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. 2019 Dec 27;5(1):655–666. doi: 10.1021/acsomega.9b03250

Structure Evolution and Deformation Behavior of Polyethylene Film during Biaxial Stretching

Qiang Chen 1, Zhiwei Wang 1, Sumei Zhang 1, Ya Cao 1, Jinyao Chen 1,*
PMCID: PMC6964308  PMID: 31956815

Abstract

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The structure evolution and deformation behavior of tenter-frame biaxially oriented polyethylene (TF-BOPE) films were investigated in this study. For sequential biaxial stretching, the original spherulites were broken into small pieces, and the fibrillar structure simultaneously formed during the machine direction (MD) stretching. Subsequently, the single fibrils were pulled away, and some fine fibrils developed via partial melting–recrystallization or lamellar rearrangement during the transverse direction (TD) stretching, which leads to the formation of the vein structure and nanosized fiber-like network. For simultaneous biaxial stretching, the fractured lamellae and the newly formed crystals were evenly distributed in the MD-TD plane like an isotropous fibrillar network. Moreover, compared with the unstretched sample, both films could achieve up to about 2 times the tensile modulus and 4.5 times the tensile strength and also exhibited the superior optical property.

1. Introduction

Tenter-frame biaxial stretching is an innovative technology for polyethylene (PE) films, where the cast sheet is drawn in two vertical directions with a single procedure (simultaneous biaxial stretching) or two subsequent procedures (sequential biaxial stretching).1 Compared with the traditional blown film, the tenter-frame biaxially oriented polyethylene (TF-BOPE) film can achieve up to twice the tensile modulus, thrice the tensile strength, twice the impact strength, thrice the puncture, and 80% less haze (www.dow.com). Moreover, the TF-BOPE film has great toughness and flex cracking resistance even under low temperature, so it is especially suitable for the frozen food packaging.2−4

The biaxially oriented films have been widely investigated, and the structure evolution can be classified in two types in terms of the crystallization capacity. On the one hand, some polymers are easy to crystallize with the cast sheet in a semi-crystalline state, such as polypropylene (PP), which is biaxially drawn at a temperature near the melting point (Tm). The spherulitic morphology of the PP film turns into the stacked lamellae via melting in the machine direction (MD, the first tensile direction in sequential biaxial stretching) and then forms a fibrillar network through the crystallographic slip in the transverse direction (TD, the second tensile direction in sequential biaxial stretching, perpendicular to MD).5,6 Moreover, a visualized microstructure of the nanometer-scale fibrillar network for the BOPP film is observed by atomic force microscopy (AFM), and the morphology can be controlled by selecting an appropriate biaxial draw ratio.7

On the other hand, other polymers, such as poly(ethylene terephthalate) (PET), polylactic acid (PLA), and polystyrene (PS), can be quenched into a quasi-amorphous state through a chill roll, which are biaxially drawn in a rubbery state (above the Tg). The model of the structure evolution during sequential biaxial stretching is proposed, and it is determined by the development of the crystal structure in the MD stretching.8,9 Simultaneously, the second population of oriented but poorly ordered crystals develops along TD.10 The relaxation behavior of the biaxially stretched PET films is divided into three regimes according to the relationship between birefringence and strain, which is directly in contact with the development of strain-induced crystallization at different biaxial ratios.11−13

There are many investigations regarding the uniaxial deformation of PE,14−22 whereas research about structure evolution for the TF-BOPE film is scanty.23−25 Indeed, biaxial stretching is a very important processing method for the fabrication of PE films in industry, but the microstructure evolution during TD stretching is still unclear.

In this work, we studied the processes of sequential biaxial stretching in detail and also studied the processes of simultaneous biaxial stretching as a comparison. The results show that the structure evolution for the two processes is conspicuously different and the simultaneously biaxially stretched film exhibited a more uniform lamellae distribution in the MD-TD plane than the sequentially biaxially stretched film. Moreover, after biaxial stretching, both films show superior mechanical and optical performance. In addition, the investigation may also provide a guide for choosing the optimal manufacture window and controlling the microstructure in industry.

2. Experimental Section

2.1. Materials and Sample Preparation

The linear low-density polyethylene used in this study was kindly offered by Guangdong Decro Film New Materials Co., Ltd., with a melt flow rate (MFR) of 1.8 g/10 min (190 °C, 2.16 kg). The number-average molecular weight (Mn) and weight-average molecular weight (Mw) were 24,900 and 116,500 g/mol, respectively. The resin pellets were molded into the 1 mm-thick sheet (190 °C, 15 MPa) and then quenched into water at room temperature.

2.2. Film Stretching

The TF-BOPE film was prepared on a biaxial stretcher (KARO IV, Brückner, Germany). The square sample (90 mm × 90 mm) was cut from the compression sheet. After being preheated in the first chamber (114 °C) for 2 min, the samples for sequential biaxial stretching and simultaneous biaxial stretching were employed with the stretching rate of 0.25 s–1 at 114 °C. Then, the film was moved to the second chamber (100 °C), setting for 2 min. As shown in Figure S1a, the film of sequential biaxial stretching was first stretched in MD with the draw ratio of 6 and constrained at the same time in TD. Subsequently, the different draw ratios of 1, 2, 3, 4, 5, and 6 in TD were chosen, which were labeled PE-6×1, PE-6×2, PE-6×3, PE-6×4, PE-6×5, and PE-6×6, respectively. Moreover, the simultaneously biaxially stretched film with the draw ratio of 6 × 6 and the original sheet were called PE-sim-6×6 and PE-0, respectively.

2.3. Scanning Electron Microscopy (SEM)

To get visual crystal morphology, the sample was chemically etched by a potassium permanganate solution.26,27 The morphology was gained by using field-emission scanning electron microscopy (SEM, S4700, Hitachi, Japan).

2.4. Atomic Force Microscopy (AFM)

The surface morphologies of the film were investigated on the Asylum Research atomic force microscope (Cypher VRS, Oxford, England) in the tapping mode. The needle type of AC240TS was used in this study with the resonance frequency of 45–95 kHz and the spring constant of 0.3–4.8 N/m. The image of a 3 μm square was gained at the scan rate of 2.44 Hz.

2.5. Differential Scanning Calorimetry (DSC)

Thermal analysis was carried out with the DSC 3+ STARe System (Mettler Toledo, Switzerland). Heat flow and temperature scales were calibrated using the high-purity indium and zinc. A nominal 5 mg of sample was heated from 25 to 170 °C at the heating rate of 10 °C/min under a nitrogen atmosphere (50 mL/min). Furthermore, the crystallinity (XDSC) is computed according to the following equation

2.5. 1

where ΔHm is the melting enthalpy, and ΔHm0 is the melting enthalpy of 100% crystalline PE (287.3 J/g).28 Moreover, the distribution of short-chain branching (SCB) for the ethylene/α-olefin copolymer was tested by the method of successive self-nucleation and annealing (SSA) fractionation, and the thermal procedures are shown in Figure S2a.29,30

2.6. Fourier Transform Infrared Spectroscopy (FTIR)

The biaxial orientation of the PE film was measured by the FTIR spectrometer (Nicolet 560, Thermo Fisher Scientific, USA) with the resolution of 2 cm–1 and accumulation of 64 scans in transmission mode.31 The polarization of the beam was performed using a zinc selenide wire grid polarizer. The sample was placed perpendicular to the FTIR beam with MD in the vertical direction and TD in the horizontal direction. Then, the measurements were performed with the polarization beam in the positions of 0° and 90°, respectively.24

The White–Spruiell biaxial orientation factors (fseg, MDB), which can quantify the orientation of interesting segments in MD and TD, are defined according to

2.6. 2
2.6. 3

where “seg” is the a, b, or c axis of the orthorhombic crystal structure for PE. In addition, the θ1 and θ2 are defined in terms of the angles between the crystallographic axis and MD/TD of the biaxially stretched film. The values of f are limited, with f = 1 for perfect orientation, f = 0 for random orientation, and f = −1 for complete perpendicular orientation. The details for this method can be found elsewhere.32,33

2.7. Wide-Angle X-ray Diffraction (WAXD)

The measurement of crystal structure and orientation was carried out using a wide-angle X-ray diffractometer (WAXD, D8 Discover, Bruker, Germany) with Cu Kα X-ray radiation (λ = 0.154 nm).34 The films were stacked to ∼0.5 mm, and their two-dimensional WAXD (2D-WAXD) patterns were gained by placing the sample surface perpendicular to the projection beams with the exposure time for 3 min. One-dimensional WAXD (1D-WAXD) data could be obtained by integrating 2D-WAXD patterns as a function of 2θ. The crystallinity (XXRD) is defined as follows

2.7. 4

where Aa and Ac are the fitted areas of amorphous and crystalline regions, respectively. Moreover, the crystallite size (L) of the various planes for PE was calculated from Debye–Scherrer’s equation

2.7. 5

where λ is the wavelength of X-ray, β is the full width of the diffraction line at half maximum, and θ is the Bragg angle.

2.8. Small-Angle X-ray Scattering (SAXS)

The SAXS experiment was conducted on the Xeuss 2.0 system (Xenocs, France) equipped with a multilayer focused Cu Kα X-ray source (λ = 0.154 nm). The sample-to-detector distance was fixed at 2500 mm, which provided the effective scattering vector q (q = (4π sin θ)/λ)) range from 0.05 to 1.15 nm–1. The long period of the lamellae (Lp) was determined from Bragg’s law35

2.8. 6

where qmax is the maximum scattering vector of the relevant intensity profile. The lamellar thicknesses (Lc) and amorphous thicknesses (La) are computed in terms of the following equations

2.8. 7
2.8. 8
2.8. 9

where ρc (1.003 g/cm3), ρa (0.855 g/cm3), and ρ are the densities of the PE crystalline phase, PE amorphous phase, and tested sample, respectively.36 Moreover, it should be noted that the equations assume a much larger lateral extent of the crystalline lamellae as compared with thickness.

2.9. Ultraviolet–Visible Spectroscopy (UV–Vis)

The luminous transmittance (T) of the PE film (∼30 μm in thickness) was measured by a UV–Vis spectrometer (UV 1800, Shimadzu, Japan) with the wavelength range from 400 to 800 nm.

2.10. Mechanical Testing

The biaxially stretched film was tailored into the rectangles (60 mm × 10 mm) along MD, TD, and diagonal direction (DD, angled 45° from MD), respectively, as shown in Figure S1b. The above samples were conducted on a universal tensile testing machine (Instron, USA) with the crosshead speed of 20 mm/min at 23 °C, and each experiment was repeated five times.

3. Results and Discussion

3.1. Biaxial Stress–Strain Behavior

The stress–strain curves of sequential and simultaneous biaxial stretching are shown in Figure 1a,b, respectively. In the first step of sequential biaxial stretching, the sheet was drawn along MD to an objective draw ratio of 6, where both curves of MD and TD exhibited a yielding behavior due to localized necking. Moreover, the tensile strength in MD was about twice larger than that in TD. Strain hardening appeared at a high draw ratio for the MD curve, which meant that some fibrillar structure developed during the first stretching. In the second step of sequential biaxial stretching, a continual decrease in stress in the MD curve was observed, which was because the original crystals along MD gradually turned toward TD. In the simultaneous biaxial stretching, the typical stress–strain behavior (yielding, strain softening, and strain hardening) of both directions was observed, which suggested that the original spherulites in the sheet were destroyed by biaxial load and then transformed into fibrils oriented along MD or TD.

Figure 1.

Figure 1

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Stress–strain curves of (a) sequential biaxial stretching30 and (b) simultaneous biaxial stretching.

3.2. Morphology Evolution during Biaxial Stretching

The morphology evolution of PE films during biaxial stretching is investigated by AFM, as shown in Figure 2. For the sample of PE-0, the isotropous lamellar structure was found in the phase image. After the MD stretching, there are plenty of crystals oriented along MD in the sample of PE-6×1. To characterize more detailed morphology for the samples of PE-0 and PE-6×1, SEM images are also obtained after selective etching shown in Figure S3. An isotropous spherulite structure and random lamellae were observed in the sheet of PE-0. Sequentially, lamellae were broken into small pieces and were well arranged along MD. Moreover, the above lamellar fragments were incorporated into the fibrillar structure, which looked like a string of pearls. It was worth mentioning that the SEM images of other samples were not obtained because the films (25 to 50 μm in thickness) were too thin to be etched.

Figure 2.

Figure 2

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AFM height and phase images (3 μm × 3 μm) of the sequentially biaxially stretched films prepared at various transverse draw ratios and the PE-sim-6×6 film. The vertical and horizontal arrows indicate MD and TD, respectively.

During the TD stretching, a separation of compact and well-arranged fibrils took place in the sample of PE-6×2, where some single fibrils were pulled away from fibrillar bundles and were slightly inclined against MD. The above morphology with the larger fibrillar bundles along MD and the divided individual fibrils against MD looked like a leaf vein and was called a vein structure in the following. From the topographic image of PE-6×3, a newly emerged structure of the nanosized fiber-like network was also observed. Also, more single fibrils were split from the vein structure, and some fine fibrils, developed by the partial melting–recrystallization or lamellar rearrangement, were highly aligned parallel to TD. With the increase in λTD from 4 to 6, the diameter of the fibrillar bundle became substantially smaller, and then the divided single fibrils rotated to TD by the transverse loading forces. Although some vein structures were not susceptible to the transverse stretching and were still parallel to MD, the majority of fibrils gradually oriented toward TD. We could conclude that the general fibrillar distribution progressively reoriented from MD to TD under the combined effects of crystal rotation and melting–recrystallization. On the other hand, the morphology of PE-sim-6×6 had some significant differences compared to that of PE-6×6, which displayed a uniform and nanosized network structure without the vein structure along MD. Moreover, the orientation of fibrils in the sample of PE-sim-6×6 was isotropous in the film plane, not like the fibrillar distribution of PE-6×6 either in MD or TD.

3.3. Thermal Analyses

The thermograms of sequentially and simultaneously biaxially stretched films are exhibited in Figure 3a. We could find that all the melting curves showed multiple-melting behavior after drawing. For the sample of PE-0, there was the central peak at 127.9 °C, the shoulder peak at 123.5 °C, and the shoulder peak at 104.7 °C, which suggested that three kinds of lamellae with different thickness values coexisted in the PE sheet in terms of the Thomson–Gibbs equation.37 Moreover, the formation of the shoulder peak at 123.5 °C was due to the quenching process. After stretching, the shoulder peak at 104.7 °C began to fade. As shown in Figure S2b, there are some peaks at low temperature, implying that the SCB distribution is heterogeneous.

Figure 3.

Figure 3

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(a) Melting curves, (b) corresponding crystallinity (XDSC), and (c) Xright of the sequentially biaxially stretched films prepared at various transverse draw ratios and the PE-sim-6×6 film.

The corresponding crystallinity of different samples is exhibited in Figure 3b, where the XDSC gradually increases until the biaxial draw ratio of 6 × 4 and then decreases for the biaxial draw ratios of 6 × 5 and 6 × 6. The change in XDSC might ascribe to the coupling effects of the recrystallization for amorphous chains and the destruction for original lamellae. Moreover, the crystallinity of PE-sim-6×6 was slightly larger than that of PE-6×6. To investigate the proportion of thick lamellae (127.9 °C), we separated all melting endotherms into the left and right areas along the dotted line at 125 °C and then calculated the crystallinity of the right peak at 127.7 °C (Xright) shown in Figure 3c. The Xright gradually reduced with the increase in λTD, suggesting that the thick lamellae were continuously destroyed by the transverse stress. Therefore, we could conclude that, during the TD stretching, the destruction of thick lamellae (127.9 °C) and the recrystallization occurred at the same time. Moreover, before the biaxial ratio of 6 × 4, the rate of recrystallization was higher than that of destruction, leading to the increase in crystallinity. When most of the amorphous molecules regularly folded into crystals, the process of crystalline elimination dominated during stretching and the crystallinity decreased at the biaxial ratios of 6 × 5 and 6 × 6.

3.4. Microstructure and Crystallinity

As displayed in Figure 4a, the (110) and (200) peaks of the 1D-WAXD profiles gradually move to a smaller scattering angle during the TD stretching. Moreover, the crystalline structure parameters (2θ), interplanar spacing (d-spacing), and crystallite size (L) are computed and listed in Table S1. The d-spacing increased and L decreased with the increase in λTD, indicating the extension of the crystal plane parallel to the chain axis and the reduction of the lamella in the direction perpendicular to the (110) and (200) planes. Besides, more lamellae gradually oriented to TD, which might develop due to effects of crystal rotation and recrystallization. As shown in Figure 4b, XXRD rises sharply at first before λTD reaches 4 and then declines for λTD = 5 and 6, which is similar to the trend of the crystallinity change from DSC.

Figure 4.

Figure 4

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(a) 1D-WAXD profiles and (b) corresponding crystallinity of the sequentially biaxially stretched films prepared at various transverse draw ratios and the PE-sim-6×6 film.

3.5. Orientation Structure

The successive 2D-WAXD patterns of different samples are shown in Figure 5a, where the equator and meridian are defined as the horizontal direction (MD) and vertical direction (TD), respectively. The diffraction rings of the (110) and (200) lattice planes were observed in the sample of PE-0 and then transformed into two pairs of concentrated and bright diffractions in the horizontal direction after the first step of stretching, implying the formation of the fibrillar structure and well-stacked lamellae along MD.38 Subsequently, the diffraction signal of PE-6×1 transformed from the meridian to the equator during the second step of stretching.

Figure 5.

Figure 5

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(a) 2D-WAXD patterns and (b) corresponding azimuthal intensity curves of the (110) lattice plane of the sequentially biaxially stretched films prepared at various transverse draw ratios and the PE-sim-6×6 film. The horizontal and vertical arrows indicate MD and TD, respectively.

The corresponding azimuthal intensity curves of the (110) lattice plane are also shown in Figure 5b, where the azimuthal angles of 0° (180°) and 90° (270°) are defined as the meridian and equator directions, respectively. For the sample of PE-6×1, two sharp peaks in the azimuthal intensity curve were found at 0° and 180°, and the crystals were highly oriented along MD. During the TD stretching, two weak peaks appeared in the azimuthal angles of 90° and 270°, and meanwhile, the intense peaks of 0° and 180° were still located at the meridian direction for the samples of PE-6×2 and PE-6×3, which indicated that another population of crystals was parallel to TD. When the sheet was stretched to the biaxial draw ratio of 6 × 4, the multiple peaks of 0° (180°) and 90° (270°) simultaneously appeared, and their intensity was almost same, indicating that the crystal orientation toward MD and TD reached a state of equilibrium. As the λTD increased to 5 and 6, the main peaks were situated at 90° and 270° and the general crystal orientation shifted from MD to TD. In addition, azimuthal intensity curves of PE-sim-6×6 showed no distinct peaks because of isotropic crystal distribution.

To describe the crystal orientation in more detail, White–Spruiell biaxial orientation characterized by FTIR was also used in this study. The triangular diagram and the values of White-Spruiell biaxial orientation factors for different samples are shown in Figure 6a and Table 1, respectively. It is worth mentioning that all the orientation points (fMDB and fTD) lie within the isosceles triangle, with the origin (0, 0) representing isotropic orientation, point (0.5, 0.5) representing equal planar orientation in the MD-TD plane, and point (−1, −1) representing the perpendicular orientation to the film surface shown in Figure S4.31

Figure 6.

Figure 6

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(a) White–Spruiell orientation triangle diagram of the sequentially biaxially stretched films prepared at various transverse draw ratios and the PE-sim-6×6 film. Arabic numerals (1 to 6) represent the transverse draw ratio of the sequentially biaxially stretched films. (b) Schematics of the assumed crystal orientation.

Table 1. White–Spruiell Biaxial Orientation Factors of Different Samples.

sample fc, MDB fc, TDB fa, MDB fa, TDB fb, MDB fb, TDB
PE-0 0 0 0 0 0 0
PE-6×1 0.94 0.03 –0.46 –0.04 –0.39 –0.08
PE-6×2 0.46 0.41 –0.38 –0.31 –0.12 –0.06
PE-6×3 0.44 0.42 –0.34 –0.38 –0.10 –0.04
PE-6×4 0.41 0.45 –0.35 –0.43 –0.05 –0.03
PE-6×5 0.32 0.46 –0.36 –0.46 0.03 0.01
PE-6×6 0.27 0.47 –0.36 –0.49 0.07 0.04
PE-sim-6×6 0.43 0.42 –0.46 –0.46 0.03 0.04
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The unstretched sheet exhibited the characteristic of isotropic lamellae arrangement (f = 0). For the sample of PE-6×1, the crystallographic c-axis was preferably oriented along the uniaxial planar-machine direction, while both the crystallographic a-axis and b-axis took a position close to the TD-ND plane (ND, perpendicular to the MD-TD plane). Simplified schematics are shown in Figure 6b with four different orientation states (I, II, III, and IV), where the crystallographic c-axis is the direction of molecular chain folding, b-axis is the direction of lamellar growth, and a-axis is the direction of lamellar stacking. Hence, it was clear that the crystal orientation of PE-6×1 was in the states of schematics I and II with an approximately equal proportion. With the increase in transverse ratio for sequential biaxial films, fc, MDB decreased and fc, TD increased simultaneously. Subsequently, the crystallographic c-axis approached a state of equal biaxial orientation in the MD-TD plane at about the ratio of 6 × 3 and then gradually oriented toward TD at a higher draw ratio. Also, the orientation factors of the crystallographic a-axis (fa, MDB and fa, TD) were stably negative, which suggested that the crystallographic a-axis was partly perpendicular to the film surface. As for the crystallographic b-axis, the fb, MDB and fb, TD progressively moved to the point of (0, 0) and reached equilibrium at the ratio of 6 × 4, implying the isotropic distribution of the crystallographic b-axis. Besides, the biaxial orientation factors of the crystallographic a-axis, b-axis, and c-axis for PE-sim-6×6 were all suited at the diagonal, stating an equal biaxial orientation in MD and TD after simultaneous biaxial stretching.

From the above analyses, we can draw a conclusion that the orientation states of I and II progressively transform into the orientation states of III and IV with the crystallographic c-axis shifting to TD during the second step of sequential biaxial stretching (shown in Figure 6b). On the other hand, although the biaxial orientation factors of the crystallographic c-axis for both films are close to the point of (0.5, 0.5), the crystal orientation of the PE-6×6 film and PE-sim-6×6 film was different. Combined with the AFM results, the c-axis of the PE-6×6 film oriented either along MD or along TD, while that of the PE-sim-6×6 film is isotropous in the film plane (shown in Figure S4). Hence, there are some differences regarding the mechanical properties of the two films.

Combined with the orientation factors and the parameters extracted from the stress–strain curves shown in Table 1 and Table S2, we can find that there may be some connection between the processing and the structure. The ratio of the yield stress in MD and TD (1.64) for the sample of PE-6×1 is consistent with the ratio of the orientation factors of the crystallographic c-axis (fc, MDB and fc, TD), which means that more lamellae oriented along MD than TD after yielding. Moreover, the ratios of the stress (at the biaxial ratio of 6 × 6) in MD and TD are 0.51 and 0.97 for the films of PE-6×6 and PE-sim-6×6, respectively. The result is also consistent with the ratios of orientation factors of the crystallographic c-axis, which signifies that more lamellae reoriented along TD after sequential biaxial stretching and isotropous lamellae orientation develops after simultaneous biaxial stretching. Hence, to some extent, it may be a simple way to get the state of crystal orientation by the parameters from the stress–strain curve.

3.6. Quantitative Analysis of Crystal Structure

The 2D-SAXS patterns for different samples are shown in Figure 7, where the equator and meridian are defined as the horizontal direction (MD) and vertical direction (TD), respectively. There was an isotropous scattering sign in the sample of PE-0 because of the random crystal distribution. After the stretching in MD, two significant signals were found in PE-6×1, where two streaks across the beam stop appearing in the meridian manifested the presence of the fibrillar structure, and two maxima lobules along the equator revealed the existence of the periodic lamellar stacks preferably oriented in MD.19,39,40 Subsequently, with the increase in λTD, the streak in the meridian disappeared and gradually transformed into the arch-shaped scattering signal, which indicated that the fibrils were either destroyed or oriented against MD. Meanwhile, the maximal signal in the equator weakened progressively, but it is still there until the biaxial draw ratio of 6 × 6. Accordingly, we found that, although the well-arranged lamellae were destroyed during the TD stretching, there were still some lamellar stacks along MD. Besides, compared with the PE-0 film, the scattering signal of the PE-sim-6×6 film was a ring, suggesting more uniform crystal thickness after simultaneous biaxial stretching.

Figure 7.

Figure 7

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2D-SAXS patterns for the sequentially biaxially stretched films prepared at various transverse draw ratios and the PE-sim-6×6 film. The horizontal and vertical arrows indicate MD and TD, respectively.

Long period (Lp), lamellar thickness (Lc), and amorphous thickness (La) for different samples are shown in Figure 8a. The Lc gradually increased before λTD = 4, while it decreased for the ratios of 6 × 5 and 6 × 6, which agreed well with the variation of crystallinity calculated from DSC and WAXD. The change in Lc was due to the coupling effects of the recrystallization for amorphous segments and the destruction for lamellae along the MD.

Figure 8.

Figure 8

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(a) Lp, Lc, and La for the sequentially biaxially stretched films prepared at various transverse draw ratios and the PE-sim-6×6 film. (b) Llateral for the sequentially biaxially stretched films prepared at various transverse draw ratios.

The lamellar lateral dimension (Llateral) of the sequentially biaxially stretched films is shown in Figure 8b, which is determined from the relation41

3.6. 10

where Δq2 is the full width at half maximum shown in Figure S5. The Llateral decreased from 25.1 nm (PE-6×1) to 19.3 nm (PE-6×6) with the increase in λTD, which manifested that the lamellar stacks along MD were gradually destroyed by the transverse stress. Moreover, the results also agreed with the change in Xright shown in Figure 3c.

To investigate the scattering signals in a specific region, the 1D-SAXS curves were acquired by integrating along the meridian and equator direction shown in Figure S6. The signals in the meridian and equator regions represented that the normal of lamellae were parallel and perpendicular to MD, respectively.8,42Figure 9a plots the SAXS scattering intensity of the meridian (Im), equator (Ie), and totality (It) as a function of λTD, where Im gradually increased while Ie decreased during the transverse stretching. Moreover, It increased at first and then decreased after the draw ratio of 3. As shown in Figure 9b, the ratio of Ie to Im sharply declines from 3.16 to 1.63 at λTD = 2 and then progressively decreases to 1.07 with the increase in λTD. In addition, the PE-sim-6×6 film revealed almost the same scattering intensity in the meridian and equator, which implied a more uniform crystal distribution of simultaneous biaxial stretching than that of sequential biaxial stretching. From the above analyses, we can draw a conclusion that some of the lamellar stacks in PE-6×1 reoriented from MD to TD under the transverse loading force.

Figure 9.

Figure 9

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(a) SAXS scattering intensity of the meridian (Im), equator (Ie), and totality (It). (b) Ratio of Ie to Im. (c) Long period in the meridian (Lm) and equator (Le) for the sequentially biaxially stretched films prepared at various transverse draw ratios and the PE-sim-6×6 film.

Figure 9c plots the long period in the meridian (Lm) and the equator (Le) versus λTD. The Lm increased rapidly from 17.2 nm at the transverse draw ratio of 1 to 21 nm at the transverse draw ratio of 6 for sequentially biaxially stretched films, while the Le showed a negligible change during the TD stretching, which might be attributed to the increase in crystalline thickness and extension of the amorphous chains along TD. Moreover, the values of Lm and Le for the PE-sim-6×6 film were similar, suggesting a homogeneous crystal distribution.

3.7. Mechanical and Optical Performance

The mechanical performances regarding Young’s modulus, tensile strength, and elongation at break of different PE films are shown in Figure 10a–c, respectively. After biaxial stretching, a considerable promotion of mechanical properties was observed, with Young’s modulus and tensile strength rising to 445.0 and 102.3 MPa in MD, 490.2 and 114.1 MPa in DD, and 713.3 and 143.8 MPa in TD for the PE-6×6 film. Moreover, Young’s modulus and the tensile strength of the PE-sim-6×6 film were increased to 579.6 and 113.4 MPa in MD, 602.7 and 115.3 MPa in DD, and 626.0 and 124.0 MPa in TD, respectively. On the contrary, the elongation at break of both the PE-6×6 film and PE-sim-6×6 film in MD, DD, and TD decreased dramatically, while it was still over 50%. Besides, the mechanical performances of PE-sim-6×6 in MD, DD, and TD were more uniform than those of PE-6×6. The reason was that a highly oriented fibrillar network formed during biaxial stretching along both MD and TD, which led to higher modulus and strength and lower ductility.

Figure 10.

Figure 10

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(a) Young’s modulus, (b) tensile strength, (c) elongation at break, and (d) luminous transmittance for the films of PE-0, PE-6×6, and PE-sim-6×6.

As shown in Figure 10d, both films of PE-6×6 and PE-sim-6×6 exhibit the excellent optical clarity with T over 80% from the wavelength of 400 to 800 nm. Moreover, the T of PE-sim-6×6 was larger than that of PE-6×6 and even achieved up to 90% at the wavelength of 570 to 800 nm. It was clear that the excellent clarity of the biaxially stretched film was due to the formation of the nanosized crystal structure shown in Figure 2. Besides, some thick fibrils toward MD were observed in the PE-6×6 film but not in the PE-sim-6×6 film, which gave rise to the stronger visible light scattering and then the lower luminous transmittance for the PE-6×6 film.

On the other hand, as shown in Figures 1 and 10b, there may be a correlation between the stretching force from the biaxial stretcher and tensile strength in MD and TD. We can obviously find that the film with the larger stretching force from biaxial stretcher has the higher tensile strength. The reason is that the higher crystal orientation results in the larger stretching force and the higher crystal orientation leads to the higher tensile strength simultaneously. Hence, the larger the stretching force, the higher the crystal orientation, and thus, the higher the tensile strength.

3.8. Mechanism of Structure Evolution and Superior Performance

Based on the analyses above, a schematic is proposed to state the structure evolution and high performance for both PE-6×6 and PE-sim-6×6 shown in Figure 11. For the sequentially biaxial stretched process, the spherulite structure is destroyed into small lamellar pieces, and then those well-ordered lamellae are incorporated into the fibrillar structure in the first tensile step. During the TD stretching, the structure evolution can be classified into three regimes according to the orientation. In the regime I (PE-6×1, PE-6×2, and PE-6×3), some single fibrils are pulled away from the fibrillar bundles, where the morphology with larger fibrillar bundles along MD and the divided individual fibrils against MD is called a vein structure. Moreover, some fibrils along TD gradually form by the partial melting–recrystallization or lamellar rearrangement, and some original lamellae are destroyed by the transverse loading force with the decrease in lateral lamellae dimension. At this situation, the orientation in MD decreases and the orientation in TD increases simultaneously, but the general orientation is still in MD. In the regime II (PE-6×4), more single fibrils are split from the fibrillar bundles, and more fine fibrils orient along TD, leading to the development of the morphology of a nanosized fiber-like network. In addition, the total biaxial orientation for MD and TD reaches a balanced state, where the crystallographic c-axis was parallel to MD or TD in the film plane, the crystallographic b-axis is randomly distributed, and the crystallographic a-axis is partly perpendicular to the film plane. In the regime III (PE-6×5 and PE-6×6), the fibrillar bundles become smaller, and even some separated single fibrils rotate from MD to TD. Although some vein structures are not susceptible to the transverse stretching and are still parallel to MD, the majority of fibrils gradually orient toward TD, which leads to the general orientation shifting to TD. On the other hand, for the simultaneously biaxial stretched process, the spherulites are destroyed into lamellar pieces, and some fibrils gradually form by the melting–recrystallization or lamellar rearrangement. Moreover, the orientation of the above lamellae in PE-sim-6×6 was isotropous in the MD-TD plane, not like the lamellar distribution of PE-6×6 either in MD or TD.

Figure 11.

Figure 11

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Schematic mechanism of structure evolution for the PE-6×6 film and PE-sim-6×6 film.

The technology of biaxial stretching gives rise to a conspicuous enhancement of modules by almost 2 times and strength by about 4.5 times for the TF-BOPE films, which is due to the fact that the crystallographic c-axis is parallel to the MD-TD plane and the molecular chains in the amorphous area are in an extended state. Although the elongation at break decreases dramatically, it is still over 50%, exhibiting good ductility. The PE-sim-6×6 film shows the more uniform mechanical property in MD, DD, and TD than the PE-6×6 film because of isotropous lamellae orientation. Moreover, the luminous transmittance of both films is over 80% due to the formation of nanosized crystals. In addition, some large fibrillar bundles along MD give rise to stronger visible light scattering and then lower luminous transmittance for the PE-6×6 film. The above fibrillar bundles are not found in the PE-sim-6×6 film, resulting in better optical clarity.

4. Conclusions

In this study, we mainly study the structure evolution of PE films in sequential biaxial stretching and also take the simultaneous biaxial stretching as a comparison. Moreover, we test the mechanical and optical performance for both TF-BOPE films. The results show that the structure evolution and performance are different for two processes. During the MD stretching in the sequentially biaxially stretched process, the original spherulites are destroyed into lamellar pieces, and then the above lamellar fragments are incorporated into the fibrillar structure. Subsequently, the separation of the single fibrils takes place, which leads to the development of the nanosized fiber-like network. The general orientation reaches a balanced state in MD and TD at the biaxial ratio of 6 × 4. On the other hand, the simultaneously biaxially stretched film (PE-sim-6×6) shows the isotropous lamellae distribution in the film plane. More importantly, both films of PE-6×6 and PE-sim-6×6 have a noteworthy increment in modulus and strength and also exhibit the excellent optical clarity. In addition, because of more uniform crystal orientation, the mechanical performance of PE-sim-6×6 is isotropous in MD, DD, and TD.

Acknowledgments

This work was supported by the National Natural Science Foundation of China (NSFC 51503134 and 51721091).

Supporting Information Available

The Supporting Information is available free of charge at https://pubs.acs.org/doi/10.1021/acsomega.9b03250.

  • Schematic diagrams of biaxial stretching and tensile spline, curves of SSA fractionation thermal procedures and results, SEM images of selectively etched samples, WAXD parameters of different samples, orientation states and the corresponding orientation factors, stress of biaxial stretching, process of obtaining Δq2, and 1D-SAXS integration curves (PDF)

The authors declare no competing financial interest.

Supplementary Material

ao9b03250_si_001.pdf (525.9KB, pdf)

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ao9b03250_si_001.pdf (525.9KB, pdf)

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