Polymorphic Crystal Transition and Lamellae Structural Evolution of

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Polymorphic Crystal Transition and Lamellae Structural Evolution of Poly(p-dioxanone) Induced by Annealing and Stretching Ying Zheng, Jian Zhou, Yongzhong Bao, Guorong Shan, and Pengju Pan J. Phys. Chem. B, Just Accepted Manuscript • DOI: 10.1021/acs.jpcb.8b12111 • Publication Date (Web): 11 Apr 2019 Downloaded from http://pubs.acs.org on April 12, 2019

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The Journal of Physical Chemistry

Polymorphic Crystal Transition and Lamellae Structural Evolution of Poly(p-dioxanone) Induced by Annealing and Stretching

Ying Zheng, Jian Zhou, Yongzhong Bao, Guorong Shan, Pengju Pan*

State Key Laboratory of Chemical Engineering, College of Biological and Chemical Engineering, Zhejiang University, 38 Zheda Road, Hangzhou 310027, China

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The Journal of Physical Chemistry 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

Abstract: Semicrystalline polymers usually undergo multilevel microstructural evolutions under high-temperature annealing and stretching deformation; this is essential to tailor the physical properties of polymer products in industrial processing. Here, we choose poly(p-dioxanone) (PPDO), a typical biodegradable, biocompatible, and bioresorbable polymer, as a model semicrystalline polymer and investigated its polymorphic structural transition and crystalline lamellar evolution under hightemperature annealing and stretching. High-temperature annealing caused the -to- phase transition of PPDO, accompanied by the improvement of crystallinity (Xc) and thickening of crystalline lamellae. Tensile strength and Young’s modulus of PPDO increased but the breaking strain decreased as the annealing temperature increased. Stretch-induced phase transition of PPDO depended strongly on the initial structure and stretching temperature (Ts). The -form PPDO transformed into its  counterpart during stretching at low Ts; this phase transition was irreversible and did not retain to  form with the release of stress. However, no phase transition took place for the -form PPDO stretched at high Ts ( 40 C). Original lamellae of -form PPDO changed into the fibrillar lamellae during stretching via melt-recrystallization mechanism.

INTRODUCTION Poly(p-dioxanone) (PPDO) is a typical semicrystalline polymer with good biodegradability, biocompatibility, bioabsorbability, and mechanical property. PPDO has been widely used in biomedical fields such as surgical suture, drug delivery system, bone and tissue fixation devices.1,2 Previous report has demonstrated that PPDO can form two crystal modifications ( and  form) with changing the crystallization temperature (Tc).3 The ′ form is a metastable phase while the  form is


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thermodynamically stable. The  form is characterized as an orthorhombic unit cell with a P212121 space group and the lattice parameters of a = 0.970 nm, b = 0.742 nm, and c (chain axis) = 0.682 nm.4 In addition, physical properties of PPDO are influenced by its crystal modification. The -form PPDO possesses slower hydrolytic degradation, higher flexibility, but lower strength and modulus than its  counterpart.3 It is well known that the crystal structure of polymorphic polymers strongly influences their physical properties such as the thermal, mechanical, barrier, and degradation properties. Understanding the relationships between crystal structure and processing condition is of fundamental importance for tuning the physical performances of final materials in processing. On the other hand, polymorphic polymers generally undergo crystal transition during thermal treatment and deformation. Metastable polymorph can transform into the thermodynamically stable one during thermal treatment, which has been demonstrated for a variety of polymorphic polymers such as isotactic polypropylene (iPP),5,6 poly(1-butene) (PB1),7,8 and poly(L-lactide) (PLLA).9,10 Thermal-induced phase transition is usually accompanied by variation of physical properties of polymer materials.10-12 Besides thermal treatment, stretching deformation can also cause the polymorphic crystal transition of semicrystalline polymers.13-19 Two types of crystal transitions have been identified in semicrystalline polymers during stretching, according to their reversibility. Most of the stretch-induced phase transitions are irreversible in the semicrystalline polymers such as iPP,20,21 PB-1,16 and poly(vinylidene fluoride);22 the crystal modification generated under stretching cannot retain to the original modification after releasing the stress. For the case of reversible phase transition, crystalline polymorph obtained under stretching can recovery into the original one with releasing the stress, which has been reported in several semicrystalline polymers including poly(ethylene 3 ACS Paragon Plus Environment


oxide),23

poly(butylene

terephthalate),24

poly(butylene

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succinate),25,26

poly(tetramethylene naphthalate),27 and poly(ethylene succinate).28 Stretch-induced crystal transition of semicrystalline polymers is usually accompanied by the changes of lamellae morphology. During stretching, semicrystalline polymers transform from an isotropic crystal structure into highly oriented fibrillar structure with the chain preferentially aligned along the stretching direction. Two processes have been proposed to explain the structural evolution mechanism of semicrystalline polymers under stretching. One is the inter- and innerlamellar slip;29,30 the other one is stress-induced melt and recrystallization.31,32 Recently, it has been demonstrated that the block slippage occurs first under stretching, followed by the stress-induced fragmentation and recrystallization at high strain.33,34 A major application of PPDO is used as the surgical suture and the surgical suture is generally prepared through melt spinning, which involves the orientation and stretching of polymer chains. Therefore, it is of fundamental importance to investigate the multilevel structural transition of PPDO during stretching, in order to tailor its physical properties in processing. In spite of recent progresses on the crystallization kinetics, crystalline structure and physical properties of PPDO,3,4,35,36 the polymorphic crystal transition of PPDO during stretching and annealing is not yet well understood. In this work, we systematically studied the polymorphic crystal transition and crystalline lamellar evolution of PPDO with annealing and stretching under different conditions. It is found that the phase transition between  and  crystals of PPDO takes place during annealing and stretching. High-temperature annealing favors the -to- phase transition; however, low-temperature stretching induces the irreversible -to- phase transition of PPDO. The underlying mechanisms for annealing and stretch-induced 4 ACS Paragon Plus Environment

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polymorphic phase transitions and crystalline lamellar evolution of PPDO were also discussed.

EXPERIMENTAL Materials. PPDO with a weight-average molecular weight of 66 kg/mol and a polydispersity of 2.2 was obtained from Tianjin Dongnan Hengsheng Medical Technology Limited Co. (China). Preparation of Annealed PPDO. To prepare the -form PPDO, PPDO was hotpressed at 140 C after melting for 3 min and then quenched to 5 C for crystallization for 12 h.3 The sample was then annealed in an oven preset to the annealing temperature (Ta, 40−90 °C) for 20 min. Preparation of Stretched PPDO. The crystallized PPDO film was cut into dumbbell specimen with a length of 35 mm, cross-section width of 3 mm, and thickness of ~0.4 mm. The specimen was stretched on a Zwick/Roell Z020 instrument at various stretching temperatures (Ts’s) under a strain rate of 5 mm/min. After stretching, the stress was released and the sample contracted to some extent in the stretching direction. True (or Hencky) stress (H) and true strain (H) of stretched sample were evaluated according to a published method.15,37 Briefly, ink points were marked on the surface of specimen; the displacements of ink points at the measured region for differential scanning calorimetry (DSC), wide-angle X-ray diffraction (WAXD), smallangle X-ray scattering (SAXS) and Fourier transform infrared (FTIR) spectroscopy were recorded by a digital camera. Width, thickness, and marked distance of the initial and stretched specimens were denoted as (W0, T0, b0) and (W, T, b), respectively. It was assumed that the volume of specimen did not change during stretching, since no whitening was observed. Therefore, W0T0b0 equaled to WTb during stretching. H and H were calculated as follows 5 ACS Paragon Plus Environment


b   b0 

 H =ln 

H =

Fb W0T0b0

Measurements. DSC Measurements. Thermal behavior of PPDO was measured on a Netzsch 214 Polyma DSC (Netzsch, Germany) equipped with an IC70 intracooler under the nitrogen gas flow (40 mL/min). The annealed or stretched sample (8−10 mg) was heated from 20 to 140 °C at a heating rate of 10 °C/min. Degree of crystallinity (Xc) was calculated from the DSC results by Xc,DSC = ΔHm/ΔHm0 × 100%, where ΔHm was the measured melting enthalpy of PPDO and ΔHm0 was the ΔHm of PPDO crystals with the infinite large lamellar thickness (ΔHm0 = 141 J/g).38 It is assumed that the  and -form PPDOs have the similar ΔHm0; such assumption has been usually used for the different polymorphs of semicrystalline polymers.39-42 WAXD and SAXS Measurements. WAXD and SAXS analyses were performed on the beamline BL16B1 of Shanghai Synchrotron Radiation Facility (SSRF). Wavelength of radiation source was 0.124 nm. WAXD and SAXS patterns were collected using a Rayonix SX-165 CCD detector (Rayonix, Illinois, USA). The sampleto-detector distances were 2.0 and 0.14 m for SAXS and WAXD measurements, respectively. For the temperature-variable WAXD analysis, the sample was sandwiched by polyimide films and heated from 35 to 120 °C at 10 C/min on a Linkam TST350 hot stage, during which WAXD patterns were recorded every 5 °C. For the in-situ WAXD and SAXS measurements of stretched sample, the dumbbell specimens, cut from the crystallized PPDO films, were stretched with a uniaxial stretching velocity of 8.3 μm/s on a Linkam TST350 hot stage with a clamping distance of 15 mm. Collection times of each WAXD or SAXS pattern were both 15 s. All SAXS data were corrected from the background, air scattering, and sample absorption. Two-dimensional (2D) WAXD and SAXS patterns were converted to the one-dimensional (1D) data by


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integration via a Fit2D software. 1D-WAXD curves were integrated from the 0−180° region of 2D-WAXD patterns; 1D-SAXS curves were integrated within ±20° along the stretching direction of 2D-SAXS patterns. Orientation degree of crystals was evaluated by the orientation parameter, fhkl, proposed by Hermans

f hkl 

3 cos 2  hkl   1 2

where φ is the angle between the normal direction of hkl crystal plane and the reference axis (stretching direction).  cos 2 hkl  is attained from the azimuthal diffraction intensity distribution by  cos hkl 2

 

 /2

0

I hkl   cos 2  sin  d



 /2

0

I hkl   sin  d

where Ihkl(φ) is diffraction intensity along the azimuthal angle φ. fhkl equals to −0.5 when all crystals are perfectly oriented with their normal perpendicular to the stretching direction and 0 when the crystals are randomly orientated. Xc was also calculated from the WAXD patterns

X c,WAXD 

Acr Acr  Aam

where Xc,WAXD, Aam, and Acr represent the crystallinity evaluated from WAXD, the areas of amorphous and crystalline peaks (2θ = 10−25°), respectively. Aam and Acr were evaluated by the curve splitting/fitting method, as illustrated in Figure S1.43 FTIR Spectroscopy. FTIR spectra were measured on a NICOLET iS50 FTIR spectrometer (Thermo Scientific, USA) at ambient temperature. All of the spectra were collected with 32 scans and a resolution of 4 cm−1. Uniaxial Tensile Test. Tensile tests were performed on a SUNS UTM2503



instrument under a strain rate of 20 mm/min at ambient temperature. The dumbbell specimen with a length of 35 mm, cross-section width of 3.0 mm, and thickness of ~0.4 mm was cut from the crystallized and annealed PPDO films. Yield stress (σy), Young’s modulus (E), and breaking strain (εb) were calculated from the engineering stress-strain curves. σy and εb corresponded to the stress at yield point and the strain at fracture, respectively; E was calculated from the slope of stress-strain curves in the elastic region (strain range: 2−4%). Seven replicated measurements were performed for each sample and the averaged results were used.

RESULTS AND DISCUSSION Annealing Effects on Thermal Behavior. Effects of annealing temperature (Ta) on melting behavior, crystalline structure, and mechanical properties of PPDOs were investigated. Figure 1a,b show the DSC heating curves and Xc,DSC of -form PPDOs annealed at various Ta’s for 20 min. For the unannealed -form PPDO crystallized at 5 C, an exotherm (Pexo, indicated by arrow) was observed prior to the major melting peak (Pm), which was arisen from the -to- transition upon heating.3 Two endotherms, Pm and annealing peak (Pa), were observed in the melting curves of PPDO annealed at different Ta’s. The small endotherm observed prior to the Pm was characteristic of Pa, indicated by asterisk in Figure 1a. The temperature of Pa was always ~10 C higher than Ta. The Pa has been detected in many semicrystalline polymers44,45 and attributed to the melt of polymers generated in secondary crystallization during long time annealing. The temperature of Pm changed little with increasing Ta. As shown in Figure 1b, Xc,DSC of PPDO improved notably after annealing or with increasing Ta. The initial -form PPDO had a low Xc,DSC of 46.5%, which increased to 60.9% after annealing at 90 °C for 20 min. The increase of Xc with Ta was attributable to the crystalline perfection and lamellar thickening during high-temperature annealing. 8 ACS Paragon Plus Environment

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b

a 6

Pm



5 

4

70



1



60

Pa 

50



40

Pexo

5

2

0

40

55 50 unannealed

unannealed

0

65 60

80



3

Ta(C) 90

Xc ()

Heat flow (mW/mg) Endo up

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


45

80

Temperature (C)

120

0

20

40

60

80

100

Ta (C)

Figure 1. DSC results of -form PPDOs after annealing at different Ta’s for 20 min: (a) DSC heating curves; (b) crystallinity. Annealing-Induced -to- Phase Transition. Polymorphic structural changes of PPDO during annealing were investigated via WAXD and FTIR spectroscopy. Figure 2a shows the WAXD patterns of -form PPDOs (Tc = 5 °C) after annealing at various Ta’s for 20 min. The -form PPDO has three characteristic diffraction peaks at 17.8°, 19.2° and 23.4° (λ = 0.124 nm), corresponding to the diffractions of (210), (020) and (310) planes, respectively.3,4 Only a diffraction peak was observed at 2 = 17.8° and the diffraction peak at 2 = 19.2° was not seen for the -form PPDO crystallized at 5 °C.3 For the -form PPDO annealed at a low Ta (50 °C), their WAXD patterns changed little compared to the unannealed sample, suggesting the absence of crystalline phase transition. After annealing -form PPDO at a high Ta (> 50 °C), the diffraction peak at 2 = 17.8° varied little but the diffraction peak at 2 = 19.2° became more pronounced and shifted to larger angle. Three diffraction peaks, characteristic of the (210), (020) and (310) planes of  form,3,4 were clearly seen for the -form PPDO annealed at 90 °C, demonstrating the occurrence of -to- phase transition during annealing at high Ta ( 50 °C).


310

10

15

20

2  Å

Ta(C) 90 80 70 60 50 40 5

2880

2957

2979

ABS (a.u.)

020

unannealed

2988

b

210

a Intensity (a.u.)

unannealed

25

Ta(C) 90 80 70 60 50 40 5

3050 3000 2950 2900 2850 2800

Wavenumber (cm-1)

1748

1734

c ABS (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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2926


Ta(C) 90 80 70 60 50 40 5

unannealed

1780 1760 1740 1720 1700

Wavenumber (cm-1)

Figure 2. WAXD and FTIR spectra of -form PPDOs after annealing at different Ta’s for 20 min: (a) WAXD patterns; (b) FTIR spectra in 3050−2800 cm−1; (c) FTIR spectra in 1790−1700 cm−1. Figure 2b,c show the FTIR spectra of -form PPDOs annealed at different Ta’s for 20 min. Notable differences in FTIR spectra were observed among the samples annealed at different Ta’s (40−90 °C), especially in the wavenumber ranges of 3050−2800 [ν(CH2) band] and 1790−1700 cm−1 [ν(C=O) band]. First, no discernible spectral splitting was observed for the -form PPDO (crystallized at 5 °C) unannealed or annealed at low Ta (40, 50 °C) in the wavenumber ranges of 3050–2800 cm−1 [ν(CH2) band] (Figure 2b). However, splitting bands at 2988 and 2979 cm−1 were seen for the -form PPDO annealed at high Ta (60−90 °C). These splitting bands are characteristic 10 ACS Paragon Plus Environment

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of the -form PPDO, indicating the relatively stronger interchain interactions within the crystal lattice.3,9,46,47 Similarly, the -form PPDO crystallized at 5 °C showed a strong absorption at 1748 cm−1 and a small shoulder at 1734 cm−1 in the wavenumber ranges of 1790−1700 cm−1 [ν(C=O) band] (Figure 2c). However, the shoulder band of 1734 cm−1 became more pronounced as Ta increased. These WAXD and FTIR results demonstrated that the ′-to- phase transition occurred after annealing at high Ta (> 50 °C) and the phase transition was facilitated by increasing Ta. The -form PPDO completely transformed into its  counterpart after annealing at 80 and 90 °C for 20 min. The annealing-induced ′-to- phase transition was accompanied by the enhancements of interchain interaction and chain packing order in the crystal lattice. Figure 3 shows the SAXS results of -form PPDOs annealed at different Ta’s. As shown in Figure 3a, the scattering peak kept almost unchanged after annealing at Ta  50 °C and shifted to small q as Ta increased from 60 to 90 °C. On the basis of SAXS profiles, long period (LP), thicknesses of crystalline lamellae and amorphous layer (dc, da) were estimated from the one-dimensional correlation function r(z) (Figure 3b)48 

 I  q  q cos  qz  dq r  z   I  q  q dq 0

2



2

0

where I(q) is the scattering intensity, q is the scattering vector, and r(z) is the electrondensity correlation length. As shown in Figure 3c, dc of -form PPDOs enhanced but its da decreased after annealing; dc of PPDO increased from 3.0 to 4.2 nm and its da decreased from 4.1 to 3.0 nm after annealing at 40 C. LP changed little with Ta at Ta  50 °C but increased notably with Ta at Ta > 50 °C. As Ta increased from 40 to 90 °C, LP and dc increased from 7.2 and 4.2 to 7.9 and 4.6 nm, respectively. These results demonstrated the crystalline thickening during the annealing-induced ′-to- phase



transition. The phase transition can undergo through either solid-to-solid or meltcrystallization mechanism. In general, dc of polymers increased during phase transition with the melt-crystallization mechanism, but varied little during the solid-to-solid phase transition.49,50 Therefore, it is considered that the annealing-induced ′-to- phase transition progressed in the melt-crystallization mechanism. Due to the partial melting of ′ crystals and subsequent recrystallization during annealing, dc of PPDO increased.

b

a

Ta (C)

Intensity (a.u.)

90 80 70 60 50 40 5

unannealed

0.5

1.0

1.5

90 80 70 60 50 40 5

r (z)

Ta (C)

unannealed

2.0

0

-1

10

q (nm )

20

30

z (nm)

c 8

Thickness (nm)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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LP

7 dc 4

da

3 2

0

10 40 50 60 70 80 90

Ta (C) Figure 3. SAXS results of -form PPDOs after annealing at different Ta’s for 20 min: (a) SAXS patterns; (b) one-dimensional correlation function; (c) long period, lamellar thickness, and amorphous thickness. Annealing Effects on Mechanical Properties. Due to the enhancement of Xc


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and the occurrence of phase transition during annealing, it expected that mechanical properties of PPDO would be influenced by annealing at elevated temperatures. Figure 4 shows the results of mechanical properties of -form PPDOs after annealing at various Tc’s for 20 min. The engineering stress-strain curves of PPDO in the smaller strain region (030%) were enlarged and shown in Figure S2. As shown in Figure 4a, the engineering stress-strain curves of PPDO were peculiar; the stress fluctuated during the uniaxial tensile test, similar as some other polymers.28,51 The -form PPDO crystallized at 5 °C was flexible and ductile, having a σy of 15.9 MPa, an E of 141.1 MPa, and a εb of 7.4. σy and E increased while εb decreased gradually as Ta increased from 40 to 90 °C. PPDO annealed at 90 C for 20 min had the σy of 34.1 MPa, E of 238.6 MPa and εb of 3.8. The annealed PPDOs had the σy of > 20.7 MPa, E of > 166.6 MPa and εb of < 6.4. Mechanical properties of PPDO could be tuned within a wide range (σy of 20.7−34.1 MPa, E of 166.6−238.6 MPa and εb of 3.8−6.4) by varying Ta. The annealing-induced stiffening of PPDO was ascribed to the enhancement of Xc and the formation of more ordered crystals during annealing. Similar variation of mechanical properties during annealing have been reported for the other polymers such as PLLA.10,52 These results further indicated that the ′-form PPDO had larger ductility but lower mechanical modulus and strength than its  counterpart. Therefore, the annealing-induced crystalline phase transition can be a feasible way to control the mechanical properties of PPDO.



b unannealed 50 C 70 C 90 C

Stress (MPa)

60

 



40

Yield stress, y (MPa)

a



20

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30

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0

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Breaking strain, b

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Ta (C)

Strain (%) Young's modulus, E (MPa)

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250 200 150

unannealed

6

4

unannealed

100

0

20

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60

Ta (C)

80

100

0

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40

Ta (C)

Figure 4. Mechanical properties of -form PPDOs after annealing at different Ta’s for 20 min: (a) engineering stress-strain curves; (b) yield strength; (c) Young’s modulus, (d) breaking strain. Stretch-Induced Polymorphic Phase Transition. Stretching Effects on Structural Changes of ′-form PPDO. Polymorphic structural transition of PPDO during stretching was investigated via synchrotron radiation WAXD. PPDOs crystallized at Tc = 5, 60 and 90 °C were stretched at Ts = 25, 40 and 60 °C, respectively. We first focused on the crystalline structural changes of ′-form PPDOs (Tc = 5 °C) during stretching at different Ts’s. Figure S3 shows the selected WAXD patterns of ′form PPDOs (Tc = 5 °C) after stretching at 25, 40 and 60 °C to different H’s; these


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patterns were all collected at 25 °C with the release of stress. For the ′-form PPDO (Tc = 5 °C) stretched at 25 °C, the diffraction peak was almost unchanged after stretching to different H’s, except for the lamellar orientation (Figure S3a). However, when Ts increased to 40 C (Figure S3b) or 60 °C (Figure S3c), the diffraction peak at 2 = 19.2°, characteristic of the (020) plane of  form, became more pronounced at a low H (~0.65), indicating the occurrence of ′-to- phase transition. As shown in Figure 2a, no phase transition was detected for ′-form PPDO (Tc = 5 °C) during annealing at 40 °C. Therefore, we concluded that stretching promotes the ′-to- phase transition of PPDO during annealing at elevated temperatures, which might be attributed to the decrease of energy barrier for polymorphic structural transition under stretching.53

b

18

21



310

020

210

H

1.39 1.37 1.26 1.21 1.19 1.02 0.92 0.76 0

 15

Intensity (a.u.)

310

020



210

a Intensity (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


1.16 1.10 1.04 0.84



24

0.65 0

15

2  1.24Å

H 1.23

18

21

24

2  1.24Å

Figure 5. WAXD patterns of -form PPDOs (Tc = 60 °C) after stretching at different Ts’s: (a) Ts = 25 °C; (b) Ts = 60 °C. The patterns were collected at 25 °C with the release of stress. Stretching-induced -to-′ phase transition. We further investigated the crystalline structural changes of -form PPDO during stretching at different Ts’s. The H-H curves of -form PPDOs (Tc = 60 °C) stretched at various Ts’s were depicted in Figure S4. At the same H, H of PPDO decreased with the increase of Ts, due to the 15 ACS Paragon Plus Environment


enhanced chain mobility. Figure 5 shows the WAXD patterns of -form PPDOs (Tc = 60 °C) after stretching at Ts = 25 and 60 °C. The patterns were collected at 25 °C with the release of stress. As shown in Figure 5, three diffraction peaks were clearly seen at 2 = 17.8°, 19.2° and 23.4° for the -form PPDO before stretching. However, for the -form PPDO stretched at 25 °C, the diffraction peak at 2 = 19.2° disappeared with

ABS (a.u.)

2880

1748

b

2926

2957

2979

2987

a

1734

stretching to H  1.26, indicating the formation of -form crystals at large H.

ABS (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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H 1.39 1.37 1.26 1.21 1.19 1.02 0.92 0.76 0

H 1.39 1.37 1.26 1.21 1.19 1.02 0.92 0.76 0

1780 1760 1740 1720 1700

3050 3000 2950 2900 2850 2800

Wavenumber (cm-1)

Wavenumber (cm-1)

Figure 6. FTIR spectra of -form PPDOs (Tc = 60 °C) after stretching at 25 °C to different H’s: (a) 3050−2800 cm−1; (b) 1790−1700 cm−1. The spectra were collected at 25 °C with the release of stress. Stretch-induced structural changes of -form PPDO were also investigated via FTIR spectroscopy. Figure 6 illustrates the FTIR spectra of -form PPDOs (Tc = 60 °C) after stretching at 25 °C to different H’s. As shown in Figure 6, characteristic splitting bands of -form PPDO were observed at 2987, 2979 and 1734 cm−1, assigned to the stretching vibrations of CH2 and C=O bond. Furthermore, the C=O stretching band (1780−1700 cm−1) became much broader with stretching. The spectrum of PPDO was featured that of -form PPDO after stretching to H = 1.39. These results were consistent with the WAXD results and further demonstrated the occurrence of -to-


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phase transition of PPDO during stretching at low Ts. Because the splitting bands are generally originated from the interchain interactions in ordered structure, the stretching decreased the intermolecular interactions and chain packing order of PPDO. It is notable that the stretch-induced -to- was irreversible and the formed  crystals could exist stably after the removal of applied stress. Since the -form PPDO has better flexibility than its  counterpart,3 the stretch-induced -to- phase transition provided an additional way to tailor the mechanical properties of PPDO through controlling the polymorphic structure. Stretch-induced phase transition of PPDO was strongly influenced by Ts. As for the -form PPDOs (Tc = 60 °C) stretched at 40 (Figure S5) and 60 °C (Figure 5b), three major diffraction peaks were observed during stretching, meaning that the -to- phase transition did not occur and the sample also kept as the  form during stretching. As mentioned above, stretching at Ts  40 C promoted the -to- phase transition of PPDO that initially contained  crystals. It is reasonable to consider that the -to- phase transition of PPDO would be prevented with stretching at high Ts. Therefore, Ts played an important role in the stretch-induced phase transition of -form PPDO. Similar phase transition has been reported in other polymers such as iPP.54-56 The form iPP transformed into oriented mesophase during stretching at low Ts; while no phase transition was seen during stretching at high Ts. To elucidate the thermal stability of  crystals formed by stretching  crystals, we investigated the structural changes of oriented  crystals upon heating by temperature-variable WAXD. Figure 7a shows the 1D-WAXD patterns of oriented  crystals collected upon heating. The oriented  form transformed to its  counterpart upon heating, corresponding to the appearance of (020) diffraction peak in WAXD



profiles. This confirmed that the  crystals of PPDO formed under stretching were metastable and retained to the thermally-stable  crystals upon heating. This was similar to the stretch and thermal-induced phase transition between the metastable and stable phases of poly(butylene adipate) (PBA).53 In the case of PBA, stretching caused the transition from stable  phase to metastable  phase; while the following heating induced the recovery of metastable  to stable  phase. Assuming that the crystals were prefect, the crystallite size obtained from the (210) diffraction peak (D210) was calculated from the half widths of (210) diffraction peak shown in Figure 7a, according to the Scherrer formula.57 As shown in Figure 7b, D210 of PPDO increased during heating, demonstrating the crystalline perfection at high temperature.

b 30  310

120C



16

18

20

D210 (nm)

020

210

a Intensity (a.u.)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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20

15

35C

22

25

24

40

2  1.24Å

60

80

100

Temperature (C)

120

Figure 7. Temperature-variable WAXD results of stretched PPDO ( form) collected during heating: (a) 1D-WAXD profile; (b) change of D210 during heating. PPDO was first crystallized at 60 °C and then stretched at 25 °C to a H of 1.39. Thermal behavior of stretched PPDO were also investigated via DSC. Figure 8a shows the DSC heating curves of -form PPDOs (Tc = 60 °C) after stretching at 25 °C to different H’s. The unstretched PPDO (H = 0) exhibited a small Pexo (indicated by an arrow) at 91.7 °C prior to the Pm, originated from the conventional meltrecrystallization mechanism.3,58 The Pexo was absent after stretching and only one Pm 18 ACS Paragon Plus Environment

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was seen in the melting region of stretched PPDO. The temperatures of Pm of stretched PPDOs with different H’s were very similar. All the stretched samples showed a small endotherm at ~45 °C, as indicated by asterisk in Figure 8a. This small endotherm might be induced by the stretch-induced secondary crystallization of PPDO.

a

b

Pm

6

70

H

5

1.39

* * * * * * * *

4 3 2 1

1.37 1.26 1.21 1.19

0

30

1.02 0.92 0.76 0

Pexo

0 60

Xc,DSC ()

Heat flow (mW/mg) Endo up

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


90

Temperature (C)

65 60 55 50

120

0.0

0.5



1.0

1.5

Figure 8. DSC results of -form PPDOs (Tc = 60 °C) after stretching at 25 °C to different H’s: (a) DSC heating curves; (b) change of crystallinity. However, several small endotherms or shoulders were present on the dominant endotherm of stretched PPDO with H  1.26, leading to the multiple melting behavior of highly-stretched PPDOs. Multiple melting behavior can be ascribed to the different crystal modifications, different lamellae morphologies, and melt-recrystallization of crystals.59 As shown in Figure 5a, only ′ crystals existed in PPDOs stretched at 25 °C to H  1.26. As reported in the literatures,56,60 both chain-folded and chain-extended crystals are present in stretched crystalline polymers. Therefore, multiple melting behavior of stretched PPDOs with H  1.26 can be ascribed to the different morphologies of crystals. Figure 8b shows the Xc’s of -form PPDOs stretched at 25 °C to various H’s. Xc improved slightly by ~5% during stretching at H = 0−1.02 and then kept almost unvaried with further stretching to H = 1.39.



Figure 9. Selected 2D-WAXD patterns of -form PPDOs (Tc = 60 °C) collected during stretching at Ts = 25, 40, and 60 °C. The stretching direction is vertical. In-Situ Investigation on Stretch-Induced Polymorphic Structural Transition. Figure 9 shows the selected in-situ 2D-WAXD patterns of -form PPDOs (Tc = 60 °C) collected during stretching at Ts = 25, 40 and 60 °C. Stretching direction was vertical and defined as meridian hereafter. As shown in Figure 9, 2D-WAXD patterns of form PPDOs collected at different Ts’s varied in similar trends under stretching. For the -form PPDO (Tc = 60 °C) stretched at Ts = 25 °C, three main isotropic diffraction rings were observed before stretching, corresponding to the (210), (020) and (310) diffractions of -form PPDO (from inner to outer), respectively.4 These diffraction rings orientated gradually to the equator with stretching to H  0.76, reflecting the orientation of crystals. However, the -form PPDOs (Tc = 60 °C) stretched at Ts = 40 and 60 °C began to orientate to the equator at H = 1.00, signifying the slower orientation of lamellae at high Ts.


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15

310

020

18

b

Ts = 25 C

Intensity (a.u.)

210

a

21

2  Å

H 1.36 1.34 1.33 1.31 1.30 1.27 1.25 1.23 1.22 1.16 1.12 1.02 0.90 0.76 0.15 0

24

Ts = 25 C

90

180

H 1.36 1.34 1.33 1.31 1.30 1.27 1.25 1.23 1.22 1.16 1.12 1.02 0.90 0.76 0.15 0

0

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300

Azimuthal angle ()

d

c

60

Xc,WAXD ()

0.0

f210/020

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Intensity (a.u.)

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-0.1 Ts = 25 C

50 40 Ts = 25 C

30

Ts = 40 C

Ts = 40 C

Ts = 60 C

-0.2

0.0

0.5

H

Ts = 60 C

1.0

20

1.5

0.0

0.5

H

1.0

1.5

Figure 10. In-situ WAXD results of -form PPDOs (Tc = 60 °C) collected during stretching at various Ts’s.: (a) 1D-WAXD patterns collected at Ts = 25 °C; (b) azimuthal diffracting intensity distribution curves for the sample collected at Ts = 25 °C; (c) orientation degree of (210)/(020) diffractions; (d) change of crystallinity. Figure 10a and S6 show the 1D-WAXD patterns of -form PPDOs (Tc = 60 °C) during stretching at 25, 40 and 60 °C, which were derived by integrating the 2D-WAXD patterns (Figure 9). As shown in Figure 10a, the (020) diffraction gradually disappeared upon stretching, confirming the occurrence of -to- phase transition with the applying of stress. Based on the 1D-WAXD results, the azimuthal intensity distributions along scattering circles of (210) and (020) diffractions were evaluated and depicted in Figure 10b. Azimuthal intensity distribution curves of (210) and (020) 21 ACS Paragon Plus Environment


diffractions kept their peaks at 90° and did not vary with Ts. According to the azimuthal intensity distributions curves, orientation degree (f) was calculated and shown in Figure 10c. It is notable that f of PPDO ranged in −0.5~0. f was 0 for an isotropic sample with random orientation and −0.5 for a perfect orientation of the lattice planes with their normal perpendicular to the stretching direction. As shown in Figure 10c, f of PPDO was close to zero before stretching, indicating the random orientation of crystals. f was small during initial stretching (H = 0−0.15), because most of the oriented crystalline chains were perpendicular to the stretching direction. Then, f dropped distinctly upon further stretching to H = 1.07; however, f varied little at H  1.07. Moreover, the variation tendency of f with H depended strongly on Ts. PPDO stretched at low Ts (25 C) had smaller f values than that stretched at high Ts (40 and 60 C) at H = 0.60−1.02, which was caused by the increased mobility of polymer chains and lower stress at high Ts. The Ts dependence of f was similar to the results of poly(butylene adipate-ranterephthalate)60 but opposite to those of iPP56 and PBS.61 Figure 10d shows the Xc,WAXD change of -form PPDOs (Tc = 60 C) during stretching at various Ts’s. Xc’s derived from WAXD showed similar H dependence but slightly smaller values than those derived from DSC for the -form PPDO (Tc = 60 C) stretched at 25 C. At Ts = 25 and 40 °C, Xc,WAXD improved slightly during stretching at H = 0−0.81, but changed little with further stretching at H  0.87. Xc,WAXD’s of PPDO increased by about 4% and 8% during stretching at 25 and 40 °C, respectively. Xc,WAXD of PPDO stretched at 40 °C was slightly larger than that of PPDO stretched at 25 °C with H  0.81, because of the formation of thicker and ordered crystalline lamellae at high Ts. However, Xc,WAXD of PPDO decreased gradually with stretching at Ts = 60 °C, indicating that the original lamellar crystals partially melted during stretching at high Ts even though the used Ts was lower than the melting temperature.19 22 ACS Paragon Plus Environment

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Some molten amorphous chains also recrystallized into the  crystals under further stretching, as demonstrated by the following SAXS results. Stretch-Induced Lamellar Structural Evolution. Figure 11 depicts the selected 2D-SAXS patterns of -form PPDOs (Tc = 60 C) collected during stretching at 25 and 60 °C. As shown in Figure 11, PPDO showed the isotropic scattering rings before stretching, reflecting the randomly distributed lamellae in the unstretched PPDO (H = 0). SAXS patterns changed obviously during stretching, because of the change of lamellar structure. For the PPDO stretched at 25 C, an elliptical-shaped scattering ring appeared in the meridian direction at H = 0.100.14, indicating the orientation of crystalline lamellae along the stretching direction. Then, the elliptical-shaped scattering rings changed into the two-bar scattering at H  0.26; this was different from the fourpoint patterns that were usually observed in the stretched semicrystalline polymers.6265

The four-point patterns are originated from the checkerboard-like arrangement of

lamellae62 or the tilting of fragmented lamellae from stretching direction.63 The absence of four-point pattern in stretched PPDO could be interpreted by the weak scattering intensity of tilted lamellae or the orientation of most lamellae perpendicular to the stretching direction.65

Figure 11. Selected in-situ 2D-SAXS patterns of -form PPDOs (Tc = 60 °C) collected at Ts = 25 and 60 °C. The stretching direction is vertical. SAXS patterns of PPDO stretched at 60 C changed in a similar manner as that



stretched at Ts = 25 C, except for the magnitude of scattering intensity. However, strong streaks were observed in the equatorial direction at H  0.14 for PPDO stretched at Ts = 25 and 60 C, indicating that the elongated heterogeneous structures (e.g., cavities or crystalline fibrils) with different electron densities were formed along the meridian direction.56,61,62 We considered that the presence of strong equatorial streaks was caused by the formation of fibrillar crystals, because no strain whitening (corresponding to the formation of cavities) was observed during stretching.56,61,66-68 Stretch-induced the formation of extended-chain crystals after melt-recrystallization has been widely reported in the literature.55,60,63 To elucidate the stretch-induced lamellae structural evolution, 1D-SAXS patterns along the meridian were integrated from 2D-SAXS patterns, as shown in Figure 12a,b. Scattering peak of PPDOs stretched at 25 and 60 C first shifted to smaller q and then returned to larger q during stretching. LP was evaluated from the 1D-SAXS patterns by Bragg’s equation (LP = 2π/q*), in which q* corresponded to the peak position of Lorentz-corrected SAXS profiles (Iq2~q plot; I, scattering intensity; q, scattering vector). Figure 12c shows the change of LP with H for the -form PPDOs stretched at different Ts’s. For PPDO stretched at Ts = 25 C, LP increased first at H  0.15 and then decreased at H = 0.15−0.28 during stretching. LP had a maximum value of ~8.8 nm at H = 0.15 and then changed little with further stretching to H  0.28. The stretchinduced increase of LP at H  0.15 was mainly caused by the elastic extension of chains in amorphous region. The decrease of LP during stretching to H = 0.15−0.28 was resulted from the destruction (i.e., stretch-induced fragmentation and melt recrystallization) of crystalline lamellae.61 The melting and recrystallization may take place at a such low H of around 0.1. In addition, the lamellar structural evolution of form PPDO during stretching underwent through melt-recrystallization mechanism. 24 ACS Paragon Plus Environment

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The amorphous chains escaped from the fragmented lamellae would form the chainextended structure along the stretching direction. Therefore, we considered that the equatorial streaks in SAXS results were induced by the formation of fibrillar crystals (i.e., chain-extended crystals) during stretching. Due to the destruction of crystalline lamellae, D210 of PPDO (Tc = 60 C) decreased during stretching at various Ts’s (Figure S7).

a

b Tc= 60 C, Ts= 25 C

Tc= 60 C, Ts= 60 C

H 1.37

H

Intensity (a.u.)

Intensity (a.u.)

1.25

0

0

0.5

1.0

q (nm-1)

0.5

1.5

1.0

-1

1.5

q (nm )

c 10 Long period (nm)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Tc= 60 C Ts= 60 C

9

8 Ts= 25 C

7

0.0

0.5

H

1.0

1.5

Figure 12. In-situ SAXS results of -form PPDOs (Tc = 60 °C) collected at Ts = 25 and 60 °C: (a) 1D-SAXS patterns collected at Ts = 25 °C; (b) 1D-SAXS patterns collected at Ts = 60 °C; (c) the change of long period. Stretch-induced variation of LP of PPDO stretched at 60 C was similar as that 25 ACS Paragon Plus Environment


of PPDO stretched 25 C. However, the PPDO stretched at 60 C showed larger LP than that of PPDO stretched at 25 C, ascribing to the formation of more ordered and thicker lamellae at high Ts. PPDO stretched at 25 C had lower scattering intensity than that stretched at 60 C, due to the formation of less ordered  crystals during stretching at low Ts. Therefore, scattered density decreased remarkably during stretching at 25 C, resulting from the generation of  crystals as well as the reduced thickness of specimen. Based on the WAXD and SAXS results, we explained the stretch-induced -to phase transition of PPDO at low Ts. As demonstrated by SAXS results, crystalline lamellar evolution of -form PPDO (Tc = 60 C) underwent through meltrecrystallization mechanism under stretching; some original lamellae melted into the amorphous chains and some polymer chains escaped from the fragmented lamellae into the amorphous region during stretching. These amorphous chains were extended and crystallized into the oriented  crystals during stretching at low Ts but  crystals at high Ts, because the  and  crystals were preferential formed at low and high Tc’s, respectively.3 Therefore, the stretch-induced -to- phase transition took place only at low Ts (e.g., 25 C). On the other hand, as reported in literatures,15,54 a critical stress is generally required for the occurrence of phase transition during stretching. A high Ts would facilitate the relaxation of polymer chains and thus lower the tensile stress. Therefore, the critical stress could be easily reached at a low Ts, thus facilitating the phase transition.

CONCLUSION In conclusion, PPDO underwent polymorphic phase transition and crystalline lamellar evolution during high-temperature annealing and stretching; which depended strongly on the initial structure of PPDO, Ta, and Ts. Metastable  crystals of PPDO


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transformed into the thermally stable  crystals during annealing at high Ta, accompanied by the improvement of Xc and LP with increasing Ta. Due to the annealing-induced crystalline structural changes, tensile strength and modulus of PPDO increased but its breaking strain decreased after annealing at high Ta. On the other hand, the -form PPDO transformed into metastable  one during stretching at low Ts (e.g., 25 C); while no phase transition took place at high Ts (e.g., 60 C). Stretch-induced to- phase transition was irreversible and the formed  crystals under stretching did not retain to  crystals with the release of stress. However, the formed  crystals under stretching recovered into their  counterparts during heating, due to its lower thermal stability. Crystalline lamellae of -form PPDO was stretched into the fibrillar lamellae during stretching through the melt-recrystallization mechanism. This study has clarified the effects of stretching and thermal annealing on polymorphic crystal transition and lamellae structural evolution of PPDO, which is essential for tailoring the physical properties of PPDO materials in processing.

ASSOCIATED CONTENT Supporting Information Partial WAXD data of stretched PPDO and tensile test.

AUTHOR INFORMATION Corresponding Author *e-mail [email protected] (P.P.). Notes The author declares no competing financial interest.

ACKNOWLEDGEMENTS This work was financially supported by the National Key R&D Program of China



(2016YFC1100801), National Natural Science Foundation of China (21674095), and Natural Science Foundation of Zhejiang Province, China (LR16E030003). WAXD and SAXS experiments were carried out at the beamline BL16B1 of SSRF, China.

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Polymorphic Crystal Transition and Lamellae Structural Evolution of

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