Structure Evolution of Polyethylene in Sequential Biaxial Stretching

Article Cite This: Ind. Eng. Chem. Res. 2019, 58, 12419−12430

pubs.acs.org/IECR

Structure Evolution of Polyethylene in Sequential Biaxial Stretching along the First Tensile Direction Qiang Chen, Dandan Chen, Jian Kang, Ya Cao, and Jinyao Chen* State Key Laboratory of Polymer Materials Engineering, Polymer Research Institute of Sichuan University, No.24 South Section 1, Yihuan Road, Chengdu, 610065, P. R. China

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S Supporting Information *

ABSTRACT: Two linear low-density polyethylene samples (PE-A and PE-B), with similar melt index and molecular weight but different molecular structures, were used to explore the structural evolution during sequential biaxial stretching. Only PE-A could be successfully stretched into a film with the biaxial draw ratio of 6 × 6. The interlamellae crossing adjacent lamellar stacks were observed in PE-A but not in PE-B, which could increase the lateral tie connection. Moreover, the amount of fibril was less in PE-A than that in PE-B and the fibrillar structure lacked tie chains between each other. In addition, the newly formed lamellae in PE-A (derived from fine lamellae) could lock the molecular chains of different crystals, which increased the lateral connection via anchoring the tie chains from adjacent lamellae. Therefore, the above specific structures increased the lateral connection and were beneficial for biaxial stretching. Furthermore, the relationship between molecular structure and condensed structure was also established.

1. INTRODUCTION Polyethylene (PE) films are widely used in packing and agriculture applications because of their excellent clarity, mechanical property, and barrier property as well as low price. From a commercial perspective, PE films are manufactured by two technologies. One is the melt blowing process,1,2 which blows up a molten PE tube in a circular orientation and simultaneously controls the rate of different rollers in a longitudinal orientation. Another is the tenter-frame process, which draws a cast sheet in two vertical directions with two successive procedures (sequential biaxial stretching) or a single procedure (simultaneous biaxial stretching). Among the tenter-frame processes, sequential biaxial stretching is more universally used in industrial manufacture due to higher productivity and lower equipment cost. The main advantages of tenter-frame biaxially oriented polyethylene (TF-BOPE) films compared with blowing films are the higher chain orientation and more flexible processing conditions, which lead to superior mechanical property, optical property, and processability.3 Moreover, due to the excellent low-temperature toughness and puncture resistance, TF-BOPE films are extremely suitable for the field of frozen food packaging. Although the uniaxial deformation of PE had been extensively investigated,4−12 there was little research about the structure−processing−property of BOPE films, especially in the microstructure evolution. Ajji et al.13 studied the orientation and shrinkage of seven linear-low density polyethylene (LLDPE) films. The results showed that the © 2019 American Chemical Society

comonomer content and high molecular weight tails have an important effect on the orientation of the BOPE films. Ratta et al.14 analyzed the stretching process of a cast sheet for highdensity polyethylene (HDPE) in the machine direction (MD, parallel to the extrusion direction). They revealed that crystalline morphology changes from the spherulite structure into lamellar stacks by localized melting and recrystallization, bulk melting, or a combination of the above ways, but subsequent stretching in the transverse direction (TD, perpendicular to MD) is not shown. Ajji et al.15 used the techniques of infrared spectroscopy, X-ray pole figures, and birefringence to measure the orientation of BOPE film. Moreover, different carbon fillers, such as carbon nanotubes, graphene, and carbon black, were incorporated to prepare PE nanocomposites. The biaxial stretching films made by those nanocomposites exhibit excellent tensile, barrier, and electronic conductivity properties.16,17 The structure evolution of other polymers was also studied, which could be divided into two categories according to the state of the original cast sheet. The first type is the semicrystalline state, such as polypropylene (PP). Lupke et al.18 had investigated the structure evolution of biaxially oriented polypropylene (BOPP) films and concluded that the drawing in MD transforms original spherulites into lamellar Received: Revised: Accepted: Published: 12419

March 30, 2019 June 2, 2019 June 11, 2019 June 11, 2019 DOI: 10.1021/acs.iecr.9b01733 Ind. Eng. Chem. Res. 2019, 58, 12419−12430

Article

Industrial & Engineering Chemistry Research

10 MPa for 5 min and then quenched into the water at room temperature. The sheets for uniaxial stretching were cut into a rectangle with 600 mm in length and 10 mm in width and then they were stretched at 114 °C with a strain rate of 300 mm/min (25%/s). The samples with the different draw ratio of 1X, 2X, 3X, 4X, 5X, 6X were labeled “PE-A/B-1/2/3/4/5/6”. For example, the sample of PE-A-5 meant that LLDPE-A was uniaxially stretched to a draw ratio of 5X. BOPE films were prepared on a biaxial stretcher KARO IV (Brückner Maschinenbau B, D-83313 Siegsdorf, Germany). The sheets 90 mm × 90 mm in size were sequentially stretched at 114 °C with a strain rate of 25%/s, and the final biaxial draw ratio was 6 × 6. In addition, the first tensile direction of sequential biaxial stretching was defined as TD-I and the second tensile direction was defined as TD-II (perpendicular to TD-I). 2.2. Gel Permeation Chromatography (GPC). Hightemperature GPC (PL220, Great Britain) was used to test the number-average molecular weight (Mn), weight-average molecular weight (Mw), and molecular distribution (Mw/Mn) of polyethylene. The solvent, solution temperature, and dissolution time were 2,4,6-trichlorobenzene, 140 °C, and 6 h, respectively. To prevent degradation, 0.3% weight of Irganox1010 was added during the process of dissolution. Besides, the IR detector and viscosity detector were employed and the flow rate was 1 mL/min. 2.3. Fourier Transform Infrared Spectroscopy (FTIR). The copolymer type of two samples was measured by FTIR spectrometer (Thermo Nicolet 560) with an accumulation of 32 scans and a resolution of 4 cm−1. Moreover, ethyl branches were characterized by the absorbance at 770 cm−1 and butyl branches were absorbed in the wavenumber of 893 cm−1.26 2.4. Differential Scanning Calorimetry (DSC). Thermal analysis was carried out with a DSC 3+ Stare System (Mettler Toledo, Switzerland). Temperature and heat flow scales were performed using high pure indium and zinc as the standard. A nominal 5 mg of sample was scanned from 25 to 180 °C under a nitrogen atmosphere with a heating/cooling rate of 10 °C/ min.27 Moreover, the crystallinity (XDSC) of PE is defined as follows

stacks through partial melting, and then the drawing in TD transforms those stacked lamellae into a fibrillar network by the action of crystal slip. Nie et al.19 observed the fiber-like network morphology of BOPP film by atomic force microscopy (AFM) and proposed that selecting an appropriate stretching ratio of MD and TD can control the structure of BOPP. Elias et al.20 pointed out that the crystalline morphology of the BOPP film is composed of an orthogonal shish structure and small lamellae, the c-axis of which is parallel to the plane of MD and TD. The second type is the amorphous state, such as poly(ethylene terephthalate) (PET), polystyrene (PS), and polylactic acid (PLA), which can be quenched into a quasiamorphous state by chill roll. Li et al.21 studied the straininduced crystallization (SIC) of biaxially oriented poly(ethylene terephthalate) (BOPET) and demonstrated that increasing the mobility of the molecular chain is significant for crystallization. Cakmak et al.22−24 had investigated the relaxation behavior of BOPET and divided the relaxation behavior into three regimes according to the degree of SIC. They also suggested that the draw ratio of MD should be suitable in order to control the content of the fibrillar structure, which influences the processability in TD stretching. In addition, the structure evolution of PLA and PET is also investigated during biaxial stretching.21,25 The process of SIC takes place in the first step of biaxial drawing and then the crystalline structure is destroyed during TD stretching, which leads to the formation of a fibrillar network. Some big challenges remain in the industrial manufacture of TF-BOPE films. On the one hand, BOPP and BOPET films have been widely used in our daily life while there are hardly special materials for BOPE. On the other hand, up to now, the mechanism of structure evolution for TF-BOPE is still unclear. In this study, two types of LLDPE resins with similar melt flow rate and molecular weight are used for sequential biaxial stretching. Offline differential scanning calorimetry (DSC), wide-angle X-ray diffraction (WAXD), small-angle X-ray scattering (SAXS), and scanning electron microscopy (SEM) were used to characterize the differences of structural evolution for two samples. Our work shows that the interlamellae, the content of fibril, and the newly formed thick lamellae are the significant factors for biaxial stretching. Moreover, gel permeation chromatography (GPC), Fourier transform infrared spectroscopy (FTIR), Raman spectroscopy, and successive self-nucleation and annealing fractionation (SSA) were employed to analyze the differences in molecular architecture for two samples. The relationship between the condensed structure and molecular structure for biaxial stretching was also established. To our knowledge, it is the first time that an appropriate evolution model of PE in sequential biaxial stretching is proposed. In addition, the structure evolution in the second tensile direction of biaxial stretching will be discussed in subsequent publication.

XDSC =

ΔHm ΔHm0

(1)

where ΔHm is the melting enthalpy for tested samples, and ΔH0m = 287.3 J/g is the melting enthalpy of 100% crystallinity for PE.28 Besides, the thermal procedures of successive selfnucleation and annealing (SSA) fractionation29 are shown in Figure S1 of the Supporting Information. 2.5. Scanning Electron Microscopy (SEM). To observe the surface morphology, the samples were chemically etched by a mixed solution of concentrated sulfuric acid, concentrated orthophosphoric acid, and distilled water (10:4:1, v/v/v), which also contained 1.3% weight of potassium permanganate.30,31 After being etched for 24 h at room temperature, the PE samples were cleaned by hydrogen peroxide and water. Lastly, the microtopography was obtained by a FEI Inspect F SEM instrument under an acceleration voltage of 10 kV.32 2.6. Two-Dimensional Wide-Angle X-ray Diffraction (2D-WAXD). 2D-WAXD experiments were conducted on a Bruker D8 Discover diffractometer equipped with a detector of general area detector diffraction system (GADDS). The Cu Kα

2. EXPERIMENTAL SECTION 2.1. Materials and Sample Preparation. Two kinds of linear low-density polyethylene (LLDPE) were used in this study. LLDPE-A (PE-A) was supplied by Guangdong Decro Film New Materials Co., Ltd. with the melt flow rate (MFR) of 1.8 g/10 min (190 °C, 2.16 kg). LLDPE-B (PE-B) was purchased from Formosa Plastic Group with the melt flow rate of 2.0 g/10 min (190 °C, 2.16 kg). The resin pellets were molded into 1 mm thick sheets at 180 °C with the pressure of 12420

DOI: 10.1021/acs.iecr.9b01733 Ind. Eng. Chem. Res. 2019, 58, 12419−12430

Article

Industrial & Engineering Chemistry Research Table 1. Molecular Architectures of PE-A and PE-B samples

Mna (g/mol)

Mwa (g/mol)

Mw/Mna

comonomer typeb

XS (%)

Tcc (°C)

Tmc (°C)

XDSCc (%)

PE-A PE-B

24900 32220

116500 108800

4.7 3.4

1-hexene 1-butene

3.7 ± 0.2 8.7 ± 0.3

115.1 106.1

128.5 122.8

51.0 44.6

a

Data were measured by GPC. bData were measured by FTIR. cData were measured by DSC.

Figure 1. SSA heating curve, corresponding fitted peak, and MSL calibration curve of (a) PE-A and (b) PE-B.

X-ray radiation (λ = 0.154 nm) was used. The experimental voltage was 40 kV and the experimental current was 40 mA, respectively. For a given hkl plane, the degree of orientation for various samples is defined by Hermans’ orientation parameter (f)33 f=

3⟨cos2 φ⟩ − 1 2

⟨cos2 φ⟩ =

∫0

π /2

Lc =

(2)

π /2

I(φ) sin φ dφ

(3)

where φ is the angle between the set of hkl plane and stretching direction, and I(φ) is the scattering intensity. In addition, the orientation parameters of PE were acquired via the orientation of (110) diffraction plane.34,35 2.7. Two-Dimensional Small-Angle X-ray Scattering (2D-SAXS). The 2D-SAXS experiments were conducted with the Xeuss 2.0 system (Xenocs, France). The 2D-SAXS images were collected with a Pilatus 300 K detector of Dectris, Swiss (680 pixels × 600 pixels, pixel size = 172 μm). The Cu Kα Xray radiation (λ = 0.154 nm) and scatterless collimating slits were used during all the measurements. Moreover, the sampleto-detector distance was fixed at 2500 mm, which provides a scattering vector q (q = (4π sin θ)/λ)) range from 0.05 to 1.15 nm−1. The parameters of the scattering geometry were calibrated using silver behenate. The average long period of the stacked lamellae (Lp) is calculated from Bragg’s law36,37

Lp =

2π q1max

ρc

(5)

La = Lp − Lc

(6)

X 1 − Xc 1 = c + ρ ρc ρa

(7)

where ρ is the density of the tested samples, ρa = 0.855 g/cm3 is the density of amorphous phase and ρc = 1.003 g/cm3 is the density of orthorhombic crystal for PE.38 Besides, it is worth noting that the equation hypothesizes a much larger lateral size and length of the lamellar crystal in relation to thickness. 2.8. Raman Spectroscopy. Raman spectroscopy was used for quantitative analysis of the weight fraction of interphase (Xi) between crystalline and amorphous layers. Raman spectrum was recorded using a DXRxi Raman imaging microscope (ThermoFisher Scientific) with laser radiation of 532 nm. Moreover, the wavenumber from 1000 to 1700 cm−1 was recorded to obtain the crystal weight fraction (Xc), the amorphous fraction (Xa) and the interphase fraction (Xi) according to the Strobl & Hagedorn method.39

2

I(φ) sin φ cos φ dφ

∫0

Xc L pρ

3. RESULTS AND DISCUSSION 3.1. Molecular Structure. For a given processing condition, the different molecular structures of samples will produce various crystal morphologies and molecular arrangements, which subsequently influence the processability and usability of the final film product. Therefore, the molecular architectures of LLDPE, containing molecular weight, molecular weight distribution, comonomer type, comonomer content, comonomer distribution, and so on, are very significant for a biaxial stretching film. The data of the molecular architectures for PE-A and PE-B are shown in Table 1 and the corresponding curves are exhibited in Figure S2 of the Supporting Information. The Mw of PE-A and PE-B was almost identical, but the Mw/Mn of PE-

(4)

where q1max is the maximum scattering vector along tensile direction (detail in Figure 8). The lamellar thickness (Lc) and amorphous thickness (La) are defined according to 12421


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Industrial & Engineering Chemistry Research A was larger than that of PE-B. In addition, the comonomer types were 1-hexene and 1-butene for PE-A and PE-B, respectively. Xylene solvent (XS) fractionation at room temperature is a relative simple approach compared with NMR to characterize the comonomer content of LLDPE.40 Moreover, the larger XS fraction represents the higher comonomer content and vice versa. The XS of PE-A (3.7%) was far less than that of PE-B (8.7%), which indicated that there were less comonomers in the PE-A chains. Hence, the crystallization temperature (Tc), melting point (Tm) and XDSC of PE-A tested by DSC were larger than those of PE-B. The distribution of short chain branching (SCB) for ethylene/α-olefin copolymers has been investigated by some quantitative methods.29,41 The SSA fractionation was used in this study and a multiple melting endotherm was gained after the sample was heated from 25 to 170 °C. The average methylene sequence length (MSL) is calculated using the following equations:42 ln(X ) = 0.3451 −

MSL =

2X 1−X

142.2 T

Figure 2. Stress−strain curves of PE-A and PE-B.

Table 2. Mechanical Properties of PE-A and PE-B (8)

(9)

where T is the melting point in K and X is the mole fraction of methylene. Plots of SSA heating curves, corresponding fitted peaks, and calibration curves (MSL-Tm) are shown in Figure 1. It is worth mentioning that the higher Tm corresponds to longer MSL, which then forms thicker chain-folded lamellae and vice versa. There were some notable differences in the SSA heating curves for PE-A and PE-B. On the one hand, although the position of each melting peak from 73.9 to 125.5 °C was nearly the same, a higher temperature peak (peak 1) about 130.5 °C only appeared in PE-A. That was to say, PE-A owned some longer crystallizable chain segments, which could develop thicker chain-folded lamellae. On the other hand, the comonomer distribution of PE-A was more nonuniform than that of PE-B. The MSL of PE-A mainly existed in the relative high and low areas showing a bimodal feature, but the MSL of PE-B focused on the relative medium area. In other words, the thicker and thinner lamellae coexisted in the sample of PE-A. 3.2. Condensed Structure. The condensed structures have an important influence on the final film properties. Moreover, crystalline morphology, orientation, and size were characterized by SEM, WAXD, SAXS, and DSC. 3.2.1. Tensile Deformation Behavior. The stress−strain curves (exhibited in Figure 2) of both PE-A and PE-B displayed a typical yielding and necking feature. As shown in Table 2, Young’s modulus and yield strength of PE-A were larger than those of PE-B owing to greater crystallinity and lamella thickness (detailed in section 3.2.4.) for PE-A. Moreover, the tendency of strain hardening was faster for PE-B, which meant that more fibrils were developed in PE-B during stretching. In addition, the biaxial stretching experiments showed that only PE-A could be successfully stretched into a film with the biaxial draw ratio of 6 × 6 while PE-B failed. The stress−strain curves and the photo of BOPE film are exhibited in Figure S3 of the Supporting Information. 3.2.2. Morphology Evolution during Stretching. The spherulitic morphologies of PE-A and PE-B, the average diameters of which are about 1 and 3 μm, respectively, are

samples

Young modulus (MPa)

yield strength (MPa)

PE-A PE-B

10.3 6.7

1.58 1.52

observed in Figure 3a and Figure 4a. The crystallographic morphologies of different draw ratios (from 1X to 6X) for PEA samples are shown in Figure 3b to Figure 3g. At the draw ratio of 1X, the spherulites were destroyed and much of the fragmented lamellae partially oriented along TD-I. For the draw ratio up to 2X and 3X, the remnants derived from the original spherulitic structure decreased and the stacked lamellae paralleling to TD-I gradually became the main morphology. Moreover, a small amount of fibrillar structures were also found, which might be attributed to the partial melting recrystallization or lamellar rearrangement. When PEA was drawn to the ratio of 4X and 5X, the structure of stacked lamellae became more regular and some specific lamellae crossing through two or more rowlike lamellar stacks were observed. The specific lamellae, like a bridge connecting adjacent lamellar stacks, were called interlamellae in this study, and could act as the stress transmission units when the sample was stretched along TD-II. Therefore, the structure of the interlamellae was beneficial for sequential biaxial stretching. Lastly, at the draw ratio of 6X, more stacked lamellae transformed into fibrils. Meanwhile, a new morphology was also observed, which looked like a slender cavity between adjacent fibrils. In addition, even at the high draw ratio, there remained some interlamellae between two neighboring fibrils, which acted as a lateral tie connection in TD-II stretching. The crystallographic morphology evolution of PE-B at different draw ratios(from 1X to 6X) is exhibited in Figure 4b to Figure 4g. At the draw ratio of 1X, the spherulitic morphology disappeared and plenty of crystal fragments oriented along TD-I. When the PE-B samples were drawn from 2X to 6X, a new morphology progressively formed, which contained a large number of stacked lamellae and highly oriented fibrils like the structure of “shish-kebab”. Moreover, compared with PE-A, some slender cavities were even observed under the low draw ratio. It is clear that cavities appear with the formation of a fibrillar structure, which leads to less lateral connection. Therefore, when the PE-B sample is drawn in TD-II, few lateral tie chains 12422


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Figure 3. SEM images of selectively etched (a) PE-A-0; (b) PE-A-1; (c) PE-A-2; (d) PE-A-3; (e) PE-A-4; (f) PE-A-5; (g) PE-A-6. The yellow arrow indicates the tensile direction.

Figure 4. SEM images of selectively etched (a) PE-B-0; (b) PE-B-1; (c) PE-B-2; (d) PE-B-3; (e) PE-B-4; (f) PE-B-5; (g) PE-B-6. The yellow arrow indicates the tensile direction.

will act as stress transmitters. On the contrary, PE-A includes some interlamellae at the high stretching ratio (6X), which is beneficial for sequential biaxial stretching. 3.2.3. Orientation Structure. The (110) and (200) diffraction rings of the orthorhombic PE unit cell are shown in Figure 5, where the meridian and equator are defined as the horizontal direction (tensile direction) and vertical direction, respectively. The homogeneous diffraction rings were observed

in the unstretched samples of both PE-A and PE-B. Moreover, the corresponding azimuthal intensity curves of the (110) diffraction plane are displayed in Figure 6a and Figure 6b, where the azimuthal angle of 180 ° represents the equatorial direction. For PE-A, when the deformation was at 1X, the patterns of (110) and (200) diffraction rings began to focus along the equator. In addition, a weak peak appeared in the azimuthal 12423


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Figure 5. 2D-WAXD patterns of PE-A and PE-B with different draw ratios from 0 to 6. The white arrow indicates the tensile direction.

Figure 6. Azimuthal intensity curves of (110) diffraction plane for (a) PE-A and (b) PE-B at various draw ratios.

Table 3. Hermans’ Orientation Parameter of PE-A and PE-B at Various Draw Ratios (DR) samples

DR = 0

DR = 1

DR = 2

DR = 3

DR = 4

DR = 5

DR = 6

PE-A PE-B

0 0

0.26 0.71

0.50 0.80

0.52 0.83

0.78 0.86

0.86 0.89

0.91 0.94

angle of 180 ° (Figure 6a) indicating the partial orientation of the crystal structure.43 For the draw ratio of 2X and 3X, both (110) and (200) diffraction arcs were situated at the equator, and the azimuthal intensity profiles showed obvious peaks at 180° (Figure 6a), which stated that the orientation degree increased during stretching. Subsequently, when PE-A was stretched from 4X to 6X, the diffraction arcs of (110) and (200) transformed into two pairs of bright concentrated diffractions in the equator indicating the development of fibrillar crystals. At the same time, the corresponding peaks at the azimuthal angles of 180 ° (Figure 6a) became stronger. According to the model of “Keller/Machin II”,44 the intense and defined diffractions of (110) and (200) located at the equator implied that the lamellar c-axis is parallel to the tensile direction and both b-axis and a-axis are perpendicular to the tensile direction. Therefore, the c-axis of PE-A lamellae oriented along the tensile direction at a high draw ratio. For PE-B, the strong and defined diffractions focusing on the equatorial direction could be found in all patterns of the draw ratio of from 1X to 6X. Simultaneously, the corresponding azimuthal intensity curves displayed a strong peak at 180° (Figure 6b). The difference between the two samples was that compared with PE-A, PE-B displayed a higher orientation and formed a fibrillar structure at the lower draw ratio. In other words, the crystals in PE-B were more inclined to orient toward the tensile direction. The Hermans’ orientation parameter (f) is calculated from the azimuthal intensity curves of (110) diffraction plane shown

in Table 3 and Figure 7. We take the value f = 1 for completely parallel orientation, f = 0 for random distribution, and f = −0.5 for completely perpendicular orientation. The f of PE-A could be divided into three regimes: Regime I (DR = 1), where the f was relatively low (about 0.26) indicating the weak orientation of PE-A at this draw ratio. Regime II (DR = 2 and 3), where f was relatively medium (about 0.50) implying the development

Figure 7. Hermans’ orientation parameter of PE-A and PE-B at various draw ratios. 12424


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Figure 8. 2D-SAXS patterns of PE-A and PE-B with different draw ratios from 0 to 6. The white arrow and rectangle indicate the tensile direction and scattering bar, respectively.

Figure 9. (a) Long period (Lp), lamellar thickness (Lc), amorphous thickness (La), and (b) lateral lamellae dimension (Llateral) for PE-A and PE-B at different draw ratios.

From the above results, compared with that of PE-B, the comonomer content of PE-A is less and the comonomer distribution of PE-A is more nonuniform, which results in high MSL. In other words, some longer crystallizable chain segments exist in PE-A and then develop into thicker chainfolded lamellae. The lamellar lateral dimension (Llateral) is present in Figure 9b which is described as follows46

of well-arranged stacked lamellae. Regime III (DR = 4 to 6), where f was relatively high (larger than 0.7) speculating the formation of fibrillar structure. As for PE-B, all the f values during stretching (DR = 1 to 6) were larger than 0.7, which meant that the fibrillar structure even formed at the low draw ratio. Moreover, the results of the orientation was almost consistent with the analysis of SEM. 3.2.4. Quantitative Analysis of Crystal Structure. The 2DSAXS patterns of PE-A and PE-B at different draw ratios are displayed in Figure 8. According to the previous researches,45−47 two significant signals can be found in the SAXS patterns. First, two maxima lobules along the meridian reveal the existence of lamellar stacks; second, two streaks across the beam stop appearing in the equator manifesting the presence of the fibrillar structure. The evolution of the SAXS patterns for PE-A and PE-B indicated the coexistence of the stacked lamellae and fibrillar structure during the stretching. However, the pattern of PE-A at the draw ratio of 1X was different from others. The elliptic scattering signal represented a transition state of structure evolution, which contained both the spherulite fragments and oriented lamellae. Long period (Lp), lamellar thickness (Lc) and amorphous thickness (La) are shown in Figure 9a, where Lp, Lc, and La increase gradually during the stretching. At a high temperature, the mobility of the molecular chains were relatively high and more coiled segments transformed into chain-folded lamellae via the strain-induced disentanglement.48 Moreover, owing to the extension of the tie chains in the amorphous area, the distance of the amorphous layer between two adjacent lamellae became farther after stretching. Besides, the Lc of PE-A was larger than that of PE-B at each deformation state.

L lateral =

2π Δq2

(10)

where Δq2 is the width of the peak at half height. The process to obtain Δq2 is provided in Figures S4 of the Supporting Information. The Llateral of both PE-A and PE-B decreased during stretching. After stretching, the effect of continuous shear, slip, and fragmentation transformed the lamellae into smaller blocks, which led to the decrease of lateral dimension. In addition, the Llateral of PE-A was larger than that of PE-B, which was due to the larger lamellar thickness49−51 and the less fraction of interphase for PE-A. Humbert et al.52 found that the fibrillar diameter retained the memory of the lamellar topological defects (cilia, loose loop, irregular chain-folds, segregated comonomer, etc.) which generally gathered in the interphase between the crystalline and amorphous layers.53,54 Moreover, the topological defects control the local stress distribution on the lamellar surface, which may generate stress concentration and lamellar fragmentation.5,55−57 Therefore, the more content of interphase there is, the greater is the concentration of topological defects, and thus the smaller is the lateral dimension of the lamellar blocks. In this study, the content of the interphase was determined by Raman 12425


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PE-A owns thicker lamellae and thinner interphase layers than PE-B. After stretching, some lamellae with larger lateral dimension (interlamellae) remains in PE-A. Moreover, the interlamellae cause a stronger lateral connection between adjacent lamellar stacks, the effect of which makes it beneficial for TD-II stretching. The process of fibrillation is done in the following steps: first, the original spherulites are destroyed and shifted to a suitable place; second, the lamellae are broken into smaller blocks; last, the lamellar stacks are incorporated into fibrils.8 The relative volume content of fibrils (Φfibril) is determined by the integrated scattering intensities of the equatorial streak and total scattering signal in the 2D-SAXS image according to the following equations:62,63

spectroscopy and SAXS. The result was shown in Table 4 and the detailed methods for calculating the fraction of interphase Table 4. Weight Fraction of Interphase According to Raman and Volume Fraction of Interphase According to SAXS for PE-A and PE-B

a

sample

Xi (wt%)a

Xi (vol%)b

PE-A PE-B

5.6 10.2

10.2 13.1

Data were calculated from Raman. SAXS.

b

Date were calculated from

by Raman spectroscopy and SAXS were provided in Figures S5 of the Supporting Information.39,58,59 Both the weight and volume fraction of the interphase for PE-B were larger than those of PE-A, which indicated that more topological defects formed in the lamellar surface for PE-B. Hence, the lamellae of PE-B were broken into smaller pieces during stretching. From the perspective of molecular structure, PE-B has more comonomer content than PE-A, which provides a larger probability to exclude the comonomer from growing crystals. Therefore, the topological defects are more likely to develop in the interphase layers for PE-B.60,61 As is shown in Figure 12, combined with the above SEM and SAXS results, it is clear that

20°

∫−20° I(q , φ) dq dφ

(11)

Ifibril Itotal

(12)

Ifibril = 2 Φfibril =

where q is the scattering vector, φ is the azimuthal angle, and Ifibril and Itotal represent the intensities of fibrils and total scattering signal, respectively. Moreover, Ifibril is proportional to the total volume of fibrils. The Φfibril of PE-A and PE-B at various draw ratios is exhibited in Figure 10a, where the Φfibril

Figure 10. (a) Relative volume content of fibril (Φfibril); (b) diameter of fibril; (c) average length of fibril (Lfibril) for PE-A and PE-B at different draw ratios. The inset shows the integrated azimuthal angle (φ) from −20° to 20°. 12426


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Figure 11. Melting curves of (a) PE-A and (b) PE-B at various draw ratios. (c) The corresponding crystallinity of PE-A and PE-B. (d) Crystallinity of PE-A integrated at the area of peak at 127.7 °C.

Moreover, the Dfibril of the PE-A is larger than that of PE-B at each tensile state. The average length of fibril (Lfibril) along the tensile direction can be determined via the formula put forward by Ruland and Perret.65−67

gradually increased during deformation indicating that more lamellae transformed into the fibrillar structure. Besides, PE-B exhibited higher Φfibril than PE-A, which suggested that there was more fibrillar structure in PE-B. The result was consistent with the analyses of stress−strain curves and SEM morphology mentioned above. The diameter of fibrils (Dfibril) is obtained from the model of cylindrical symmetry where the radius size of the scattering unit is calculated by applying the Guinier’s approximation function to the distribution of scattering intensity along the equator streak.64 ij −q 2R g2 yz I = I0 expjjjj 2 zzzz j 3 z k {

(13)

Dfibril = 2R g

(14)

Bobs = Bφ +

2π Lfibril q

(15)

where Bobs and Bφ represent the integral breadth of the azimuthal integrated curve and misorientation of fibrillar structure, respectively. The method for obtaining Bobs and Lfibril is provided in Figures S7 of the Supporting Information. As is shown in Figure 10c, the Lfibril of PE-A is 484 nm at the draw ratio of 2X and increases to 637 nm at the draw ratio of 6X. For PE-B, Lfibril increased from 305 to 376 nm when samples were stretched from 1X to 6X. The differences between the two fibrillar structure of PE-A and PE-B are summarized as follows: there are higher relative volume content, larger diameter, and larger average length of fibrils in PE-A than in PE-B. Therefore, the amount of fibril in the PE-B sample is more than that of the PE-A. It is clear that few tie chains exist between adjacent fibrils, which leads to much less lateral connection for PE-B.

where Rg is the radius of gyration and q2 is the distribution of scattering intensity along the equatorial streak. The detailed method for calculating Rg is provided in Figures S6 of the Supporting Information. As shown in Figure 10b, the Dfibril gradually decreases with increasing draw ratio, which implies that the lamellae in the fibril are broken into smaller blocks. 12427


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Industrial & Engineering Chemistry Research 3.2.5. Thermal Analysis. The melting curves of both PE-A and PE-B at various draw ratios are shown in Figure 11a and Figure 11b, respectively. For PE-A, all the thermograms exhibited a multiple-melting behavior. For the heating curve of PE-A-0, the melting point at 127.7 °C and the small shoulder peak at 104.7 °C were found, which suggested that there were two kinds of lamellae with different thicknesses. The lamellar thicknesses were calculated according to the Thomson-Gibbs equation.68 When PE-A was stretched to the draw ratio of 6X, the shoulder peak (104.7 °C, 4.7 nm) moved to a higher temperature (118.1 °C, 7.4 nm). Moreover, the relative area of the melting peak at 118.1 °C increased gradually indicating that more original fine lamellae thickened during stretching. As shown in Figure 11c, the crystallinity of PE-A slightly decreases from 47.4% to 46.7% when the sample was stretched to the ratio of 1X. At this draw ratio, the sample was in the state of yielding with a number of spherulites transforming into small lamellar blocks, which lead to the decline of crystallinity. Subsequently, the crystallinity of PE-A gradually increased to 49.0% at the draw ratio of 6X due to the partial melting recrystallization or the rearrangement of coiled chains. To manifest that the melting peak (118.1 °C) of PE-A mainly came from the shoulder peak (104.7 °C), we separated all heating curves into left and right areas along the dotted line at 122.2 °C (shown in Figure 11a). The higher melting temperature peak at 127.7 °C was called the “area of peak at 127.7 °C”. Furthermore, the areas of the peak at 127.7 °C are integrated and transformed into crystallinity as shown in Figure 11d. It was clear that the crystallinity of the area of peak at 127.7 °C exhibited a negligible change during stretching. Hence, we could draw a conclusion that the newly formed lamellae (7.4 nm) were mainly derived from original fine lamellae (4.7 nm). Compared with those of PE-A, some differences of thermograms for PE-B were observed. There was only a melting peak in all heating curves and no shoulder peak at low temperature (shown in Figure 11b). In addition, the melting point decreased from 122.8 to 119.2 °C after stretching, which supported that some thick lamellae were destroyed into small pieces. The crystallinity of PE-B gradually increased from 42.3% to 44.7%, which was less than that of PE-A at each draw ratio. From the above analyses, we can conclude that the thick lamellae (7.4 nm) of PE-A mainly form by the melting and recrystallization of fine lamellae (4.7 nm). The corresponding comonomer distribution of PE-A is exhibited in Figure 1, where the MSL mainly exists in the relative high and low areas showing a bimodal feature. Hence, both thick and thin lamellae develop in PE-A owing to this specific molecular structure. As is shown in Figure 12, the newly formed lamellae can lock molecular chains between adjacent lamellae, which increases the lateral connection via anchoring tie chains and is beneficial for TD-II stretching. 3.3. The Deformation Mechanism for Biaxial Stretching. With the structural analyses above, we propose a deformation model to illustrate why only PE-A can be biaxially stretched. The structure of unoriented PE can be simplified as several elements. On the whole, the system is filled with isotropous spherulites and amorphous chains. In the interior of the spherulites, there are some tie molecules and entangle chains between two parallel lamellae as stress transmitters. During the deformation, the lamellae are broken into small pieces, which are incorporated in fibrils. Moreover, the fibrils

Figure 12. Scheme of schematic mechanism under TD-I stretching.

own a lot of axial connections but very few lateral tie connections, leading to the failure in TD-II stretching. As shown in Figure 12, three scenarios of structure evolution may occur during stretching. In scenario I, PE-A owns thicker lamellae and thinner interphase layers (topological defects) as compared with PE-B. Hence, the lamellar lateral dimension of PE-A is larger than that of PE-B. In addition, some interlamellae crossing adjacent lamellar stacks are also observed in PE-A, which increases the lateral tie connection. In scenario II, the amount of fibril in the PE-B sample is more than that in the PE-A. It is clear that the fibrillar structure consists of highly oriented lamellae and taut tie molecules along the tensile direction, but the tie chains between neighboring fibrils are scanty. Therefore, the overabundant fibrils are an important factor leading to the failure in TD-II stretching. In scenario III, some fine lamellae exist in the PE-A sample, which develop into thicker lamellae by melting and recrystallization during stretching. The newly formed lamellae can lock molecular chains between adjacent lamellae. Hence, such a structure increases the lateral connection via anchoring tie chains of different lamellae. The differences in structural evolution during stretching are mainly determined by molecular structure. Compared with PEB, PE-A owned longer crystallizable chain segments which then developed thicker chain-folded lamellae. In addition, PE-B has more comonomer content providing a larger probability to exclude the comonomer from growing crystals. Therefore, the topological defects are more likely to develop in the interphase layers for PE-B. On the contrary, PE-A owns a thicker crystalline phase and thinner interphase layer than does PE-B, which leads to the larger lateral dimension of lamellae and the more content of fibrils after stretching. Besides, the heterogeneous distribution of comonomer allowed PE-A to form some short MSL segments, which subsequently developed into fine lamellae. After deformation at high temperature, the fine lamellae turned into thick lamellae by melting and recrystallization, acting as a lateral connection during TD-II stretching. In short, it is significant to design a suitable molecular structure for BOPE films.

4. CONCLUSIONS In this study, we have investigated the processability of biaxial stretching for two kinds of LLDPE. The results show that PE-A is successfully stretched into a film with the draw ratio of 6 × 6 while PE-B fails. The molecular structure is characterized by the GPC, FTIR, SSA, and XS fraction. The condensed structure is investigated by SEM, WAXD, SAXS, and DSC. Moreover, a deformation model is proposed to illustrate why only PE-A can be biaxially stretched. First, the interlamellae 12428


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crossing adjacent lamellar stacks are observed in PE-A but not in PE-B, which increases the lateral tie connection. Second, the amount of fibril is less in PE-A than that in PE-B and the fibrillar structure lacks tie chains between each other. Third, the newly formed lamellae derived from the original fine lamellae can lock molecular chains between adjacent crystals, which increases the lateral connection via anchoring the tie chains from different lamellae. In addition, the relationship between molecular structure and condensed structure is also established.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.iecr.9b01733. Thermal procedures of SSA fractionation; curves of GPC, IR, and DSC; stress−strain curves of sequential biaxial stretching; process of obtaining Δq2; methods for determining the weight and volume fraction of interface according to the Raman and SAXS; method for getting Dfibril; and process of computing Bobs (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Fax: +86-28-8540-6333. Tel.: +86-28-8540-6333. ORCID

Qiang Chen: 0000-0003-0046-7340 Jian Kang: 0000-0002-3888-2462 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the National Natural Science Foundation of China (NSFC 51503134, 51721091) and the State Key Laboratory of Polymer Materials Engineering (Grant No. SKLPME 2017-3-02).

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REFERENCES

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Structure Evolution of Polyethylene in Sequential Biaxial Stretching

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