stretching (as shown in Fig. 6 in the following), these two peaks probably both come from the scattering of periodical lamellar stacks and can be regarded as the first-order peak and the second-order peak, respectively. Based on qmax, the corresponding long period can be obtained through Bragg Law. For better comparison, micro-strain of lamellar stacks is defined by the variation rate of long period. Fig. 3(a2c2) illustrate the evolutions of micro-strains (εm-1 and εm-2) during stretching, corresponding to the first- and second- order scattering, respectively. Since the first-order scattering peak might be greatly disturbed by the enlargement of the strong scattering within small q range at large strain as we observed in the SAXS patterns, it would be more reliable to use εm-2 to measure the deformation amount of microscopic lamellar stacks. The evolutions of εm-1 are provided here as references, and the zoomed-in plots of Fig. 3(a2-c2) in a small strain range are provided as Fig. S2 in the Supporting Information. Corresponding well with the movement of peaks in Fig. 3(a1), εm-2 with stretching speed of 0.04 mm/s in Fig. 3(a2) immediately increases at first and the increment is suddenly accelerated after εy. At the very beginning of strain hardening, εm-2 even obtains a value of around 32%, which is significantly larger than the employed macro-strain of around 20%. Note that with strain increasing into strain hardening zone, it is hard to determine the exact value of εm-1. And εm-2 turns to slow down and keep
dimensional (2D) SAXS patterns collected in-situ during tensile deformation with three typical stretching speeds (0.04 mm/s, 0.8 mm/s and 4 mm/s). The initial pattern before stretching holds two-point scattering along meridian, which means the initial structure is composed of oriented lamellar stacks as mentioned in our previous studies [12,13,25,39]. Taking the situation of 0.04 mm/s as the example, the two-point scattering gradually moves close to beamstop combined with the appearance of diamond-shaped scattering along meridian near yield point. Due to the strengthening of scattering intensity, the second-order scattering can be also observed by naked eyes. Further stretching will greatly increase the scattering intensity and at the same time new streak signal in equator shows up, of which the intensity is gradually increased. Note that with strain larger than 125% at the later stage of strain hardening zone in stress-strain curve, the two-point signal is distorted obviously and the second-order scattering disappears, probably indicating that lamellar crystals are no longer arranged along machine direction and massive slipping of lamellar crystals or crystal broken might be motivated. The evolutions of SAXS patterns at stretching speeds of 0.8 and 4 mm/s are basically similar to the results observed at v = 0.04 mm/s, while the distorted two-point scattering is never observed even at strain of 250%. In order to distinguish the difference of structural evolutions among the three different stretching speeds, we will further extract quantitative structural parameters. Fig. 3(a1-c1) present the one-dimensional integration curves during stretching at the three typical stretching speeds, which are calculated within meridian area as indicated by the inset image of Fig. 3(a1). One
Fig. 2. The representative in-situ 2D SAXS patterns during stretching with three typical stretching speeds: 0.04 mm/s, 0.8 mm/s and 4 mm/s. The numbers in the first row denote strain with unit of percentage. 3
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Fig. 3. The 1D SAXS integration profiles calculated within meridian region (1), the evolutions of calculated micro-strain (2) and 1D SAXS integration profiles calculated within equatorial region (3) during stretching at three typical stretching speeds: (a1-a3) 0.04 mm/s, (b1-b3) 0.8 mm/s and (c1-c3) 4 mm/s, respectively.
previous studies on hard-elastic polypropylene films [12,25], suggesting that stress induced density fluctuation or phase separation in interlamellar amorphous phase might occur around the yielding point. Furthermore, the intensities of scattering signal close to beamstop (Ilow) and the second-order scattering signal (Im-2) are extracted in Fig. 4. The values are corrected with the consideration of the variation of sample's thickness. Here Ilow is probably originated from the occurrence of low-density regions between lamellar stacks, while Im-2 is related to the arranging and concentration of periodical lamellae. Note that for low strain rates like 0.04 mm/s, the second-order peak disappeared at large strain and the Im-2 is just roughly estimate the intensity of the corresponding area. Based on the mask protocols of integrating areas as shown in Fig. 4(a), Fig. 4(b) and (c) present the corresponding evolutions of Ilow and Im-2 during stretching with different stretching speeds. As presented in Fig. 4(b), the increase of Ilow is motivated immediately at the beginning of stretching and suddenly accelerated once entering non-linear zones, followed by a decreasing trend after reaching a maximum with stretching speeds smaller than 0.8 mm/s. And further increasing stretching speed introduces a new plateau of Ilow before reaching the maximum. The strain of the maximum is increased with the increase of stretching speeds, while the maximal intensity is gradually decreasing except for 0.04 mm/s. In Fig. 4(c), Im-2 follows a two-step increasing trend to reach a maximal intensity and then decreases until stretching is stopped. Both the intensities of the maximal value of Im-2 and the corresponding strain
a constant value of around 35% during further stretching, which also disappears with strain larger than 160% probably due to the coupled effects of the bad periodical arrangement of lamellar stacks and the interference of the strong scattering near beamstop. With v increasing to 0.8 and 4 mm/s, the deformation seems more uniform so that we can calculate εm-2 during the whole stretching. As strain is in the linear elastic region and the beginning of strain hardening region, εm-2 follows a similar non-linear increasing trend as that at 0.04 mm/s and increases to around 100% before the occurrence of double yielding in stressstrain curves. With further stretching, εm-2 turns to follow a two-stage decreasing trend until tensile deformation is stopped, during which the turning point of decreasing slope might be consistent with the end of double yielding behavior. Note that during stretching with the three different stretching speeds, micro-strains always perform a self-propelled increasing trend and hold values larger than macro-strains once passing through yielding region, which might indicate the occurrence of microphase separation within amorphous phase between lamellae. In order to further check whether microphase separation occurs or not during stretching, we obtain the corresponding 1D integration profiles calculated within equatorial regions during stretching. As shown in Fig. 3(a3-c3), there exists a vague and wide peak, of which the intensity is increased significantly during stretching. Due to the interference of strong scattering at small q range, we can only roughly estimate the average length of the periodical domains as around 10–20 nm with qmax ranging from 0.3 to 0.5 nm−1. This phenomenon is similar to our 4
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Fig. 6. SEM images of stretched hard-elastic polyethylene films at different conditions: (a) 0.04 mm/s, 50%; (b) 0.04 mm/s, 125%; (c) 4 mm/s, 50%; (d) 4 mm/s, 125%. Fig. 4. (a) Mask protocol in 2D SAXS patterns to define Ilow and Im-2. (b-d) The evolutions of several structural parameters during stretching at different stretching speeds: (b) Ilow, (c) Im-2, and (d) Φ.
SAXS patterns. Correspondingly, in Fig. 5(b) with strain increasing, the sample gradually transforms from transparent to whitening state with v = 0.04 mm/s, while samples preserve specific transparency at strains of 25% and 50% but become a little white at strain of 125% with v = 4 mm/s. To further check the microscopic deformation effect, Fig. 6 presents the SEM surface morphologies of stretched samples mentioned above. As the two-point scattering in meridian presented in Fig. 6, most lamella are basically arranged along stretching direction despite of the strain rates and stretching strains. With v = 0.04 mm/s, several slender crazes perpendicular to the stretching direction can be observed (Fig. 6(a)) with strain of 50%. Further increasing strain to 125%, lamellae are further separated and more cavities exist with the support of fibrillar bridges (Fig. 6(b)). Differently, with v = 4 mm/s no cavities exist with strain of 50% (Fig. 6(c)), but slender crazes perpendicular to the stretching direction is observed and distributed homogeneously when the sample is stretched to 125% (Fig. 6(d)). Considering that the inhomogeneity of microscopic deformation also plays an important role on the macroscopic mechanical property, we further compare the evolutions of elastic recovery rate (ER50) and the ratio of Ilow and Im-2 (Φ50) at preset strain of 50% as presented in Fig. 7(a). With stretching speed increasing (in log scale), ER50 follows a linear increasing trend, while Φ50 decreases almost linearly at the same time. It is suggested that the hard elasticity indeed has positive correlation with the homogeneity of deformation. Fig. 7(b) summarizes the evolutions of ER at different pre-set strains during stretching with the three typical stretching speeds (0.04, 0.8 and 4 mm/s). ER obviously increases with stretching speed increasing even when pre-set strain increasing to 250%, while ER all perform a nonlinear decreasing trend
perform an increasing trend. Furthermore, we plot the evolutions of the ratio of Ilow and Im-2 (Φ) in Fig. 4(d), which can be used to roughly estimate the nonuniformity and uniformity of deformation, respectively. Φ starts to increase non-linearly as the stress-strain curves enter non-linear zones and the increasing slope decreases significantly with the increase of stretching speed especially with v smaller than 0.8 mm/ s. These phenomena might indicate that the deformation becomes more and more homogeneous with the increase of stretching speed. Fig. 5(a) and (b) present the typical 2D SAXS patterns and the photos of stretched hard-elastic polyethylene films, respectively. Note that these samples are prepared through cyclic tensile deformation with typical pre-set strains (25%, 50% and 125%) and stretching speeds (0.04 and 4 mm/s). At small strains like ε = 25% and 50%, rhombic scattering along stretching direction near beamstop and two-point scattering coexist with stretching speed of 0.04 mm/s, while only twopoint scattering is observed with stretching speed of 4 mm/s. At large strain of 125%, the intensity in meridian near beamstop is greatly strengthened combined with the appearance of streak signal in equator with v = 0.04 mm/s, suggesting the existence of massive cavities and fibrillar bridges, while with v = 4 mm/s the two-point scattering combined with the rhombic scattering in meridian is observed in the 2D
Fig. 7. (a) The evolutions of elastic recovery rate (ER50) and the ratio of Ilow and Im-2 at strain of 50% as functions of stretching speeds. (b) The evolutions of elastic recovery rates (ER) with increase of pre-set strain at three typical stretching speeds.
Fig. 5. (a) Typical 2D SAXS patterns and (b) the photos of stretched hard-elastic polyethylene films. The stretching direction is vertical. 5
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Fig. 8. We will first discuss the situation at high stretching speed region II (4 mm/s) since the deformation is more homogeneous and simpler according to the relative SAXS and SEM results. Based on the hyperelastic mechanical behavior and recoverable structural parameters, elastic deformation is proposed to take responsibility for the deformation mechanism of hard-elastic polyethylene films at room temperature, especially for the initial nonlinear mechanical behavior. As shown in Fig. 8(b), lamellae are separated through the elastic extension of amorphous chains between lamellar layers in the linear elastic zone before yielding. Considering that the increase of microscopic strain εm-2 is accelerated (Fig. 3(c2)) combined with the appearance of the wide peak in Ie-q curve (Fig. 3(c3)) in the following yielding and short stress plateau zone, it can be proposed that microphase separation within amorphous phase in lamellar stacks triggers the initial nonlinear mechanical behavior (see Fig. 8(c)), which is further developed in strain hardening region, resulting in the further separation of lamellae. Based on the recoverable property of SAXS patterns and film photo, no massive and irreversible crystal broken occurs at strain smaller than 50%. Further stretching induces a strain softening behavior, during which εm2 turns to decrease and the slim streak signal in equatorial gradually forms. With strain increasing to 125% (strain hardening zone), irrecoverable plastic deformation might dominate the main deformation since the strong scattering near beamstop in the SAXS pattern and the stress whitening phenomenon of sample after cyclic stretching are observed in Fig. 6(b). Partial lamellar crystals are broken and new fibrillar bridges are formed as shown in Fig. 8(d), thus resulting in the massive formation of cavities. As for samples stretched with low stretching speed (region I like 0.04 m/s), the microscopic structural evolutions also go through the three steps as discussed above: lamellar separation, microphase separation in amorphous and the formation of fibrillar bridges together with cavitation. Differently, the deformation is not so homogeneous, and the irreversible deformation occurs earlier and more massively if compared with what we observed at high stretching speed. As the accelerated increase of εm-2 will be found together with the appearance of the diffusive and wide peak in Ie-q curve once stress-strain curve enters the nonlinear mechanical zone (yielding and stress plateau zones), microphase separation of amorphous phase is triggered at this condition. Then with the development of microphase separation in the following strain hardening zone, the increase of εm-2 turns to slow down and the scattering from large-scale low-density region and lamellae are more and more difficult to separate. Note that the strong diamondshaped scattering near beamstop in SAXS pattern still exists after sample is recovered mechanically with preset strain of 25% (early stage of strain hardening), while stress whitening phenomenon is more distinct and the film is even no longer transparent with preset strain of 125% (after secondary yielding). Therefore, we propose that microphase separation within amorphous phase has gradually developed into
with pre-set strain increasing. Taking the situation of 4 mm/s as an example, ER decreases almost linearly from 100% to around 50% during stretching, and the decrease of ER seems to slow down with strain larger than 100%. 4. Discussion With the combination of in-situ and ex-situ SAXS experiments and SEM measurements, some interesting findings can be extracted. (i) Based on the evolutions of structural parameters (Lm-1, Lm-2, Ie and Φ) and mechanical data (σ and ER), the stretching speed range can be divided into two regions with 0.8 mm/s as the approximate boundary: low stretching speed region I (0.04–0.8 mm/s) and high stretching speed region II (0.8–4 mm/s). (ii) In region I, the stress-strain curves only go through four typical stages, i.e., linear elastic zone, yield and stress plateau zone, strain hardening zone and weak second-yielding zone. The corresponding deformation of microscopic structures is obviously inhomogeneous as presented by the relatively diffusive scattering signal in meridian and near beamstop, high value of the ratio of the intensity in low-density region and the second-order scattering (Φ), and the non-uniformly distributed cavities with different size. (iii) In high stretching speed region II, a new strain softening behavior appears at large strain and seems more and more important with v increasing from 1.6 to 4 mm/s. Meanwhile, lamellar separation, microphase separation within amorphous area and even the distribution of cavity are more homogeneous in this region. (iv) Based on the inverse variation relationship between ER and Φ50, the elastic recovery property has positive correlation with the deformation homogeneity. And films stretched at high stretching speed can obtain higher elastic recovery rate, which decreases with the increase of strain. Before discussion, it is necessary to know that the initial hard-elastic polyethylene films are composed of highly oriented lamellar stacks based on the surface morphology and SAXS results. At room temperature, lamellar crystals possess lowest molecular mobility and the modulus is probably two or three orders of magnitude larger than that of amorphous phases. That is, the deformation ability of amorphous chains is probably more sensitive to the strength of external stress field. Therefore, we will focus on the effects of stretching speed on the structural evolutions during stretching of hard-elastic polyethylene films, especially the deformation in amorphous phases in the following discussion section. 4.1. Structural evolutions in regions I and II It would be the first step to figure out the structural evolutions of hard-elastic polyethylene films during the whole tensile deformation with a certain stretching speed. In order to make the description easier to be followed, we provide the schematic image of the deformation mechanism of oriented polyethylene lamellar stack during stretching in
Fig. 8. The schematic image of the deformation mechanism of oriented polyethylene lamellar stack. 6
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strain, which results in the inhomogeneous deformation during further stretching. With stretching speed increasing into region II, the relaxation of molecular chains cannot catch up with the variation of external tensile field and microphase separation distributed more uniformly and gradually. Therefore, the deformation tends to be homogeneous, which provide better hard elasticity. This study will deepen the understanding of the deformation mechanism of oriented lamellar stacks in specific external tensile field, which can also guide the post-stretching process in the production of microporous membranes by dry process.
cavitation at the initial strain hardening zone. And the formation of fibrillar bridges together with partial lamellar broken occurs at the same time. Further stretching might induces massive slipping of deformation unity like lamellar stacks plotted in Fig. 8, resulting in the distorted scattering signal in SAXS pattern at strain of 250%. 4.2. Coupled effects of chain mobility and external field at different stretching speeds The different structural evolutions at different stretching speeds are probably originated from the coupling effects of molecular mobility and motivated external tensile field. As we discussed in our previous studies about hard-elastic polypropylene films, the triggering of microphase separation within amorphous phase occurs gradually and lamellar stacks with and without microphase separation coexist especially in the initial nonlinear mechanical zone. In low stretching speed region I (0.04–0.8 mm/s), molecular chains in amorphous phase have enough time to relax during stretching. Amorphous chains within lamellar stacks without microphase separation might be relaxed thoroughly, while in lamellar stacks with microphase separation, amorphous phases with high and low density would be separated more distinctly due to the relaxation effects and self-propelled microphase separation process. Therefore, the stress concentration phenomenon becomes more obvious in large scale, which make it much easier for microphase separation to develop into cavitation. Here microphase separation within amorphous phase denotes two amorphous phases with high and low density, respectively, while cavitation is the extreme case, where low density decreases to 0 (vacuum). If we accelerate the stretching speeds into region II (0.8–4 mm/s), two possibilities should be taken into consideration and the distribution of stress concentration will be different. (i) The whole degree of microphase separation at small strains is relatively low since amorphous chains do not have enough time to relax thoroughly, which can be supported by the smaller value of εm-2 (observed from enlarged plot of Fig. 3). (ii) Based on the chain relaxation state, the degree of microphase separation in those lamellar stacks at the early nonlinear stage is low so that they consume less stress and energy. Then the transmitting of stress or work done by the external force in other lamellar stacks nearby without microphase separation would be seldomly affected. This stress situation can introduce more lamellar stacks to join in microphase separation behavior within a shorter time or strain window. That is, the distribution of stress from microscopic point of view is more homogeneous even at large strains (125%). Therefore, the distributions of lamellae and cavitation are more uniform and the hard-elastic mechanical property performs more perfectly in high stretching speed region II.
Acknowledgements This work is supported by the National Natural Science Foundation of China (51633009, 51790503), the China Postdoctoral Science Foundation (2018M643076) and the National Key R&D Program of China (2018YFB0704200). And the experiment is carried out in Shanghai Synchrotron Radiation Facility beamline BL19U2. Appendix A. Supplementary data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.polymer.2019.121579. References [1] B. Sprague, Relationship of structure and morphology to properties of “hard” elastic fibers and films, J. Macromol. Sci., Part B: Physics 8 (1–2) (1973) 157–187. [2] S. Cannon, G. McKenna, W. Statton, Hard‐Elastic fibers.(A review of a novel state for crystalline polymers), J. Polym. Sci. Macromol. Rev. 11 (1) (1976) 209–275. [3] S.-Y. Lee, S.-Y. Park, H.-S. Song, Lamellar crystalline structure of hard elastic HDPE films and its influence on microporous membrane formation, Polymer 47 (10) (2006) 3540–3547. [4] R. Hosemann, I. Schulze, Refraction effects and structural changes of hard elastic polypropylene (HEPP) during stretching, Colloid Polym. Sci. 265 (8) (1987) 686–695. [5] E. Clark, A mechanism of energy-driven elasticity in crystalline polymers, in: R. Lenz, R. Stein (Eds.), Structure and Properties of Polymer Films, vol. 1, Springer US, 1973, pp. 267–282. [6] H.D. Noether, Factors affecting the formation of hard elastic fibres, Polym. Eng. Sci. 19 (6) (1978) 427–432. [7] H.D. Noether, I.L. Hay, Small-angle X-ray diffraction studies and morphology of microporous materials and their 'hard' elastic precursors, J. Appl. Crystallogr. 11 (5) (1978) 546–547. [8] I. Park, H. Noether, Crystalline “hard” elastic materials, Colloid Polym. Sci. 253 (10) (1975) 824–839. [9] W. Ren, Hard elastic polypropylene-nature, internal friction, and surface energy, Colloid Polym. Sci. 270 (10) (1992) 943–955. [10] J. Xie, R. Xu, C. Lei, Uniaxial stretching induced pore nucleation and growth in rownucleated crystalline hard-elastic polypropylene film: the effect of activation volume and stretching work, Polymer 158 (2018) 10–17. [11] C. Lei, R. Xu, Z. Tian, H. Huang, J. Xie, X. Zhu, Stretching-induced uniform micropores formation: an in situ SAXS/WAXS study, Macromolecules 51 (9) (2018) 3433–3442. [12] Y. Lin, X. Li, L. Meng, X. Chen, F. Lv, Q. Zhang, L. Li, Stress-induced microphase separation of interlamellar amorphous phase in hard-elastic isotactic polypropylene film, Polymer 148 (2018) 79–92. [13] Y. Lin, F. Tian, L. Meng, X. Chen, F. Lv, Q. Zhang, L. Li, Microbuckling: A possible mechanism to trigger nonlinear instability of semicrystalline polymer, Polymer 154 (2018) 48–54. [14] H. Noether, W. Whitney, X-ray diffraction and morphology of crystalline, hard, elastic materials. Aktuelle Probleme der Polymer-Physik IV, Springer, 1973, pp. 991–1005. [15] K. Matsui, N. Hosaka, K. Suzuki, Y. Shinohara, Y. Amemiya, Microscopic deformation behavior of hard elastic polypropylene during cold-stretching process in fabrication of microporous membrane as revealed by synchrotron X-ray scattering, Polymer 70 (0) (2015) 215–221. [16] E.S. Clark, C.A. Garber, Effect of industrial processing on the morphology of crystalline polymers, Int.J.Polym. Mater.Polym.Biomater. 1 (1) (1971) 31–45. [17] I. Kuryndin, V. Lavrentyev, V. Bukošek, G. Elyashevich, Percolation transitions in porous polyethylene and polypropylene films with lamellar structures, Polym. Sci. 57 (6) (2015) 717–722. [18] D. Gäritz, F. Müller, Rückstellungsmechanismus von elastischen Hartfasern, Colloid Polym. Sci. 252 (10) (1974) 862–870. [19] C.J. Chou, A. Hiltner, E. Baer, The role of surface stresses in the deformation of hard elastic polypropylene, Polymer 27 (3) (1986) 369–376. [20] I.A. Okkelman, A.A. Dolgova, S. Banerjee, J.P. Kerry, A.L. Volynskii, O.V. Arzhakova, D.B. Papkovsky, Phosphorescent oxygen and mechano-sensitive nanostructured materials based on hard elastic polypropylene films, ACS Appl.
5. Conclusion The effects of stretching speed on the structural evolutions of oriented polyethylene lamellar stacks are systematically studied with hard-elastic polyethylene films via SAXS, SEM and tensile tests with a wide stretching speed range (0.04–4 mm/s). Based on the evolutions of microscopic strain εm-2, equatorial intensity Ie, the ratio of the intensities of low-density and second-order scattering Φ, elastic recovery rate ER and other parameters, the stretching speed space can be divided into two regions: low stretching speed region I (0.04–0.8 mm/s) and high stretching speed region II (0.8–4 mm/s). Whether stretched with stretching speeds located in regions I or II, the deformation of microscopic structures will go through three stages: lamellar separation, microphase separation within amorphous phase and the formations of fibrillar bridges and cavities. But stretching speed exactly plays an important role in how microphase separation occurs and develops into cavities. With a low stretching speed in region I, the stress distribution is not uniform since thorough chain relaxation make it easier to differentiate lamellar stacks with and without microphase separation. In this case, microphase separation tends to develop into cavities at small 7
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Deformation mechanism of hard elastic polyethylene film during uniaxial stretching: Effect of stretching speed
T
Yuanfei Lina,b, Xueyu Lib, Xiaowei Chenb, Minfang Ana, Qianlei Zhangb, Daoliang Wangb, Wei Chenb, Panchao Yina, Lingpu Mengb,**, Liangbin Lib,* a
South China Advanced Institute for Soft Matter Science and Technology, South China University of Technology, Guangzhou, 510640, China National Synchrotron Radiation Laboratory, Anhui Provincial Engineering Laboratory of Advanced Functional Polymer Film, CAS Key Laboratory of Soft Matter Chemistry, University of Science and Technology of China, Hefei, 230026, China
b
H I GH L IG H T S
of stretching speed on structures and mechanical property was studied. • Effect speed is of importance in how microphase separation occurs and develops. • Stretching deformation is inhomogeneous at low stretching speed region I. • The • Plastic deformation occurs earlier at low stretching speed.
A R T I C LE I N FO
A B S T R A C T
Keywords: Hard-elastic polyethylene films Deformation mechanism Stretching speed
The effects of stretching speed on the structural evolutions and mechanical behaviors of hard-elastic polyethylene films are studied with in-situ and ex-situ small-angle X-ray scattering (SAXS), scanning electronic microscope (SEM) and tensile tests in a wide stretching speed range (0.04–4 mm/s). Based on the evolutions of structural parameters extracted from SAXS results and the surface morphologies from SEM experiments, the stretching speed space can be divided into two regions with the boundary of 0.8 mm/s. Stress induced microphase separation of amorphous phase triggers the yielding behavior, which distributes more homogeneously with the increase of stretching speed. In region I, microphase separation tends to develop into cavities at smaller strain due to the thorough relaxation process of molecular chains in amorphous phases, which results in the inhomogeneous deformation during further stretching. In region II, the relaxation of molecular chains is not enough to response to the variation of external tensile field, thus inducing the uniform distribution of the occurrence of microphase separation.
1. Introduction Polypropylene (PP) and polyethylene (PE) microporous membranes produced by dry process serve as one of the most key materials in Lithium battery production, namely battery separator. With the quick development of electrical vehicles, high-performance Lithium battery separator in the market is required to possess great mechanical property and high porosity at the same time, which brings large challenge for industry due to the multi-step and multi-parameter properties of film processing. Generally, the production of the microporous membranes is based on the cold and hot stretching of hard-elastic precursor films, during which the relationship between structural evolutions and non-
*
linear mechanical properties is still controversial. For decades, the deformation mechanisms of hard-elastic materials have attracted much attention from polymeric researchers for their perfect hyperelasticity and the importance in the post-stretching processing of microporous membranes. The microporous structures supported by fibrillar bridges are indeed originated from the coupled effects of deformation temperature, strain and strain rate. Note that most related studies were based on ex-situ experimental techniques, including tensile tests, small- and wide-angle scattering (SAXS/WAXS), scanning electronic microscope (SEM) and atomic force microscope (AFM) and so on [1–11]. With the combination of SAXS and SEM results, the hard-elastic materials (like PE or PP) have been proved to be
Corresponding author. Corresponding author. E-mail addresses: [email protected] (L. Meng), [email protected] (L. Li).
**
https://doi.org/10.1016/j.polymer.2019.121579 Received 23 April 2019; Received in revised form 8 June 2019; Accepted 13 June 2019 Available online 18 June 2019 0032-3861/ © 2019 Published by Elsevier Ltd.
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composed of oriented lamellar stacks were supplied by Wuhan Bosheng Technology Company, Ltd. Based on differential scanning calorimetry measurement, the melting temperature is 134.5 °C. And the initial crystallinity is around 64.8% according to wide-angle X-ray scattering result. The thickness and initial length of the films are 25 μm and 30 mm, respectively. The other detailed structural information can be found in our previous studies [37–39].
composed of row-nucleated lamellae, i.e., highly oriented lamellae [1,12–15]. Sprague and Clark tried to use leaf spring model (energy model) combined with interlamellar separation to explain roughly the elastic deformation of crystals at room temperature [1,5,16,17], which was also partially verified by the theoretical calculation results based on a microbuckling model for idealized hard-soft (lamella-amorphous) laminated composite structures recently [13]. Besides, Goritz and coworkers proposed another hypothesis on entropic effects of interlamellar layers to account for the hard elasticity, based on the analysis of thermal effects in the amorphous part during tensile deformation [2,18]. And void or cavitation observed from SEM pictures of samples after going through further hot stretching was believed to form from interlamellar amorphous region [10,11,14,15,19,20]. Since the synergetic effects of temperature, strain and strain rate play important roles on the intrinsic deformation of oriented lamellar stacks, it remains as a large challenge to figure out what exactly happens during tensile deformation at a specific external field. Generally, molecular chains in crystal or amorphous phases possess different chain relaxation dynamics at a typical temperature, while the variation of strain rate might accelerate or decelerate the dynamics processes, to some extent, resulting in different deformation mechanisms [21–23]. Benefiting from the development of ultrafast in-situ SAXS/WAXS techniques, it has been found that at temperatures below alpha relaxation temperature (Tα) stress-induced amorphization or microphase separation of interlamellar amorphous phase might dominate the initial nonlinear mechanical behavior [24,25], while crystals slipping and melting-recrystallization phenomenon can be observed at higher deformation temperatures [25–29]. However, few studies focused on the effects of strain rate or stretching speed on the structural evolutions during tensile deformation of hard-elastic materials, which require more efforts. In fact, strain rate also serves as one of the important roles during processing of other semi-crystalline polymers, which will affect the structure and mechanical property of plastic production [30–34]. Based on the previous studies on compression-molded or injected samples, crystals intends to deform plastically before the occurrence of cavitation in amorphous phase at low strain rates, while yielding stress increases with the increase of strain rate [30,35,36]. Due to the complicated and hierarchic property from the existence of spherules, the intrinsic mechanical behavior of lamellar stacks is hard to be obtained. And although the oriented lamellar stacks cannot fully represent that of lamellar stacks in spherulites due to their different microscopic chain conformation, but they after all exclude the interference of orientation factor or larger-scale spherulite are excluded, to some extent. Thus, considering that hard-elastic materials are composed of highly oriented lamellar stacks, they can be regarded as good but probably not perfect sample for the study on the deformation mechanisms of semi-crystalline polymers. In this work, with the combination of in-situ synchrotron radiation SAXS techniques and ex-situ SEM measurement, the structural evolutions of hard-elastic PE precursor films with highly oriented lamellar stacks during uniaxial tensile deformation is systematically studied at room temperature with different stretching speeds ranging from 0.04 to 4 mm/s (over two orders of magnitude). The corresponding initial strain rate is ranging from 0.003 to 0.3 s−1. The effects of stretching speed (strain rate) on the deformation mechanisms and structural evolutions during stretching are discussed, which play an important role to understand deeply the relationship between structure and mechanical property of hard-elastic materials and might guide the processing of microporous membranes by dry process in industry.
2.2. Experimental procedures Uniaxial tensile deformation of HDPE precursor films was performed with a home-made tensile apparatus [40]. The deformation direction was along the machine direction, i.e., the normal direction of lamellar stacks. The deformation temperatures were set at room temperature and the employed stretching speeds of each clamp were ranging from 0.04 to 4 mm/s. Note that the corresponding initial strain rate was calculated as 0.003 to 0.3 s−1. Due to the limitation of the apparatus, larger initial strain rate cannot be achieved. Ten different stretching speeds were chose in the experiment. In-situ two-dimensional (2D) SAXS measurements with a wavelength of 0.103 nm were carried out during the whole stretching at the beamline BL19U2 in Shanghai Synchrotron Radiation Facility (SSRF) with the aid of Pilatus 1M detector [41]. The sample-to-detector distances for SAXS and WAXS tests were calibrated as 5937 mm and 260 mm, respectively. The time resolution of the detector was set based on the strain resolution of around 0.4%, which was different for specific stretching speed to obtain high scattering intensity as possible as we can. Fit2D was employed to analyze the SAXS data, which were corrected by subtracting the contributions of background scattering from air and sample holder. The 2D SAXS data were integrated to obtain 1D scattering profile as a function of q = 4πsinθ/λ. And long period of the lamellar structure was obtained through Bragg Law L = 2π/qmax, where qmax was the position of the scattering peak. The ex-situ SAXS experiments of typical stretched samples were performed using an in-house setup in Hefei [42], providing parallel beam with a wavelength of 0.154 nm. And a Pilatus 300K detector was used to collect 2D SAXS patterns. The images of surface morphology of microporous membranes were taken using a Field Emission Scanning Electron Microscope (FE-SEMZEISS ΣIGMA) in Institute of Nuclear Energy Safety Technology (Chinese Academy of Sciences). The voltage was 1 kV in this study to minimize the electronic damage, and samples were not sputter-coated with a gold ion beam before tests. 3. Results Fig. 1(a) shows the engineering stress (σ)-strain (ε) curves of HDPE films at room temperatures with different stretching speeds ranging from 0.04 to 4 mm/s. And the corresponding enlarged image in small strain range is provided in Fig. 1(b) to observe the characteristics more clearly. Fig. S1 in the Supporting Information presents the corresponding true stress-strain curves as a reference. With v ≤ 0.8 mm/s, the stress-strain curves behave with similar variation trend: stress σ increases linearly at the very beginning and then goes through a short plateau or weak softening, followed by a strain hardening, while further increasing stretching speed introduces a significant double yielding at large strain and the stress-strain curves seem no longer sensitive to stretching speed with 0.8 mm/s < v < 4 mm/s. Besides, Fig. 1(c) extracts the evolutions of yield strain (εy), the onset strain for strain hardening (εh) and Young modulus (E) during stretching with different stretching speed. Both εy and εh decreases with v increasing, especially with v smaller than 0.4 mm/s, while E increases almost linearly from 1.2 GPa to 2.0 GPa with the increase of v (in logarithm coordinate). With the nonlinear mechanical properties at different stretching speeds bearing in mind, we will turn to the corresponding structural evolutions in the following. Fig. 2 shows several representative two-
2. Experimental 2.1. Material and characterization The annealed high-density polyethylene (HDPE) precursor films 2
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Fig. 1. (a) The engineering stress-strain curves of hard-elastic polyethylene films at room temperature with different stretching speeds. (b) The zoomed-in image of stress-strain curves. (c) The evolutions of yield strain (εy), the onset strain for strain hardening (εh) and Young modulus (E) as a function of stretching speed v.
broad peak (at qmax ≈ 0.2 nm−1) and a shoulder peak (at qmax ≈ 0.4 nm−1) can be easily observed before stretching and maintain the ratio of 1:2 during stretching in a small strain range. Combined with the unchanged structure of lamellar structure before and after stretching (as shown in Fig. 6 in the following), these two peaks probably both come from the scattering of periodical lamellar stacks and can be regarded as the first-order peak and the second-order peak, respectively. Based on qmax, the corresponding long period can be obtained through Bragg Law. For better comparison, micro-strain of lamellar stacks is defined by the variation rate of long period. Fig. 3(a2c2) illustrate the evolutions of micro-strains (εm-1 and εm-2) during stretching, corresponding to the first- and second- order scattering, respectively. Since the first-order scattering peak might be greatly disturbed by the enlargement of the strong scattering within small q range at large strain as we observed in the SAXS patterns, it would be more reliable to use εm-2 to measure the deformation amount of microscopic lamellar stacks. The evolutions of εm-1 are provided here as references, and the zoomed-in plots of Fig. 3(a2-c2) in a small strain range are provided as Fig. S2 in the Supporting Information. Corresponding well with the movement of peaks in Fig. 3(a1), εm-2 with stretching speed of 0.04 mm/s in Fig. 3(a2) immediately increases at first and the increment is suddenly accelerated after εy. At the very beginning of strain hardening, εm-2 even obtains a value of around 32%, which is significantly larger than the employed macro-strain of around 20%. Note that with strain increasing into strain hardening zone, it is hard to determine the exact value of εm-1. And εm-2 turns to slow down and keep
dimensional (2D) SAXS patterns collected in-situ during tensile deformation with three typical stretching speeds (0.04 mm/s, 0.8 mm/s and 4 mm/s). The initial pattern before stretching holds two-point scattering along meridian, which means the initial structure is composed of oriented lamellar stacks as mentioned in our previous studies [12,13,25,39]. Taking the situation of 0.04 mm/s as the example, the two-point scattering gradually moves close to beamstop combined with the appearance of diamond-shaped scattering along meridian near yield point. Due to the strengthening of scattering intensity, the second-order scattering can be also observed by naked eyes. Further stretching will greatly increase the scattering intensity and at the same time new streak signal in equator shows up, of which the intensity is gradually increased. Note that with strain larger than 125% at the later stage of strain hardening zone in stress-strain curve, the two-point signal is distorted obviously and the second-order scattering disappears, probably indicating that lamellar crystals are no longer arranged along machine direction and massive slipping of lamellar crystals or crystal broken might be motivated. The evolutions of SAXS patterns at stretching speeds of 0.8 and 4 mm/s are basically similar to the results observed at v = 0.04 mm/s, while the distorted two-point scattering is never observed even at strain of 250%. In order to distinguish the difference of structural evolutions among the three different stretching speeds, we will further extract quantitative structural parameters. Fig. 3(a1-c1) present the one-dimensional integration curves during stretching at the three typical stretching speeds, which are calculated within meridian area as indicated by the inset image of Fig. 3(a1). One
Fig. 2. The representative in-situ 2D SAXS patterns during stretching with three typical stretching speeds: 0.04 mm/s, 0.8 mm/s and 4 mm/s. The numbers in the first row denote strain with unit of percentage. 3
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Fig. 3. The 1D SAXS integration profiles calculated within meridian region (1), the evolutions of calculated micro-strain (2) and 1D SAXS integration profiles calculated within equatorial region (3) during stretching at three typical stretching speeds: (a1-a3) 0.04 mm/s, (b1-b3) 0.8 mm/s and (c1-c3) 4 mm/s, respectively.
previous studies on hard-elastic polypropylene films [12,25], suggesting that stress induced density fluctuation or phase separation in interlamellar amorphous phase might occur around the yielding point. Furthermore, the intensities of scattering signal close to beamstop (Ilow) and the second-order scattering signal (Im-2) are extracted in Fig. 4. The values are corrected with the consideration of the variation of sample's thickness. Here Ilow is probably originated from the occurrence of low-density regions between lamellar stacks, while Im-2 is related to the arranging and concentration of periodical lamellae. Note that for low strain rates like 0.04 mm/s, the second-order peak disappeared at large strain and the Im-2 is just roughly estimate the intensity of the corresponding area. Based on the mask protocols of integrating areas as shown in Fig. 4(a), Fig. 4(b) and (c) present the corresponding evolutions of Ilow and Im-2 during stretching with different stretching speeds. As presented in Fig. 4(b), the increase of Ilow is motivated immediately at the beginning of stretching and suddenly accelerated once entering non-linear zones, followed by a decreasing trend after reaching a maximum with stretching speeds smaller than 0.8 mm/s. And further increasing stretching speed introduces a new plateau of Ilow before reaching the maximum. The strain of the maximum is increased with the increase of stretching speeds, while the maximal intensity is gradually decreasing except for 0.04 mm/s. In Fig. 4(c), Im-2 follows a two-step increasing trend to reach a maximal intensity and then decreases until stretching is stopped. Both the intensities of the maximal value of Im-2 and the corresponding strain
a constant value of around 35% during further stretching, which also disappears with strain larger than 160% probably due to the coupled effects of the bad periodical arrangement of lamellar stacks and the interference of the strong scattering near beamstop. With v increasing to 0.8 and 4 mm/s, the deformation seems more uniform so that we can calculate εm-2 during the whole stretching. As strain is in the linear elastic region and the beginning of strain hardening region, εm-2 follows a similar non-linear increasing trend as that at 0.04 mm/s and increases to around 100% before the occurrence of double yielding in stressstrain curves. With further stretching, εm-2 turns to follow a two-stage decreasing trend until tensile deformation is stopped, during which the turning point of decreasing slope might be consistent with the end of double yielding behavior. Note that during stretching with the three different stretching speeds, micro-strains always perform a self-propelled increasing trend and hold values larger than macro-strains once passing through yielding region, which might indicate the occurrence of microphase separation within amorphous phase between lamellae. In order to further check whether microphase separation occurs or not during stretching, we obtain the corresponding 1D integration profiles calculated within equatorial regions during stretching. As shown in Fig. 3(a3-c3), there exists a vague and wide peak, of which the intensity is increased significantly during stretching. Due to the interference of strong scattering at small q range, we can only roughly estimate the average length of the periodical domains as around 10–20 nm with qmax ranging from 0.3 to 0.5 nm−1. This phenomenon is similar to our 4
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Fig. 6. SEM images of stretched hard-elastic polyethylene films at different conditions: (a) 0.04 mm/s, 50%; (b) 0.04 mm/s, 125%; (c) 4 mm/s, 50%; (d) 4 mm/s, 125%. Fig. 4. (a) Mask protocol in 2D SAXS patterns to define Ilow and Im-2. (b-d) The evolutions of several structural parameters during stretching at different stretching speeds: (b) Ilow, (c) Im-2, and (d) Φ.
SAXS patterns. Correspondingly, in Fig. 5(b) with strain increasing, the sample gradually transforms from transparent to whitening state with v = 0.04 mm/s, while samples preserve specific transparency at strains of 25% and 50% but become a little white at strain of 125% with v = 4 mm/s. To further check the microscopic deformation effect, Fig. 6 presents the SEM surface morphologies of stretched samples mentioned above. As the two-point scattering in meridian presented in Fig. 6, most lamella are basically arranged along stretching direction despite of the strain rates and stretching strains. With v = 0.04 mm/s, several slender crazes perpendicular to the stretching direction can be observed (Fig. 6(a)) with strain of 50%. Further increasing strain to 125%, lamellae are further separated and more cavities exist with the support of fibrillar bridges (Fig. 6(b)). Differently, with v = 4 mm/s no cavities exist with strain of 50% (Fig. 6(c)), but slender crazes perpendicular to the stretching direction is observed and distributed homogeneously when the sample is stretched to 125% (Fig. 6(d)). Considering that the inhomogeneity of microscopic deformation also plays an important role on the macroscopic mechanical property, we further compare the evolutions of elastic recovery rate (ER50) and the ratio of Ilow and Im-2 (Φ50) at preset strain of 50% as presented in Fig. 7(a). With stretching speed increasing (in log scale), ER50 follows a linear increasing trend, while Φ50 decreases almost linearly at the same time. It is suggested that the hard elasticity indeed has positive correlation with the homogeneity of deformation. Fig. 7(b) summarizes the evolutions of ER at different pre-set strains during stretching with the three typical stretching speeds (0.04, 0.8 and 4 mm/s). ER obviously increases with stretching speed increasing even when pre-set strain increasing to 250%, while ER all perform a nonlinear decreasing trend
perform an increasing trend. Furthermore, we plot the evolutions of the ratio of Ilow and Im-2 (Φ) in Fig. 4(d), which can be used to roughly estimate the nonuniformity and uniformity of deformation, respectively. Φ starts to increase non-linearly as the stress-strain curves enter non-linear zones and the increasing slope decreases significantly with the increase of stretching speed especially with v smaller than 0.8 mm/ s. These phenomena might indicate that the deformation becomes more and more homogeneous with the increase of stretching speed. Fig. 5(a) and (b) present the typical 2D SAXS patterns and the photos of stretched hard-elastic polyethylene films, respectively. Note that these samples are prepared through cyclic tensile deformation with typical pre-set strains (25%, 50% and 125%) and stretching speeds (0.04 and 4 mm/s). At small strains like ε = 25% and 50%, rhombic scattering along stretching direction near beamstop and two-point scattering coexist with stretching speed of 0.04 mm/s, while only twopoint scattering is observed with stretching speed of 4 mm/s. At large strain of 125%, the intensity in meridian near beamstop is greatly strengthened combined with the appearance of streak signal in equator with v = 0.04 mm/s, suggesting the existence of massive cavities and fibrillar bridges, while with v = 4 mm/s the two-point scattering combined with the rhombic scattering in meridian is observed in the 2D
Fig. 7. (a) The evolutions of elastic recovery rate (ER50) and the ratio of Ilow and Im-2 at strain of 50% as functions of stretching speeds. (b) The evolutions of elastic recovery rates (ER) with increase of pre-set strain at three typical stretching speeds.
Fig. 5. (a) Typical 2D SAXS patterns and (b) the photos of stretched hard-elastic polyethylene films. The stretching direction is vertical. 5
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Fig. 8. We will first discuss the situation at high stretching speed region II (4 mm/s) since the deformation is more homogeneous and simpler according to the relative SAXS and SEM results. Based on the hyperelastic mechanical behavior and recoverable structural parameters, elastic deformation is proposed to take responsibility for the deformation mechanism of hard-elastic polyethylene films at room temperature, especially for the initial nonlinear mechanical behavior. As shown in Fig. 8(b), lamellae are separated through the elastic extension of amorphous chains between lamellar layers in the linear elastic zone before yielding. Considering that the increase of microscopic strain εm-2 is accelerated (Fig. 3(c2)) combined with the appearance of the wide peak in Ie-q curve (Fig. 3(c3)) in the following yielding and short stress plateau zone, it can be proposed that microphase separation within amorphous phase in lamellar stacks triggers the initial nonlinear mechanical behavior (see Fig. 8(c)), which is further developed in strain hardening region, resulting in the further separation of lamellae. Based on the recoverable property of SAXS patterns and film photo, no massive and irreversible crystal broken occurs at strain smaller than 50%. Further stretching induces a strain softening behavior, during which εm2 turns to decrease and the slim streak signal in equatorial gradually forms. With strain increasing to 125% (strain hardening zone), irrecoverable plastic deformation might dominate the main deformation since the strong scattering near beamstop in the SAXS pattern and the stress whitening phenomenon of sample after cyclic stretching are observed in Fig. 6(b). Partial lamellar crystals are broken and new fibrillar bridges are formed as shown in Fig. 8(d), thus resulting in the massive formation of cavities. As for samples stretched with low stretching speed (region I like 0.04 m/s), the microscopic structural evolutions also go through the three steps as discussed above: lamellar separation, microphase separation in amorphous and the formation of fibrillar bridges together with cavitation. Differently, the deformation is not so homogeneous, and the irreversible deformation occurs earlier and more massively if compared with what we observed at high stretching speed. As the accelerated increase of εm-2 will be found together with the appearance of the diffusive and wide peak in Ie-q curve once stress-strain curve enters the nonlinear mechanical zone (yielding and stress plateau zones), microphase separation of amorphous phase is triggered at this condition. Then with the development of microphase separation in the following strain hardening zone, the increase of εm-2 turns to slow down and the scattering from large-scale low-density region and lamellae are more and more difficult to separate. Note that the strong diamondshaped scattering near beamstop in SAXS pattern still exists after sample is recovered mechanically with preset strain of 25% (early stage of strain hardening), while stress whitening phenomenon is more distinct and the film is even no longer transparent with preset strain of 125% (after secondary yielding). Therefore, we propose that microphase separation within amorphous phase has gradually developed into
with pre-set strain increasing. Taking the situation of 4 mm/s as an example, ER decreases almost linearly from 100% to around 50% during stretching, and the decrease of ER seems to slow down with strain larger than 100%. 4. Discussion With the combination of in-situ and ex-situ SAXS experiments and SEM measurements, some interesting findings can be extracted. (i) Based on the evolutions of structural parameters (Lm-1, Lm-2, Ie and Φ) and mechanical data (σ and ER), the stretching speed range can be divided into two regions with 0.8 mm/s as the approximate boundary: low stretching speed region I (0.04–0.8 mm/s) and high stretching speed region II (0.8–4 mm/s). (ii) In region I, the stress-strain curves only go through four typical stages, i.e., linear elastic zone, yield and stress plateau zone, strain hardening zone and weak second-yielding zone. The corresponding deformation of microscopic structures is obviously inhomogeneous as presented by the relatively diffusive scattering signal in meridian and near beamstop, high value of the ratio of the intensity in low-density region and the second-order scattering (Φ), and the non-uniformly distributed cavities with different size. (iii) In high stretching speed region II, a new strain softening behavior appears at large strain and seems more and more important with v increasing from 1.6 to 4 mm/s. Meanwhile, lamellar separation, microphase separation within amorphous area and even the distribution of cavity are more homogeneous in this region. (iv) Based on the inverse variation relationship between ER and Φ50, the elastic recovery property has positive correlation with the deformation homogeneity. And films stretched at high stretching speed can obtain higher elastic recovery rate, which decreases with the increase of strain. Before discussion, it is necessary to know that the initial hard-elastic polyethylene films are composed of highly oriented lamellar stacks based on the surface morphology and SAXS results. At room temperature, lamellar crystals possess lowest molecular mobility and the modulus is probably two or three orders of magnitude larger than that of amorphous phases. That is, the deformation ability of amorphous chains is probably more sensitive to the strength of external stress field. Therefore, we will focus on the effects of stretching speed on the structural evolutions during stretching of hard-elastic polyethylene films, especially the deformation in amorphous phases in the following discussion section. 4.1. Structural evolutions in regions I and II It would be the first step to figure out the structural evolutions of hard-elastic polyethylene films during the whole tensile deformation with a certain stretching speed. In order to make the description easier to be followed, we provide the schematic image of the deformation mechanism of oriented polyethylene lamellar stack during stretching in
Fig. 8. The schematic image of the deformation mechanism of oriented polyethylene lamellar stack. 6
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strain, which results in the inhomogeneous deformation during further stretching. With stretching speed increasing into region II, the relaxation of molecular chains cannot catch up with the variation of external tensile field and microphase separation distributed more uniformly and gradually. Therefore, the deformation tends to be homogeneous, which provide better hard elasticity. This study will deepen the understanding of the deformation mechanism of oriented lamellar stacks in specific external tensile field, which can also guide the post-stretching process in the production of microporous membranes by dry process.
cavitation at the initial strain hardening zone. And the formation of fibrillar bridges together with partial lamellar broken occurs at the same time. Further stretching might induces massive slipping of deformation unity like lamellar stacks plotted in Fig. 8, resulting in the distorted scattering signal in SAXS pattern at strain of 250%. 4.2. Coupled effects of chain mobility and external field at different stretching speeds The different structural evolutions at different stretching speeds are probably originated from the coupling effects of molecular mobility and motivated external tensile field. As we discussed in our previous studies about hard-elastic polypropylene films, the triggering of microphase separation within amorphous phase occurs gradually and lamellar stacks with and without microphase separation coexist especially in the initial nonlinear mechanical zone. In low stretching speed region I (0.04–0.8 mm/s), molecular chains in amorphous phase have enough time to relax during stretching. Amorphous chains within lamellar stacks without microphase separation might be relaxed thoroughly, while in lamellar stacks with microphase separation, amorphous phases with high and low density would be separated more distinctly due to the relaxation effects and self-propelled microphase separation process. Therefore, the stress concentration phenomenon becomes more obvious in large scale, which make it much easier for microphase separation to develop into cavitation. Here microphase separation within amorphous phase denotes two amorphous phases with high and low density, respectively, while cavitation is the extreme case, where low density decreases to 0 (vacuum). If we accelerate the stretching speeds into region II (0.8–4 mm/s), two possibilities should be taken into consideration and the distribution of stress concentration will be different. (i) The whole degree of microphase separation at small strains is relatively low since amorphous chains do not have enough time to relax thoroughly, which can be supported by the smaller value of εm-2 (observed from enlarged plot of Fig. 3). (ii) Based on the chain relaxation state, the degree of microphase separation in those lamellar stacks at the early nonlinear stage is low so that they consume less stress and energy. Then the transmitting of stress or work done by the external force in other lamellar stacks nearby without microphase separation would be seldomly affected. This stress situation can introduce more lamellar stacks to join in microphase separation behavior within a shorter time or strain window. That is, the distribution of stress from microscopic point of view is more homogeneous even at large strains (125%). Therefore, the distributions of lamellae and cavitation are more uniform and the hard-elastic mechanical property performs more perfectly in high stretching speed region II.
Acknowledgements This work is supported by the National Natural Science Foundation of China (51633009, 51790503), the China Postdoctoral Science Foundation (2018M643076) and the National Key R&D Program of China (2018YFB0704200). And the experiment is carried out in Shanghai Synchrotron Radiation Facility beamline BL19U2. Appendix A. Supplementary data Supplementary data to this article can be found online at https:// doi.org/10.1016/j.polymer.2019.121579. References [1] B. Sprague, Relationship of structure and morphology to properties of “hard” elastic fibers and films, J. Macromol. Sci., Part B: Physics 8 (1–2) (1973) 157–187. [2] S. Cannon, G. McKenna, W. Statton, Hard‐Elastic fibers.(A review of a novel state for crystalline polymers), J. Polym. Sci. Macromol. Rev. 11 (1) (1976) 209–275. [3] S.-Y. Lee, S.-Y. Park, H.-S. Song, Lamellar crystalline structure of hard elastic HDPE films and its influence on microporous membrane formation, Polymer 47 (10) (2006) 3540–3547. [4] R. Hosemann, I. Schulze, Refraction effects and structural changes of hard elastic polypropylene (HEPP) during stretching, Colloid Polym. Sci. 265 (8) (1987) 686–695. [5] E. Clark, A mechanism of energy-driven elasticity in crystalline polymers, in: R. Lenz, R. Stein (Eds.), Structure and Properties of Polymer Films, vol. 1, Springer US, 1973, pp. 267–282. [6] H.D. Noether, Factors affecting the formation of hard elastic fibres, Polym. Eng. Sci. 19 (6) (1978) 427–432. [7] H.D. Noether, I.L. Hay, Small-angle X-ray diffraction studies and morphology of microporous materials and their 'hard' elastic precursors, J. Appl. Crystallogr. 11 (5) (1978) 546–547. [8] I. Park, H. Noether, Crystalline “hard” elastic materials, Colloid Polym. Sci. 253 (10) (1975) 824–839. [9] W. Ren, Hard elastic polypropylene-nature, internal friction, and surface energy, Colloid Polym. Sci. 270 (10) (1992) 943–955. [10] J. Xie, R. Xu, C. Lei, Uniaxial stretching induced pore nucleation and growth in rownucleated crystalline hard-elastic polypropylene film: the effect of activation volume and stretching work, Polymer 158 (2018) 10–17. [11] C. Lei, R. Xu, Z. Tian, H. Huang, J. Xie, X. Zhu, Stretching-induced uniform micropores formation: an in situ SAXS/WAXS study, Macromolecules 51 (9) (2018) 3433–3442. [12] Y. Lin, X. Li, L. Meng, X. Chen, F. Lv, Q. Zhang, L. Li, Stress-induced microphase separation of interlamellar amorphous phase in hard-elastic isotactic polypropylene film, Polymer 148 (2018) 79–92. [13] Y. Lin, F. Tian, L. Meng, X. Chen, F. Lv, Q. Zhang, L. Li, Microbuckling: A possible mechanism to trigger nonlinear instability of semicrystalline polymer, Polymer 154 (2018) 48–54. [14] H. Noether, W. Whitney, X-ray diffraction and morphology of crystalline, hard, elastic materials. Aktuelle Probleme der Polymer-Physik IV, Springer, 1973, pp. 991–1005. [15] K. Matsui, N. Hosaka, K. Suzuki, Y. Shinohara, Y. Amemiya, Microscopic deformation behavior of hard elastic polypropylene during cold-stretching process in fabrication of microporous membrane as revealed by synchrotron X-ray scattering, Polymer 70 (0) (2015) 215–221. [16] E.S. Clark, C.A. Garber, Effect of industrial processing on the morphology of crystalline polymers, Int.J.Polym. Mater.Polym.Biomater. 1 (1) (1971) 31–45. [17] I. Kuryndin, V. Lavrentyev, V. Bukošek, G. Elyashevich, Percolation transitions in porous polyethylene and polypropylene films with lamellar structures, Polym. Sci. 57 (6) (2015) 717–722. [18] D. Gäritz, F. Müller, Rückstellungsmechanismus von elastischen Hartfasern, Colloid Polym. Sci. 252 (10) (1974) 862–870. [19] C.J. Chou, A. Hiltner, E. Baer, The role of surface stresses in the deformation of hard elastic polypropylene, Polymer 27 (3) (1986) 369–376. [20] I.A. Okkelman, A.A. Dolgova, S. Banerjee, J.P. Kerry, A.L. Volynskii, O.V. Arzhakova, D.B. Papkovsky, Phosphorescent oxygen and mechano-sensitive nanostructured materials based on hard elastic polypropylene films, ACS Appl.
5. Conclusion The effects of stretching speed on the structural evolutions of oriented polyethylene lamellar stacks are systematically studied with hard-elastic polyethylene films via SAXS, SEM and tensile tests with a wide stretching speed range (0.04–4 mm/s). Based on the evolutions of microscopic strain εm-2, equatorial intensity Ie, the ratio of the intensities of low-density and second-order scattering Φ, elastic recovery rate ER and other parameters, the stretching speed space can be divided into two regions: low stretching speed region I (0.04–0.8 mm/s) and high stretching speed region II (0.8–4 mm/s). Whether stretched with stretching speeds located in regions I or II, the deformation of microscopic structures will go through three stages: lamellar separation, microphase separation within amorphous phase and the formations of fibrillar bridges and cavities. But stretching speed exactly plays an important role in how microphase separation occurs and develops into cavities. With a low stretching speed in region I, the stress distribution is not uniform since thorough chain relaxation make it easier to differentiate lamellar stacks with and without microphase separation. In this case, microphase separation tends to develop into cavities at small 7
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