Structure–Property Evolution of Poly(ethylene terephthalate) Fibers in

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Structure−Property Evolution of Poly(ethylene terephthalate) Fibers in Industrialized Process under Complex Coupling of Stress and Temperature Field Jianping Ma,† Liang Yu,† Shichang Chen,*,† Wenxing Chen,*,‡ Yongjun Wang,† Shanshan Guang,† Xianming Zhang,‡ Wangyang Lu,‡ Ying Wang,† and Jianna Bao† †

College of Materials and Textiles and ‡National Engineering Laboratory for Textile Fiber Materials and Processing Technology (Zhejiang), Zhejiang Sci-Tech University, Hangzhou 310018, P. R. China

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S Supporting Information *

ABSTRACT: The structure−property evolution of commercial poly(ethylene terephthalate) (PET) fibers obtained from the different drawing and heat-setting stages in industrial processing was systematically investigated. Upon combination of thermal analysis (DSC and DMA) with crystallization and orientation (WAXD and SAXS), the variation of crystallization and microstructures mainly containing lamellar and microfibrillar crystals following the processes were discussed in connection with properties. Results indicate the significant tenacity increase of fiber in the drawing process is mainly attributed to the orientation development of the interlamellar amorphous region, the interfibrillar extension of amorphous molecular chains, and its entanglement with the lamellae. Accordingly, a decline of shrinkage can be seen as a fact of the coiling of amorphous molecular chains, the formation of rigid amorphous fraction, and the increase of crystallinity. Thus, a new four-phase model has been proposed to clarify the structure−property relationships of the commercial PET industrial fibers.

1. INTRODUCTION As one of the most versatile thermoplastic polymers, poly(ethylene terephthalate) (PET) has been widely applied in plastic, film, and fiber industries because of its high melting point, good mechanical properties, and sound chemical resistance.1 The mechanical properties of PET products are mainly dependent on the crystallization and orientation caused by complex coupling of stress and temperature field in processing. Generally, the structural variation of the industrialized PET is completed in tens of milliseconds because of the fast preforming; it is difficult to obtain intermediate samples and understand the intermediate state of polymer chain in products through the normal method. Although many structural models have been proposed for semicrystalline polymers,2−4 the structure−property evolution mechanism of commercial PET remains unclear. © XXXX American Chemical Society

Extensive studies of the structure−property relation of PET under the drawing and heat-setting processes have been performed over the past decades. Kawakami et al.5−8 conducted a systematic study of the structural evolution of completely amorphous PET sample of dog-bone-shaped during uniaxial drawing at different temperatures below 110 °C. No crystallization behavior was found when the microfibrillar structure appeared in the plastic deformation region. In the strain hardening zone, the crystallization was captured since the further nucleation and growth of crystallite facilitated the lamellar crystals into microfibrillar crystals with strain increasing, while the lamellar could fragment under large strain, Received: July 20, 2018 Revised: December 14, 2018

A

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by using small-angle X-ray scattering (SAXS). Comprehensive analysis of the thermal behavior, aggregation structure, and mechanical property of the fiber as well as the process− structure−property relationships was conducted, and a new four-phase model was proposed to understand the evolution mechanism of the structure−property relationship.

which led to further microfibrillar splitting and catastrophic breakage. Similarly, the strain-induced structural formation of completely amorphous PET film at 90 and 100 °C was discussed by Okada et al.9 It was noted that the new lamellar crystals appeared during drawing, and the isolated lamellae changed to stacked lamellae. As is well-known, the initial structure of PET molecular chains has a great influence on the structural change in processing with coupling stress and temperature field. It was particularly evident in the industrial processing of chemical fibers. Keum et al.10 performed a study of the amorphous PET fibers in the drawing of high-speed melt-spun upon elevating the temperature to 250 °C; a transient mesophase was found before the triclinic crystal structure formed. They also elucidated the microstructure’s variation of oriented chains in view of the effect of different spinning speeds on the thermal deformations.11 Furthermore, Haji et al.12 investigated crystallinity, molecular orientation, and another microstructure during the hot multistage drawing at 90−130 °C of the low oriented PET fibers. Murthy et al.13 investigated the structure changes of four types of PET fibers with high intrinsic viscosity and high tenacity. It seemed that tenacity of sample was largely determined by the diameter of lamellar stacks and the tilting angle of the lamellae. Recently, Liu et al.14,15 focused on the microstructure development of commercial PET industrial fibers during uniaxial deformation and thermal annealing at various temperatures in the instrument platform. Results indicated that the new lamellar crystals appeared and the isolated lamellae changed into stacked lamellae, and the annealing process at different temperatures had an impact on the crystallinity, thickness of amorphous region, and the long period of fibrillar crystals. However, most of the research for the structures evolution of PET is limited to a low deformation speed (as draw speed 4 mm/min14,15) and a long heat treatment time in in-situ laboratory conditions. It is noticeable that the draw rate can reach 7.2 m/min as described by Mahendrasingam et al.16−20 Nevertheless, it is far from the actual industrial processing draw rate of 3000 m/min. Although some online studies of the highspeed spinning process of PET had been reported by using insitu synchrotron X-ray diffraction experiments at take-up speeds up to 4200 m/min,21−23 it only indicated the changes of crystallization and orientation induced by the neck-like deformation during high-speed spinning, which had a significant difference compared with the multistage drawing and heat-setting deformation of PET industrial fibers.21−25 Consequently, the knowledge is still insufficient to bridge the gap of the structure−property relationship of PET. To the best of our knowledge, until now, there is no study of the structure−property evolution of PET fibers in an actual industrialized forming process under complex coupling of stress and temperature field. In the present paper, we focus on the deformation mechanisms of superlow shrinkage (SLS) type PET industrial fibers in the different forming stages of drawing and heatsetting process based on a special sampling method to acquire fiber samples during actual industry production. The crystallinity and motion ability of polymer molecular chain were investigated by using differential scanning calorimetry (DSC) and dynamic mechanical analysis (DMA). Twodimensional wide-angle X-ray diffraction (WAXD) was used to analyze the crystallization and orientation of PET fibers. Additionally, the lamellar and fibrillar structures were explored

2. EXPERIMENTAL SECTION 2.1. Materials. The PET industrial fibers (SLS) were obtained from Zhejiang Guxiandao Green Fiber Co., Ltd. The samples were taken from different forming stages of the industrialized process, which were marked as Y1, Y2, Y3, Y4, and Y5. The intrinsic viscosity ([η]) of the PET melt used for spinning in this paper was 0.98 dL/g, and the Tm was 256.12 °C. The drawing and heat-setting process flow is shown in Figure 1 (more detailed information can be found in the Supporting Information).

Figure 1. Drawing and heat-setting process flow of PET industrial fibers. It is necessary to note that the forming process of each sample was completed within 10−2 s. The corresponding critical process parameters are listed in Table 1. For the purpose of technology

Table 1. Processing Parameters in Different Forming Stages of Fiber Samples process parameters

sample

forming stage

Y1

undrawn, spinning speed 350−600 m/min first drawing second drawing and tension heatsetting first relax heatsetting second relax heatsetting

Y2 Y3 Y4 Y5

drawing ratio

drawing temp (°C)

2−5 1.2−3

100−150 200−250

relax ratio (%)

heatsetting temp (°C)

4−9

200−250

3−7

180−250

confidentiality, the accurate values of the process parameters are not displayed. Additionally, the general physical properties of fiber samples are also listed in Table 2. 2.2. Thermal Analysis. Thermal analysis of samples was performed by using a differential scanning calorimeter (DSC) (Mettler Toledo, Switzerland) under a nitrogen atmosphere. The samples were initially heated from 25 to 300 °C at 10 °C/min. Then, the crystalline mass fraction (XC,D) was calculated by the equation26 XC,D =

ΔHm ΔHm*

(1)

where ΔHm is the enthalpy of the melting of the PET fiber. The crystals melting enthalpy (ΔHm*) of PET was recorded as 117.6 J/g in most public works.14,27 Dynamical mechanical analysis (DMA) (TA, American) was conducted to obtain the storage modulus (E′), loss modulus (E″), and loss factor (tan δ), the temperature of maximum tan δ (tan δmax) B

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The crystalline orientation ( fc) was calculated as follows:30

Table 2. Mechanical Properties and Shrinkage of Fiber Samplesa sample

linear density (dtex)

tenacity (cN/ dtex)

elongation (%)

shrinkage (%)

initial modulus (cN/dtex)

Y1 Y2 Y3 Y4 Y5

34.00 8.12 5.86 6.31 6.64

0.84 3.75 7.82 7.06 6.72

629.01 31.66 23.11 28.68 31.36

24.0 18.0 4.1 2.9 1.9

6.40 35.44 46.50 44.47 43.65

fc =

Mechanical properties were tested by a Model YG021D electronic single-fiber strength tester. The shrinkage was measured in the oven under hot air at 177 °C for 15 min without load. All test results are the average data from 10 pcs filaments of the individual sample (Y1 to Y5).

(Tα), and the temperature range at the location of half-height width of internal friction peaks (TH/2) of samples, at a constant frequency of 1 Hz within the temperature range of 30−270 °C at a heating rate of 5 °C/min. 2.3. Sonic Measurement. Sonic measurement was performed on a digital fiber sound velocimeter (SCY-III, China) designed by Donghua University. The orientation of the amorphous phase was obtained according to the two-phase sonic modulus theory proposed by Samuels.28 2.4. Wide-Angle X-ray Diffraction (WAXD). A two-dimensional wide-angle X-ray diffractometer equipped with a D8 Discover diffraction instrument (Bruker, Germany) was used to obtain the crystallization and orientation indexes. PET fibers were cut into 30 mm in length, and then they was put on the sample platform horizontally. The test voltage and electricity were 40 kV and 40 mA, respectively. The X-ray wavelength (λ) was fixed at 0.15418 nm. The coupled 2θ/θ scanning was employed, and the scanning procedures were divided into three steps, with corresponding diffraction angle (2θ) being 20°, 40°, and 60°. The time for each scanning step was 70 s, so the total experiment time was 210 s. The crystallinity (XC,W) and crystallite size (D(hkl)) were calculated by the following equations:29,30

∑ Ic ∑ Ic + ∑ Ia

(2)

D(hkl) =

kλ β cos θ

(3)

(4)

where hi is the full width at half-maximum of crystal peaks in orientation (in degrees). 2.5. Small-Angle X-ray Scattering (SAXS). Experiments were performed using a SAXS/WAXS laboratory system (Xeuss, France) with a microfocus sealed X-ray tube (λ = 0.15418 nm) at 60 kV and 0.6 mA. The system was equipped with a Pilatus3 R200K detector, and the sample-to-detector distance was 0.45 m. The scattering experiments were conducted at room temperature. The fiber bundle was put in a capillary tube and horizontally mounted in sample stage, and the acquisition time for SAXS measurement was set to 1200 s. The effective scattering vector q = 4π sin(θ)/λ. The 2D images were treated with the software Fit2D and were integrated and converted to one-dimensional scattering data of scattering intensity I(q) parallel and perpendicular to the fiber axis, which are defined as q1 and q2, respectively.

a

XC,W =

360 − ∑ hi 360

3. RESULTS AND DISCUSSION 3.1. Thermal Properties. Thermal behaviors of PET fiber samples at different process stages were studied by using DSC and DMA. The heating curves of DSC and tan δ with change of the temperature are shown in Figures 3 and 4, respectively.

Figure 3. DSC thermograms of PET fibers in N2.

where Ic represents the crystallite integral intensity, Ia refers to amorphous integral intensity, k is the Scherrer constant (k = 0.89), and β is the full width at half-maximum of corresponding crystal plane. A PET fiber was cut into about 30 mm length and then placed on the specimen holder. The specimen holder was put vertically in the testing stage with vertical beam irradiating the sample fibers. Step (with count limit) scanning mode was adopted, and the efficient scan time was set to 210 s. The orientation measuring system is schematically shown in Figure 2. Figure 4. Loss factor of PET fiber in N2.

In Figure 3, the glass transition temperature (Tg) and cold crystallization temperature (Tc) of undrawn fiber (Y1) can be obviously observed. It is known that Tg of the fiber is usually inspected in the condition of spinning speed below 5000 m/ min.31 The unclear Tg and Tc from Y2 to Y5 indicate the existence of crystalline structure and high level of crystallinity. The thermal analysis data collected from DSC are listed in Table 3. It is noted that the XC,D of PET fibers increases gradually, and the melting peaks tend to be sharp and narrow from Y1 to Y5. It reveals that the crystalline structure of fibers

Figure 2. Schematic diagram of the WAXD orientation measuring system. C

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Y1

Y2

Y3

Y4

Y5

Tm (°C) Tm − Tm,onset (°C) ΔHm (J/g) XC,D (%)

254.22 15.81 −35.32 30.04

257.37 5.32 −50.60 43.03

261.90 3.11 −52.62 44.74

262.00 3.01 −56.23 47.82

260.00 2.58 −60.68 51.06

is becoming more and more perfect with the development of drawing and relax heat-setting processes. Thus, the difference between melting peaks (Tm) temperature and initial melt temperature (Tm,onset) (Table 3) is decreased from Y1 to Y5. It is worth mentioning that a minute melting peaks appears before the main melting peaks for samples from Y3 to Y5, which suggests some small imperfect crystals emerged in the process of crystallization improvement under heat-setting at high temperature.32,33 It would affect the crystal structure, and it is usually not mentioned in public works. When the amorphous structure of polymer is subjected to the couple effect of stress and temperature field, it causes the energy loss and exhibits viscoelasticity, which can be characterized by multiplicity and instability. In that way, the structural characteristics of amorphous region of fiber could be studied by the DMA method. Figure 4 shows the tan δ of PET fibers at different forming stages (the E′ and E″ curves of samples can be seen in Figure S1). Huisman et al.34 reported a mechanical response of the molecular mobility was based on the temperature where E″ reached a maximum level, i.e., transition temperature, and there is a good linear relationship between E″ and Tα. As listed in Table 4, The Tα increases from

Figure 5. (a) Two-dimensional WAXD crystalline diffractions. (b) One-dimensional intensity distribution profiles of PET fiber. (c) Deconvolution of the intensity distribution profile of Y5 as an example.

suggests no crystalline structure in undrawn fibers. It had been proved that was amorphous structure for an undrawn fiber when the spinning speed was located within 350−850 m/ min.36 With regard to Y2, the weaker diffractions have already been recognizable corresponding to three characteristic PET planes ((010), (−110), and (100)), which means the crystallization starts during the first drawing. By contrast, the more clear and complete diffraction halos from Y3 to Y5 imply a more perfect crystalline structure with high molecular orientation. This indicates the second drawing at high temperature has a significant impact on the molecular chains structure. A small but nonobvious difference among Y3, Y4, and Y5 shows the primary microstructure of PET fibers almost remains unchanged after two-stage relax heat-setting process with high temperature, but crystallization was gradually improved from one-dimensional WAXD intensity distribution profiles as shown in Figure 5b. The diffraction peaks information about PET fibers from Y1 to Y5 in Figure 5b is completely in conformity with the halo characteristics in Figure 5a. To obtain detailed crystalline structures, a curve fitting procedure was performed to separate the crystalline diffractions from amorphous scattering, for example, the treatment of Y5 as shown in Figure 5c. Detailed crystalline structural parameters are summarized in Table 5. The XC,W increases following the drawing and heat-setting process, which is consistent with the trend of XC,D (Table 3). It is noted that the XC,W is higher than XC,D (Figure S3). This discrepancy could be attributed to the role of rigid amorphous fraction (RAF).37 And it tends to be large in the second drawing and almost keeps same level in two-stage relax heatsetting processes. This indicates that the RAF of interlamellar amorphous increases to different extents in whole processes. Herein, the densities (ρ) of the samples were obtained using a solid density instrument, which was different from the traditional density gradient methods. As seen in Table 5, the variation trend of ρ of fiber samples in different forming stages is consistent with that of XC,W.

Table 4. DMA Data of PET Fibers parameters

Y1

Y2

Y3

Y4

Y5

Tα (°C) tan δmax TH/2 (°C)

96.7 2.00 10.68

147.9 0.27 48.12

155.2 0.20 43.75

147.2 0.20 42.19

145.2 0.19 40.50

Y1 to Y3 but decreases gradually from Y3 to Y5. The Tα of the fiber can be considered to be coincided with the Tg.35 It is obvious the undrawn fiber (Y1) has the best molecular mobility. With the development of the two-stage drawing processes, the molecular chains experience palpable orientation along the fiber axis, the large molecular chains tend to be packed closely, and then a small amount of crystals appears. After that, the extended molecules in amorphous orientation regions have sufficient mobility to obtain their most probable configuration in the two-stage relax heat-setting processes, which means these molecules are likely to coil up, indicating a better molecular chains mobility from Y3 to Y5. With regard to tan δ, there is a positive relation between the tan δmax and the content of amorphous region. Consequently, the amorphous region reduces dramatically in drawing stages but slightly in relax heat-setting stages. This is consistent with the variation of tan δmax in Table 4. Samui et al.4 found that a lower tan δmax can lead to more tie molecules. In addition, the constant decrease of TH/2 from Y2 to Y5 indicates the crystallization structure of fiber becomes perfected gradually. 3.2. Crystallization and Orientation Analysis. The crystallization and orientation results of PET fibers obtained from two-dimensional WAXD are vividly depicted in Figures 5 and 6, respectively. An annular halo pattern of Y1 (Figure 5a) D

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Figure 6. (a) Two-dimensional patterns of PET fiber. (b) One-dimensional intensity distribution profiles obtained from WAXD. (c) Deconvolution of the intensity distribution profile of Y5 as an example.

shown in Figure 6b (the detailed integrate signal processing is illustrated in Figure S2). It was said that the drawing process could promote the amorphous molecular orientation whereas heat-setting diminished the orientation.15 As presented in Table 6, the sonic orientation parameter (fs) increases from Y1 to Y3 (in drawing) but decreases from Y3 to Y5 (in relax heat-setting). The sonic orientation parameter (fs) of Y3 is slightly higher than that of Y2; the amorphous orientation factor (fa) is opposite. Noticeably, the amorphous molecular orientation in drawing at high temperature (from Y2 to Y3) is actually determined by the competition between drawing and tension heat-setting process. It is reasonable to conclude that the coiling of amorphous molecules induced by tension heatsetting takes a competitive advantage compared with the extending of amorphous molecules caused by drawing. 3.3. Structure Evolution of PET Fibers. The typical twodimensional SAXS patterns of PET fibers are shown in Figure 7. There are two kinds of diffuse scattering features that one is near the beam spot and the other one is the symmetrical lamellar peaks on both sides of the beam stop. Obviously, the latter will be the focus of research. The lamellar peaks in the meridian direction (q1) at both sides of the beamstop were caused by periodic electron density distribution in the oriented lamellar structure.38 It reveals that the lamellar stacks structure exist in the fiber through drawing. The lamellar peaks often appeared as a bar, which implied a coexistence of two-spot and four-spot patterns.13 When the lamellar surface tilted away from the chain axis, a four-spot pattern emerges.39 Because the two lamellar peaks from Y2 to Y5 all have four scattering maxima shown in Figure 7, a fourspot pattern always exists in all forming stages of drawing and heat-setting process of the PET industrial fibers. Apparently, the pictorial images are deservedly well in agreement with WAXD graphics (Figures 5 and 6), indicating the same crystalline performance. The corresponding scattering intensity scans of lamellar peaks along and perpendicular to the fiber axis from Y2 to Y5 are illustrated as Figure 8 (q1) and Figure 10 (q2), respectively (Y1 is not discussed because of no lamellar peaks). The scattering intensity has a sudden increase from Y2 to Y3 during the second drawing at high temperature, indicating the

Table 5. Crystalline Structural Parameters of PET Fibers crystallite size (Å) sample

XC,W (%)

ρ (g/cm3)

D(010)

D(−110)

D(100)

Vc (nm3)

N (1018/cm3)

Y2 Y3 Y4 Y5

46.6 68 68.5 74.4

1.336 1.342 1.345 1.365

30.8 59.9 57.7 56.4

38.3 48.9 45.6 45.8

28.6 44.3 43.2 43.6

918 1274 1291 1344

1.23 0.44 0.48 0.46

After first drawing, the crystallite size of Y2 is the smallest, while its crystal number (N) is the most, since the orientation induces crystallization at low temperature. It is easy to understand that the drawing process at low temperature only results in a large number of small crystals, while the crystallite size of Y3 (after the second drawing at a high temperature) increases remarkably and the numbers of crystals decreased significantly. Interestingly, the crystallite size of Y4 and Y5 decreases, but the volume of crystal (Vc) increases in the twostage relax heat-setting processes. This suggests that the crystal grows at different directions in relax heat-setting process compared with in drawing process. From Y3 to Y5, the value of N almost remains unchanged, suggesting the crystallinity increase is mainly caused by the growth of crystals. Similar to the crystalline diffraction images (Figure 5a), no diffraction ring can be seen for the undrawn Y1 in Figure 6a that there is no crystalline orientation; the variation of definition and brightness of signal from Y1 to Y5 is in agreement with the diffraction crystalline images. Correspondingly, the orientation in the crystalline region (fc) continues to increase as listed in Table 6. DIFFRAC.EVA software was used to integrate the pattern from 0° to 360° to obtain the onedimensional profiles through the circular integral method, as Table 6. Orientation Parameters of PET Fibers samples

fs

fa

fc

Y1 Y2 Y3 Y4 Y5

0.711 0.847 0.850 0.800 0.780

0.711 0.823 0.782 0.571 0.383

0.878 0.885 0.905 0.916 E

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Figure 7. Two-dimensional SAXS patterns of PET fibers at different forming stages.

Figure 8. (a) One-dimensional SAXS profiles along the meridian direction of PET fiber; the inset in (a) shows the integrating direction using Fit 2D in the analysis. (b) Electron density correlation function that the effective long period (L*) and amorphous thickness (La) can be derived as shown in the inset of (b); then crystalline thickness (Lc) can be calculated based on the relation Lc = L* − La.

(L*) tends to be large under complex coupling of stress and temperature field. The linear crystallinity (XC,L) can be calculated as the ratio of Lc to L* (the data of L, L*, Lc, La, and XC,L from Y2 to Y5 are listed in Table S1). It is helpful to understand the distribution of crystalline lamellae and amorphous region in the fibril.32,40 Figure 9 shows the long period (L, L*), crystalline thickness (Lc), and amorphous thickness (La) of PET fiber from Y2 to

difference of electron density between the amorphous and crystalline phase increases rapidly. This has been verified by the sharp increase of XC,W (Table 5) and the decrease of fa (Table 6). By comparing the scattering intensity between Y3 and Y4, a slight ascend of scattering intensity can be seen after the first relax heat-setting process, suggesting the decrease of electron density ratio between crystalline and amorphous layers along the fiber axis. The theoretical long period (L), which represents the repeat unit distance of crystalline and amorphous phase in the fiber structure, can be calculated by the peak position (q1,max) of I(q1) according to Bragg’s law:32 L=

2π q1,max

(5)

From Y2 to Y5, q1,max shifts to a smaller value (Figure 8a) so that the long period increased gradually in the drawing and heat-setting process. To understand the thickness of amorphous and crystalline phase along the meridian direction, the one-dimensional electron density correlation function γ(z) is adopted to analyze the scattering profiles:32 Figure 9. Variation of long period, crystalline thickness, and amorphous thickness following the forming stages.

∞

γ(z ) =

∫0 I(q1) cos(q1z) dq1 ∞

∫0 I(q1) dq1

(6)

Y5. The L is larger than L*; it is in agreement with a previous study.41 It is noted that L and L* follow the same trend as the laminar peak intensity in Figure 7. In the following discussion, L* will be adopted to substitute for L because both Lc and La are derived from the correlation function. In the above bars charts, L* and Lc increases gradually along with the drawing and heat-setting process, whereas La increases with the drawing process (from Y2 to Y3) but decreased after entering the heat-setting process (from Y3 to Y5). Even

where z denotes the fiber axis. The correlation function was obtained via Fourier transformation of the SAXS profile. The Lorentz correction where I(q1) was multiplied by q12 was not applied because of the anisotropic orientation of the lamellar in fibers.3,33 In Figure 8b, it can be seen from the correlation functions that the first peak moves toward a larger z following the industrial processes, indicating that the effective long period F

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Figure 10. (a) Scattering intensity of the lamellar peaks along q2. (b) Curve-fitting procedure used for evaluation of lamellar. (c) The lateral size of lamellar. (d) The tilting angle of lamellar surface.

where Δq2 is defined as the full width at half-maximum (fwhm) of the fitting Pearson VII function, Δχ represents the separation distance between the centers of the two Pearson VII functions, and q1,max is the position of lamellar peak along q1 at the maximum intensity of I(q1). The variation of lateral size of lamellar (DSAXS) and tilting angle of lamellar surface (Φ) from Y2 to Y5 are shown in Figures 10c and 10d, respectively; the detailed data are also listed in Table S1. With the development of drawing and heatsetting, the DSAXS decreases whereas the Φ first increases in drawing and then decreases in relax heat-setting. The remarkable changes of both the DSAXS and the Φ can be seen in the second drawing process, indicating the major variation of crystalline and lamellar structures occurs from Y2 to Y3 again. In addition, the variation of Φ in the whole drawing and heat-setting process suggests the lamellar surface always keeps a constant angle with the fiber axis. This indicates that a four-spot always exists in the forming stages of PET industrial fiber, which is consistent with the results of four scattering maxima in Figure 7. Up to this point, the microstructure of PET fiber and its relationship with processing are gradually being understood; all that remained is trying to clarify the relation of structure and property. The evolution of minute structure and the corresponding tenacity and shrinkage from Y1 to Y5 are presented in Figure 11. It is evident that the declared twophase models of crystalline phase and amorphous phase are no longer effective to explain the variations of tenacity and shrinkage from Y2 to Y3. According to three-phase models (the ordered crystalline lamellae, interlamellar amorphous region, and interfibrillar extended noncrystalline molecules),37,42 the third phase of Y3 was speculated to be further extended and entangled with adjacent microfibrils, which led to tenacity increase. Meanwhile, the extended molecules in interlamellar amorphous might be coiled sharply. Nevertheless,

though Y1 is not presented because of its noncrystalline structure, it is still worth considering that the process from Y1 to Y2 refers to the first drawing with the high draw ratio and low temperature, implying the formation of the initial lamellar structure configuration. Both Lc and La are significantly increased from Y1 to Y2 due to the extension of the amorphous phase and the strain-induced crystallization. On account of the second drawing process with high temperature, Lc and La of Y3 were further increased. As mentioned before, the high temperature and strain induction led to the recombination and enhancement for many small crystalline lamella and the interlamellar amorphous region. In addition, the competition between extending and coiling of the molecular chains in interlamellar amorphous region occurred during the second drawing stage, which the molecular chain coiling is a notch above. When entering into the two-stage relax heat-setting process (Y3 to Y5), L* and Lc increased while the corresponding La decreased gradually. As we know, the structural changes in relax heat-setting process mainly include the crystallization of coiled molecules in the amorphous phase and the growth of crystal lamellar instead of the molecular chains extending. Figure 10 reveals the scattering results of lamellar peaks along q2 of PET fiber. It is clearly observed that the scattering intensity in Figure 10a is a basic one-to one correspondence with the two-dimensional SAXS patterns in Figure 7. Here, two parameters, the lamellar lateral size (DSAXS) and the tilting angle between lamellae surface and fiber axis (Φ), are provided to illustrate the lamellae structure and calculated according to the equations32,33 DSAXS =

2π Δq2

Φ = tan−1

Δχ 2q1,max

(7)

(8) G

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tenacity continues to increase, while the shrinkage decreases keenly. In a macroscopic sense, the property changes of Y3 should be attributed to the second drawing at a higher temperature. The drawing process leads to an increase of amorphous orientation; on the contrary, the heat-setting process induces the reduce of amorphous orientation. From the slight decrease of fa listed in Table 6, it can be known that the coiling of molecules in interlamellar amorphous has a competitive advantage over the extension of noncrystalline molecules from Y2 to Y3. Thus, the maximum reduction of shrinkage from Y2 to Y3 should be attributed to the coiling of interlamellar amorphous molecules, the greater RAF formation, and the higher crystallinity. After the drawing process, the perfection of crystalline lamellae structure is increasingly improved, and the amorphous phase orientation is steadily decreased during the following two-stage relax heat-setting processes, which cause decrease of shrinkage and tenacity. Besides, the slight increase of RAF content in relax heat-setting stages results in the shrinkage decrease. In the first relax heatsetting process (from Y3 to Y4), the fa decreased dramatically, which is mainly due to the coiling of interlamellar and intermicrofibrillar amorphous molecular chains. In the second relax heat-setting process, the structure of Y5 displays a similar evolution with Y4 as that the fa decreases substantially, suggesting a further decrease of tenacity and shrinkage. According to the above process−structure−property relationship, the theoretical guidance can be proposed to understand the regulation of structures and desired properties of polymer fibers. In the multistage drawing and heat-setting process, the first drawing is more important to increase tenacity. The draw ratio should be as large as possible, but for the temperature it is completely opposite, so that a higher molecular orientation with a relatively lower crystallinity can be obtained, which makes contribution to the more increase of tenacity in the further drawing process. The thermal shrinkage is mainly dependent on the second drawing process in which the drawing temperature needs to be set as high as possible. In that way, a high setting temperature and a high relax ratio in the multistage heat-setting process are the prerequisites to obtain superlow shrinkage (SLS) fibers.

Figure 11. (a) Three-dimensional structural models from Y1 to Y5 in multistage drawing and heat-setting process of PET industrial fiber. In the model of Y5, A is the tilting fold surface of crystalline lamellae, B is the crystalline lamellae (Lc represents its thickness), C is the interlamellar amorphous (La represents its thickness) which contain the rigid amorphous fraction (RAF) marked in green color and mobile amorphous fraction (MAF), D is the extended noncrystalline molecules, and L* denotes the long period (L* = Lc + La). (b) The pliers-shaped evolution for the properties of Y1 to Y5 corresponding to the structure evolution of PET fiber.

it still suffers something difficult in the explanation of RAF formation. Based on the traditional “three-phase” model, a new “fourphase” structural model is proposed in the present study. Figure 11a shows the visually clear and compendious threedimensional structural models of the formation stage during drawing and heat-setting. The four-phase model includes crystalline lamellae, interlamellar rigid amorphous fraction (RAF, partial enlarged view in Figure 11a), mobile amorphous fraction (MAF), and the interfibrillar extended noncrystalline molecules. The packed crystalline lamellae is simplified as cuboid. RAF and MAF constitute the interlamellar amorphous region. The extended noncrystalline molecules located between the microfibrils consisting of crystalline lamellae and interlamellar amorphous region. In addition, the length, width, and height of the simplified cuboid are estimated as the average of the three equator crystallite sizes ((D010 + D−110 + D100)/3), the lamellar lateral size (DSAXS), and the crystalline thickness (Lc), respectively. In Figure 11b, the cross curves of tenacity and shrinkage can be termed a pliers-shaped evolution of properties. Previous studies proved that the tenacity and shrinkage are highly governed by the amorphous phase.34,40,43,44 A low tenacity and high shrinkage of undrawn fiber (Y1) are attributed to the lower amorphous orientation and absence of lamellar stacks. After first drawing with the large draw ratio and low temperature, the microfibrils composed of the repeated long period (L*) appear in Y2. Here, the remarkable increase of tenacity might be attributed to the higher amorphous orientation of the interlamellar amorphous and extended noncrystalline molecules, while the shrinkage decreasing from 24% to 18% (Table 2) is mainly due to the formation of a large number of small crystal nuclei. This provides another proof that the crystals formation can inhibit thermal shrinkage effectively. In addition, the RAF can hardly be seen in Y2 as the difference of crystallinity obtained from WAXD (XC,W) and DSC (XC,D) is very minute (Figure S3). From Y2 to Y3, the

4. CONCLUSIONS Inspired by online sampling, commercial PET industrial fibers collected from different forming stages of drawing and heatsetting process were analyzed to provide an insight of process− structure−property relationship. Structure−property evolution of PET fibers, based on the actual industrial processes under complex coupling of stress and temperature field, was investigated by DSC, DMA, WAXD, and SAXS. With the development of drawing and the heat-setting process, the crystallinity of PET fibers increased gradually, and the crystalline structures became more and more perfect. The impact of drawing on these was even greater than relax heatsetting. This could be confirmed by any of the three methods (DSC, WAXD, and SAXS). Results of WAXD diffraction peak showed significant changes in crystalline structure in drawing process, and it almost was maintained in the following relax heat-setting processes. The crystalline orientation factor nearly increased as a constant in each processing stage. However, it was complex that the amorphous orientation factor increased during the first drawing and then decreased slightly during the second drawing and decreased significantly during the twoH

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Macromolecules

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stage relax heat-setting process. It seemed that a high temperature led to a decline of amorphous orientation. Similarly, the long period and crystalline thickness increased gradually along with the drawing and heat-setting process, whereas the amorphous thickness increased during the second drawing stage and then decreased in the two-stage relax heatsetting processes. In different processing stages, the mechanical performance referred to different evolution mechanism of microstructure of PET fibers; generally, it could be concluded to be closely dependent on the amorphous phase structures. Consequently, a new “four-phase” structural model mainly involving the added interlamellar rigid amorphous fraction was proposed to explain the process−structure−property relationship of the typical industrialized fiber. This study for the first time proposes a comprehensive explanation which can be used to fully describe the process−structure−property relationships of the typical industrialized fiber and film under the complex coupling of stress and temperature field.

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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.macromol.8b01561.

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Industrial process and sampling points; storage modulus and loss modulus of PET fibers; orientation measuring method, comparison of crystallinity obtained from different methods; structural parameters of fiber samples obtained by SAXS (PDF)

AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected] (S.C.). *E-mail: [email protected] (W.C.). ORCID

Shichang Chen: 0000-0002-3974-6913 Wenxing Chen: 0000-0002-4554-1455 Wangyang Lu: 0000-0002-5650-0675 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was sponsored by the National Key Research and Development Program of China (No. 2016YFB0303000), National Natural Science Foundation of China (No. 51803187), and Zhejiang Provincial Natural Science Foundation of China (No. LQ18E030011).

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DOI: 10.1021/acs.macromol.8b01561 Macromolecules XXXX, XXX, XXX−XXX