Article pubs.acs.org/Macromolecules
Strain-Induced Crystallization of Segmented Copolymers: Deviation from the Classic Deformation Mechanism Ping Zhu,†,‡ Xia Dong,*,†,‡ and Dujin Wang†,‡ †
Beijing National Laboratory for Molecular Sciences, CAS Key Laboratory of Engineering Plastics, Institute of Chemistry, Chinese Academy of Sciences, Beijing 100190, China ‡ University of Chinese Academy of Sciences, Beijing 100049, China S Supporting Information *
ABSTRACT: The classic deformation mechanism of the Gaussian model of Haward and Thackray was utilized to treat the true stress−strain behaviors of a family of novel polyether-bamide segmented copolymers based on the crystalline hard segments of polyamide1012 and the amorphous soft segments of poly(tetramethylene oxide). The results showed that the deformation behaviors of the plastic copolymers abided by the Gaussian model, causing the modulus G of the strain-induced hardening process to depend on the weight percentage of the polyamide segments in the copolymers. By contrast, the deformation of elastomeric copolymers deviated from the model because of the occurrence of strain-induced crystallization in the soft polyether sections at large strains, which negated the Gaussian assumption; i.e., the random coil conformation was maintained even under substantial stretching. The onset point of deviation was, for the first time, quantitatively identified by in situ FTIR and further confirmed by in situ WAXD/SAXS.
■
INTRODUCTION Deformation behavior and tensile properties are two important characteristics of semicrystalline polymers. Generally, external work by tensile stretching breaks and transforms the original spherulitic morphology into a highly oriented fibrillar state, where the alignment of the polymeric chains is preferentially parallel with the drawing direction.1−4 Strain hardening can be observed during this process for semicrystalline polymers and manifests as a rapid boost of stress against strain at large extensions.5,6 G’Sell et al. carried out research on strain hardening for several plastics by plotting true stress−strain curves from tensile deformation experiments with a constant strain rate, which was controlled according to in situ measurements of the diameter in the neck.7,8 Haward and Thackray put forward a model to describe the deformation mechanism of semicrystalline polymers during strain hardening, which consisted of a Hookean spring in series with a dashpot in parallel with another spring.9−11 The Hookean spring and the dashpot correlate with the movement of the crystalline domain, such as the movement through lamellar coupling and slipping, whereas the rubbery spring represents the entangled amorphous domain. This model can be further modified by a Gaussian treatment, in which the polymer coil cannot be fully stretched; thus, the elastic stress can be represented by the Gaussian equation. Two lines can be calculated from the curve of the true stress (σ) against λ2 − λ−1, yielding the Hookean elastic constant E and the rubber elastic constant G, respectively, where λ is the draw ratio. Haward successfully applied this model to several semicrystalline polymers, © XXXX American Chemical Society
including polyethylene (PE), polypropylene (PP), poly(ether ether ketone) (PEEK), polyamide 6, and polyamide 66. The Gaussian equation has been successfully used to treat the deformation of several glassy polymers, including polycarbonate (PC) and polystyrene (PS), the strain hardening of which resulted from the increase of the rate of nonaffine displacement.12 Men et al. used this model to treat miscible blends of semicrystalline polymers and amorphous polymers and concluded that the incorporation of a second noncrystallizable polymer does not change the crystal thickness of the semicrystalline polymer but expands the amorphous intercrystalline regions for poly(ε-caprolactone) (PCL), poly(vinyl methyl ether) (PVME), and PCL/styrene acrylonitrile (SAN) blends.13 Hong et al. studied the deformation behavior of a copolymer, poly(ethylene-co-12% vinyl acetate) (PEVA), and proposed a three-component model to decompose the stress into the contributions of the crystal skeleton, the amorphous network, and the viscosity.14 The microscopic mechanism underlying the deviation behavior has not been clearly explored because of insufficient understanding of the deformation behaviors of a series of semicrystalline copolymers with welldefined macrostructures. The question whether the Gaussian model can be applied to the deformation of segmented copolymers or not has not been confirmed yet. Received: December 31, 2016 Revised: April 17, 2017
A
DOI: 10.1021/acs.macromol.6b02747 Macromolecules XXXX, XXX, XXX−XXX
Article
Macromolecules
TFAA reacted with active hydrogen atoms in amide, amine, hydroxy, and carboxy function groups, making the reaction mixture trifluoroacetylated (tfaed) and dissolved in CDCl3 (Figure 1). The chemical shift (CF) of the representative protons can be easily identified after trifluoroacetylating. The CF of methylene protons in the α-position of trifluoroacetylated OH groups is 4.55 ppm; by contrast, 4.3 ppm is that of the CH2 protons in the α-position of the oxygen atom of the ester junction. The CF of methylene protons in the α-position next to the carbonyl group of trifluoroacetylated amide groups is 2.96 ppm, which shows a constant signal across the reaction; by contrast, 2.52 ppm is that of the CH2 protons in the α-position of the carbonyl group of ester junction. The integrated area of signals ascribed to protons adjacent to the newly formed ester junction between the PA1012 HS and PTMO SS can be calculated against that of respective protons close to terminal function OH/COOH groups in two oligomers. As the reaction proceeded, the proton signals brought by newly formed ester bond intensified, such as signals at 4.3 and 2.52 ppm. It can be employed to trace the reaction process against time either by the ratio of integrated area of 4.55 ppm proton and 4.3 ppm proton defined as R4.55/4.3 or by the ratio of integrated area of 2.52 ppm proton and 2.96 ppm proton defined as R2.52/2.96. For A600G2000, the esterification of DTPA1012-600 and PTMEG2000 was studied at an optimized condition of 240 °C in the presence of Ti(OBu)4 (0.5 wt % of reactive mixture) under vacuum, with reference to existing kinetics study of esterification reactions catalyzed by organometallic derivatives.15 In the first 2 h, R4.55/4.3 decreased rapidly, while R2.52/2.96 increased remarkably. The two ratios reached a plateau or an equilibrium status, when the reaction continued for 3 h. This methodology was used to determine the optimum reaction time for the rest of copolymer samples. The composition was predetermined by the molecular weight of the oligomeric PA1012 and PTMEG charged for specific polymerizations. It was finally calculated from the molar ratio obtained from the 1H NMR spectrum for each sample. A representative 1H NMR spectrum of A1700G2000 is shown in Figure S1. The reference peaks are the methylene peaks next to the oxygen on PTMO (ca. 3.8 ppm), next to the nitrogen of PA1012 (ca. 3.6 ppm), and next to the oxygen on the ester conjunction between PA1012 and PTMO (ca. 4.3 ppm).17,18 The PTMO microdomains impart the copolymers with elastomeric flexibility and extensibility, whereas crystalline PA1012 microdomains hold the matrix together via interchain hydrogen bonds, serving as thermally stable physical cross-linking sites that impart dimensional stability and mechanical strength. Samples are denoted in an abbreviated form AxxxxGyyyy, indicating the oligomeric molecular weight of each segment as a subscript (xxxx and yyyy). The composition and molecular weight of a family of nine samples of PA1012−PTMO copolymers, each with a different polyamide weight percentage (WPA), are listed in Table 1. The molecular weight of the all copolymers was determined by gel permeation chromatography (GPC, equipped with a refractive index
Therefore, the main aim of present investigation is to probe into the essence of deformation mechanism of segmented copolymers with crystalline hard segments (HS) and amorphous soft segments (SS) and to try to correlate the microstructure variation and macroscopic mechanical properties. For this purpose, we designed and synthesized a series of polyether-b-amide (PEBA) polymers based on polyamide1012 (PA1012) and poly(tetramethylene oxide) (PTMO) by polycondensation at the melt state.15 The chemical structures for the two segments are shown in Scheme 1. The Gaussian Scheme 1. Chemical Structures for the Repeating Units of the PA1012 Hard Segment (HS) and PTMO Soft Segment (SS)
model was utilized to treat the true stress−strain curves of both elastomeric and plastic copolymers. Abidance and deviation were observed for different samples, the mechanism of which was discussed in detail. The onset point of deviation from the Gaussian model has been, for the first time, quantitatively identified with a combination of Fourier transform infrared spectroscopy (FTIR), wide-angle X-ray diffraction (WAXD), and small-angle X-ray scattering (SAXS).
■
EXPERIMENTAL SECTION
Materials and Sample Preparation. PA1012−PTMO copolymers were synthesized using a two-step method. Diacid-terminated oligomeric-PA1012 (DTPA1012) was first synthesized with different oligomeric molecular weights, and PA1012−PTMO was subsequently synthesized with different compositions from oligomeric PA1012 and poly(tetramethylene ether glycol) (PTMEG) via melt polycondensation.16 The polymerization can be monitored by solubilizing the reaction mixture in a cosolvent of CDCl3/trifluoroacetic anhydride (TFAA).15
Figure 1. Structures of tfaed DTPA1012/PTMEG/PA1012−PTMO with CF for methylene protons; R4.55/4.3 and R2.52/2.96 vs time. Inset: NMR spectra at different times. B
DOI: 10.1021/acs.macromol.6b02747 Macromolecules XXXX, XXX, XXX−XXX
Article
Macromolecules
A600G2000, fibrils and dendritic crystals with lamellae are surrounded by PTMO domains. Big spherulites can be found from the image of A3530G1000, where PA1012 domains are continuous. The diameter of spherulites in A1700G2000 is around 8 μm, less than 10 μm in A1330G1000. Higher WPA leads to bigger spherulite, which can influence the deformation behaviors. All samples were dried under 100 °C vacuum, melted, and pressed under 220 °C and 50 MPa for 3 min and then quickly cooled to room temperature through cool compression. Films with the thickness of 0.5 mm and 40 μm were prepared for mechanical property tests and FTIR measurements, respectively. The small dumbbell-shaped specimens were cut using a mold. Characterization. The melting and crystallization behaviors of PA1012−PTMO copolymers were examined using a DSC (TA Instruments, Q2000). The instrument was calibrated with indium before measurements. Temperature scans were performed in the temperature range from −60 to 210 °C under a nitrogen atmosphere with a heating/cooling rate of 10 °C/min. Under the assumption of a constant volume during the elongation, the true strain was used as a measurement for the deformation, as described by Strobl and co-workers.22,23 The width, thickness, and marked distance of the initial specimen and stretched specimen at specific test points are labeled as (W0, T0, b0) and (W, T, b), respectively. The true stress σ, draw ratio λ, and true strain εH are described as follows:
Table 1. Compositions and Molecular Weights of the Samples composition in copolymer by 1H NMR
GPC after Ntrifluoroacetylation
sample code
sample name
WPA
Mn/ 104
Mw/ 104
D
1 2 3 4 5 6 7 8 9 10
A600G2000 A1330G2000 A1700G2000 A1330G1000 A2000G1000 A2800G1000 A3530G1000 A2800G650 A4630G650 PA1012
0.24 0.35 0.47 0.56 0.67 0.73 0.78 0.81 0.87 1
1.77 2.57 3.24 3.14 3.08 2.95 2.84 2.46 2.20 NAa
2.47 3.95 10.41 7.14 8.14 7.14 7.27 5.21 4.08 NA
1.39 1.54 3.21 2.27 2.56 2.42 2.55 2.12 1.85 NA
a
NA means not applicable.
detector) using chloroform as the solvent and monodisperse polystyrene as the calibration standard (Table 1). Each sample was treated with N-trifluoroacetylation to dissolve in chloroform, a common GPC solvent.19 Homopolymeric PA1012 with a melting index of 22 g/10 min was also tested in parallel, coded as sample 10. The FTIR spectra and band assignment are shown in Figure S2 and Table S1, respectively, and the band assignments are in good agreement with those of PEBA with polyamide12 (PA12) as HS.20 Differential scanning calorimetry (DSC) curves are shown in Figure S3. AFM in taping mode can give real images about the morphology of PEBA.21 Samples were dissolved in dimethylacetamide at elevated temperatures, and then thin films were spin-coated on mica by dropping the hot solution. To probe the influence of WPA on the morphologies at ambient temperature, selective images of A600G2000, A1700G2000, A1330G1000, and A3530G1000 are shown in Figure 2. For
σ=
Fb W0b0T0
λ = b/b0
εH = ln λ
Using the above equations, the true stress−strain curves were obtained at a constant crosshead speed of 6.0 mm/min by an optical-assisted measuring method.24 In situ polarized FTIR spectra were obtained by using a Nicolet 6700 FTIR spectrophotometer operated in transmission mode and equipped with a DTGS detector. Small dumbbell bars, 25 mm long, 3 mm wide, and 40 μm thick, were mounted onto a Linkam TST350 hot stage installed in the test chamber of the spectrophotometer. The spectra were collected by accumulating 16 scans at a resolution of 4
Figure 2. AFM images of A600G2000 (a), A1700G2000 (b), A1330G1000 (c), and A3530G1000 (d). C
DOI: 10.1021/acs.macromol.6b02747 Macromolecules XXXX, XXX, XXX−XXX
Article
Macromolecules cm−1, and the polarizer was alternatively switched between the parallel and perpendicular directions relative to the drawing axis. The orientation functions of PTMO macromolecular chains and PA1012 macromolecular chains, denoted as f SS and f HS, respectively, were calculated on the basis of the absorption at 1366 cm−1 (the CH2 wagging mode) and 3308 cm−1 (the amide A mode), respectively. To characterize the molecular orientations in the soft and the hard segments, the angles ψ of two absorption peaks with respect to the direction of the molecular main-chain axis were taken as 0° and 90°, respectively:20 f=
3⟨cos2 θ⟩ − 1 2
f=
3A R 0 + 2 3 − A R0 − 1
A=
A↑ A→
curves (Figure S4), however, show manifested characters. Samples with a WPA less than 0.5 are elastomers with large elongation-at-break and no yielding point, whereas samples with a WPA greater than 0.5 are typical thermoplastics with evident yielding points. For the sake of clear comparison, curves of σ plotted against λ2 − λ−1 for plastic samples 4−10 are shown in Figure S5. The curves for samples 4−10 are found to follow the Gaussian model; thus, the modulus G for these samples can be determined by fitting the curves under large deformation. All the linear fits result in sound determination coefficients no less than 0.99. In Figure 4, the modulus G of
R 0 = 2 cot2 ψ
Here, θ represents the orientation angle of a chain segment with the drawing direction, and R0 is the perfect dichroic ratio based on the angle ψ. A↑ and A→ represent the band absorption intensity when the polarizer is perpendicular or parallel to the deformation, respectively. In situ WAXD measurements were carried out at the BL14B1 beamline in the Shanghai Synchrotron Radiation Facility (SSRF). The wavelength of the radiation source was 1.24 Å. The small tensile bars were stretched at a crosshead speed of 6.00 mm/min on a Linkam TST350 hot stage. Diffraction patterns were collected using an MAR 165 detector with a resolution of 3072 × 3072 pixels (pixel size: 79 × 79 μm2). The image acquisition time was 15 s, and the sample-todetector distance was 194 mm. The Herman’s orientation function of specific diffraction lattice was defined as
3⟨cos2 Φ⟩ − 1 f= 2
2
cos Φ =
∫0
π /2
I(Φ) cos2 Φ sin Φ dΦ
∫0
π /2
I(Φ) sin Φ dΦ Figure 4. Modulus G against WPA for plastic samples 4−10: 4, A1300G1000; 5, A2000G1000; 6, A2800G1000; 7, A3530G1000; 8, A2800G650; 9, A4630G650; 10, PA1012. Inset: Gaussian model of Haward−Thackray and the corresponding Haward−Thackray curve with the calculated values of E and G.
Here, Φ is the angle between the normal direction of the crystal plane and the reference axis (equator direction). In situ SAXS measurements were carried out at the BL16B1 beamline of SSRF. The wavelength of the radiation source was 1.24 Å. The small tensile bars were stretched at a crosshead speed of 6.00 mm/min on a Linkam TST350 hot stage. An MAR 165 CCD detector with the resolution of 2048 × 2048 pixels (pixel size: 73 × 73 μm2) was used to collect the scattering patterns. The image acquisition time was 15 s, and the sample-to-detector distance was 1535 mm. All of the X-ray patterns were corrected for detector noise, air scattering, and sample absorption.
samples 4−10 also bears a linear relationship with WPA. Meanwhile, the Gaussian model and the corresponding Haward−Thackray curves for the calculation of E and G are shown in the inset. Derivations from the Gaussian model are observed on the curves for samples 1, 2, and 3, especially for sample 1 (Figure 5a). The analysis of the large strain deformation of the semicrystalline polymers is usually based on the molecular network, which consists of trapped chain entanglements in the amorphous phase and crystallites acting as physical crosslinking points. Several factors affect the modulus G. An increase in the crystallinity brings higher modulus G values for a series of PE and PEVA samples, which can be explained as an effective contribution to the network density of chains hinged on the adjacent crystallites.23 The modulus G extracted from the deformation behavior of isotactic poly(1-butene) and its ethylene copolymers also bear a positive relationship with the crystallinity.25 Lamellar coupling also plays a role in the decrease of modulus G with temperature because the viscous friction between the lamellar blocks becomes weaker with increasing temperature, while the entangled amorphous network is drawn.26 The modulus G of the blends of PCL and PVME is mainly determined by the PCL because the swelling of the amorphous regions with the flexible PVME chains does not alter the entanglement density.23 Inclusion of the stiffer SAN chains reduces the entanglement density and thus reduces the modulus G. Men et al. found that the amorphous phase as a whole determined the large deformation mechanism of the blends, though the noncrystallizable polymer
■
RESULTS AND DISCUSSION Deformation Behaviors and Gaussian Model. The true stress−strain curves of all samples are shown in Figure 3. The results indicate no obvious difference in the deformation process for all ten samples. The engineering stress−strain
Figure 3. True stress−strain curves of PA1012−PTMO copolymers and PA1012: 1, A600G2000; 2, A1330G2000; 3, A1700G2000; 4, A1330G1000; 5, A2000G1000; 6, A2800G1000; 7, A3530G1000; 8, A2800G650; 9, A4630G650; 10, PA1012. Note: samples 1−3 did not break when the clamps of the TST350 reached the maximum. D
DOI: 10.1021/acs.macromol.6b02747 Macromolecules XXXX, XXX, XXX−XXX
Article
Macromolecules
possibility of SIC of PTMO in those systems. Thus, the PTMO chains are kept in amorphous state at large deformations. The new phase plays a key role on the deformation behaviors. As representative samples, A600G2000 and A3530G1000 are further examined by in situ WAXD/SAXS/FTIR to determine why the Gaussian model fails or works. In Situ WAXD Analysis and Strain-Induced Crystallization. Haward−Thackray curves of A600G2000 and A3530G1000 are displayed in Figure 5, and the representative WAXD patterns are incorporated and compared for further analysis of the deviation and abidance. The two samples bear strikingly different morphological characteristics: dendritic crystals are confined and surrounded by amorphous PTMO matrix for A600G2000, while well-organized spherulites dominate, with strictly packed lamellae for A3530G1000 (Figure S7). Onedimensional (1D) integrated curves are presented in Figure 6
Figure 6. 1D integrated WAXD curves in the meridional direction of A600G2000 (a) and A3530G1000 (b). True strain listed for each curve, and the arrows indicating the increase of strain.
Figure 5. True stress vs λ2 − 1/λ and representative WAXD patterns during deformation of A600G2000 (a) and A3530G1000 (b). Stretching is along the horizontal direction.
was not linked to the crystal parts of the semicrystalline polymer by chemical bonds.13 Our experimental observations and the aforementioned theoretical analysis indicate that the amorphous PTMO chains and the crystalline PA1012 domains together determine the deformation mechanism under large deformation. The Tg of the PA1012 and its copolymers with a high WPA is approximately 60−70 °C, higher than the experiment temperature (Figure S3). Therefore, we deduce that the crystalline part and the glassy part of the PA1012 in homopolymer and copolymers can together influence the modulus G for the deformation behaviors at room temperature when all PTMO chains are still amorphous according to the DSC results (Figure S6). The PTMO chains dilute the PA1012 (crystalline part and glassy part) through linking to noncrystalline PA1012 chains. A lower WPA results in a lower entangled amorphous network density and thus a lower modulus G of the respective copolymers. The deformation behavior of A600G2000 shows pronounced deviation from the linearity of the Gaussian model at large strains, where a red dashed line is drawn to illustrate the deviation clearly (Figure 5). The thermal behaviors of the broken samples after deformation tests were also characterized by DSC. Peaks are observed at approximately 45 °C in the heating scan curves of A600G2000, A1300G2000, and A1700G2000, with the polyether oligomer molecular weight of 2000 g/mol, indicating a new highly oriented phase, presumably, SIC of PTMO (Figure S6).17,27,28 No new peaks are found in the DSC curves of the other deformed samples, which rules out the
for comparison of the deformation behaviors of the elastomer and plastics. WAXD images are completely isotropic with an amorphous halo in the initial states. Crystalline PA1012 in A600G2000 and A3530G1000 adopt α phase, manifested itself by (100) peak and (110)/(010) peak with d-spacings around 0.44 and 0.37 nm, respectively.29,30 It should be noted that the dspacings of (100) and (010) peaks for PTMO crystal are reported around 0.45 and 0.38 nm, respectively.28,31 There are substantial overlaps for d-spacings of crystalline PA1012 and PTMO crystals. The sharpness of original α(100) and α(010)/ (110) diffractions of A3530G1000 vanishes and is replaced with a broad one at large strains (Figure 5b). This behavior is the same as the crystal transition observed during the deformation of homopolymer PA1012.29 The broad peak can be denoted as the transient phase of PA1012. No additional WAXD peaks are identified through the entire deformation process, indicating the absence of SIC of PTMO for A3530G1000. During deformation, a definitive difference is observed in the oriented A600G2000 sample as compared to the original pattern. Instead of the amorphous PTMO halo, anisotropic A600G2000 exhibits two strong vertical reflections, of which the d-spacing values are 0.44 and 0.37 nm, respectively (Figure 5a). SIC of PTMO has been well documented for PA12−PTMO copolymers (commercialized on early 1980s), and the diffraction peaks of PTMO crystals can be distinguished from γ(001) diffraction of PA12 with d-spacing around 0.4 nm.17,32 Whether SIC occurs to PTMO or not should be carefully analyzed for A600G2000 since the superposition of crystalline PA1012 diffractions and PTMO crystal diffractions. E
DOI: 10.1021/acs.macromol.6b02747 Macromolecules XXXX, XXX, XXX−XXX
Article
Macromolecules
Figure 7. Normalized 1D integrated WAXD curves of A600G2000 under selective λ (a) and 1D integrated WAXD curves at λ = 1 and 5.5, unloaded state, annealed state, and the difference between the latter two states (b).
Figure 8. Selective FTIR spectra of A600G2000 at different λ (a) and development of absorbance around 1009 cm−1 (b).
which has previously been used to characterize the SIC of PTMO in polyurethane−urea.31 Selective FTIR spectra of A600G2000 at different λ are displayed in Figure 8a, while the development of absorbance around 1009 cm−1 is highlighted in Figure 8b. The band at 1105 cm−1 corresponding to the C−O− C antisymmetric stretching mode is saturated, since the relative content of PTMO is very high in this sample. The variable f SIC is defined as the integrated area of the absorption band at 1009 cm−1 divided by that of 2860 cm−1 to eliminate the thickness effect. The value of f SIC begins to increase remarkably after λ = 1.73 (Figure S9). By contrast, soft segment orientation function ( f SS) increases after the deformation starts (Figure 9). This behavior means that the PTMO chain is first stretched from a random coil conformation to align with the external force and
First, the diffraction intensity of one crystalline part will definitely decrease during stretching if the relative content is constant, since the specimen thickness decreases. Given a new phase can be induced by strain, its diffraction intensity may not decrease but increase. In order to quantitatively analyze the intensity, all 1D integrated curves of WAXD patterns must be normalized by sample thickness T as below.24 T = T0 b0 /b
The normalized intensity of the new phase increases with draw ratio λ (Figure 7a). Thus, the new phase cannot be ascribed to the original crystalline PA1012 but a strain-dependent component or strain-induced crystallization of PTMO. Second, after the A600G2000 was stretched to λ = 5.5, the sample specimen was unloaded and the pronounced signals of newly formed crystals were still present (Figure 7a). The sample was heated from room temperature to 55 °C, which can fuse the oriented phase according to DSC (Figure S6), but exerting little influence on PA1012 HS. Difference between two 1D integrated WAXD curves can highlight the difference caused by external fields.18 The pink curve in Figure 7b is obtained by subtracting the curve after annealing from the unloaded one (picture enlarged in Figure S8). It shows the diffraction of amorphous PTMO increases remarkably at the cost of two peaks at 2θ = 16o (d-spacing = 0.44 nm) and 2θ = 19.5o (d-spacing = 0.37 nm); thus, the latter two can be ascribed to the SIC of PTMO. Third, the crystallization of PTMO can be monitored by following the variation of absorption band around 1009 cm−1,
Figure 9. f HS and f SS vs λ for sample A600G2000. F
DOI: 10.1021/acs.macromol.6b02747 Macromolecules XXXX, XXX, XXX−XXX
Article
Macromolecules
intrafibrillar scattering. The two streaks in the transverse direction are caused by interfibrillar scattering.33 The reestablishment of the ordered structures with different q position is possibly consistent with the further development of SIC of PTMO (Figure S11a). For A600G2000, azimuthal scans are taken for the two SAXS patterns at λ = 5.5 and the relaxed state, which are in the range from −90° to 90° and at respective q position of maximum intensity (Figure S12). The singe maxima split after the stress removed, indicating a four-point 2D pattern for the “unloaded” pattern in Figure 10.37 The structure may consist with two sets of crystalline components: not only relaxed and tilted PA1012 crystals but also PTMO crystals, because SIC of PTMO is not fully fused upon retraction (Figure 7b). The evolution of the lamellae structure in A3530G1000 is similar to that during stretching of PA1012 and PA1212 at room temperature.29,38 The structure change is so quick that the tilting of broken lamellae is not observed. Strong signals at low q can be ascribed to the microvoids formed during stretching, whereas no evident signals about fibrillated structures are detected (Figure 9b and Figure S11b). A twolobe pattern is witnessed since the γ-phase component can be induced and the lamellae fibrillated during PA1012 stretching at 100 °C. This phenomenon is not observed for A3530G1000 because it lacks the thermal activation at room temperature. TEM images can help explain the difference between the lamellae evolution of the two samples during deformation (Figure S7). Onset Point of the Deviation from Gaussian Model and Microstructural Orientation. On the basis of the above analysis, it can be concluded that the deformation behavior of PEBA deviates from the Gaussian model because of PTMO SIC. Can the starting point of deviation be identified? To the best of our knowledge, a quantitative description of the influence of PTMO SIC on the deformation behavior has never been reported. However, the exact amount of the PTMO crystal content cannot be extracted from the WAXD data for multiple components (PTMO crystals, amorphous PTMO, crystalline PA1012, and amorphous PA1012), the peaks of which are convoluted in the 2θ range from 16° to 20° under large strains. We observed that the role of PTMO SIC can be evaluated by plotting the true stress against f SIC. A turning point is clearly found at λ = 3.8, and soon afterward, the slope of the true stress vs f SIC increases and a linear relationship is observed (Figure 11). Therefore, the point at λ = 3.8 could be defined as the onset of the deviation from the linearity of the Gaussian model at large strains. The second clue can be found in the orientation function of the hard segment chains, f HS, which becomes positive after the initial stage of the negative trend at λ = 3.8 (Figure 9). This negative−positive transition can be explained by Bonart’s model of orientation in segmented polyurethane with a low hard-segment content which can be either in a fibrillar or lamellar morphology.34 The fragmentation from lamellae to fibrils morphology is completed at the negative−positive transition point. Thermoplastic polyester elastomers with PTMO as the soft segments and poly(butylene succinate) (PBS) as the hard segments also demonstrate similar orientation behaviors during stretching. Lee et al. postulated that the negative orientation of the crystalline polyester hard segments implied that the integrity of the hard segments was maintained for PBS−PTMO elastomers.35
then is further stretched to a regular packing conformation, which is a prerequisite for SIC. Moreover, the onset of f SIC coincides with the point where f HS reaches its negative minimum, which will be discussed together with SAXS observation later on. All aforementioned three evidence clearly and explicitly prove that PTMO SIC occurs during deformation of A600G2000. The in situ WAXD patterns of the elastomeric samples A1330G2000 and A1700G2000 and the plastic samples A1330G1000 and A4630G650 are shown in Figure S10. The first group shows SIC signals similar to those of A600G2000 (Figure S10a,b). The signals of the second group bear similarity to that of A3530G1000 (Figure S10c,d). These results further support our hypotheses. The analysis of the in situ WAXD of A1700G1000, A2800G1000, and A2800G650 is not given here because their evolution trends are the same as that of A3530G1000. In Situ SAXS and FTIR Analysis. Analysis of in situ SAXS and FTIR can provide additional information to build a whole picture of the microscopic evolution during deformation. Several representative SAXS patterns of A600G2000 and A3530G1000 are displayed in Figure 10. 1D-SAXS distribution
Figure 10. Representative SAXS patterns during deformation for A600G2000 (top) and A3530G1000 (bottom). The numbers indicate λ. Stretching is performed along the horizontal direction.
files integrated along horizontal direction are shown in Figure S11. The evolution of SAXS patterns for A600G2000 is similar to that previously reported for a commercial PA12-PTMO elastomer, Pebax 3533.33 The scattering patterns first convert the circular ring into an ellipsoidal form with the ellipse radius decreasing in the horizontal direction. The value of long period in the stretch direction increases as expected because the PA1012 lamellae are stretched further apart. After the λ reaches 1.73, SAXS pattern changes into two arcs on the horizontal direction, along with a scattering streak looms on the vertical direction. This drastic change is in good agreement with the observation by in situ FTIR: the f HS decreases from zero, reaching negative minimum at λ = 1.73, where the original lamellae of PA1012 break down and SIC of PTMO starts to take place (Figure 9). The negative orientation of the hard segment has been observed in polyetherester and polyurethane.34−36 The gradual development of streak indeed confirms the formation of microfibrillar structure. After the λ reaches 2.8, two arcs change into two lobes on the horizontal direction, along with the streak on the vertical direction turns sharpened. The lobes result from G
DOI: 10.1021/acs.macromol.6b02747 Macromolecules XXXX, XXX, XXX−XXX
Article
Macromolecules
also analyzed based on the WAXD analysis. The Herman’s orientation functions of the lattice diffractions with d-spacing = 0.44 nm (f H‑0.44) for A600G2000 and A3530G1000 are shown in Figure S14, further advising the different deformation mechanisms determined by different WPA. The f H‑0.44 of PA1012 lattice in plastic A3530G1000 responses to external work very sensitively and plunges after λ reaches 1.5. No straininduced crystallization occurs to A3530G1000 during stretching (Figure 6b). The f H‑0.44 for the elastomeric A600G2000 does not turn negative remarkably until λ reaches 2.8. The information from f H‑0.44 for the elastomeric A600G2000 denotes the crystalline lattice orientation not only from crystalline PA1012 but also from PTMO crystals, since there is substantial overlap between the d-spacings of crystalline PA1012 and PTMO crystal at 2θ = 16°. The crystalline PA1012 may be inert, since it can be read from the 2D WAXD patterns that the orientation is mainly due to the strain-induced crystallization of PTMO (Figure 5a). Quantitative comparisons between f SS based on FTIR and the Herman’s orientation functions based on WAXD are also carried out. After plotting the f H.PTMO(010) against f SS, a twostage correlation between the two parameters can be seen (Figure 13). In the first stage from the beginning of
Figure 11. Correlation curve of the true stress vs f SIC of PTMO under large strains.
The third clue can be found by tracing Herman’s orientation function ( f H) of a diffraction lattice, which can evaluate the angular position of specific lattice against drawing direction.39,40 Referring the stretching direction as the zero degree, the Herman’s orientation of PTMO (010) (f H.PTMO(010)) based on the azimuthal scans of WAXD diffraction at 2θ = 19.5° is shown in Figure 12. The normal of lattice is perpendicular to the
Figure 13. Plot of f H.PTMO(010) vs f SS.
Figure 12. Plot of f H.PTMO(010) vs λ.
deformation to λ = 1.73, the increase of f H.PTMO(010) is very slow. f SS takes a huge leap forward from λ = 1.73 to λ = 1.9. Consequently, the slope of the curve turns bigger after the latter point, which marks the onset of the second stage. The critical point at λ = 1.73 is also just the very point where f HS reaches the negative minimum as well as f SIC turns positive. This correlation also supports one of our findings that the random coils of PTMO are first stretched and aligned along the deformation direction, leading to the increase of f SS, and further stretching induces crystallization after the turning point. Multistage Deformation Mechanism. The microscopic deformation mechanism is illustrated in Figure 14 for A600G2000. During the deformation process from λ = 1 to 1.73, the PA1012 domains are oriented by a continuum mechanical stress, transferred from the PTMO domains, and finally rotate along their long axes toward the drawing direction. The PA1012 domains with fibrillar morphology are positively oriented, while the PA1012 domains with lamellar morphology are negatively oriented. At the beginning stages, the negative orientation dominates, leading to a negative trend (Figure 9). Larger stresses destroy the lamellar hard-segment domains into smaller fragments of hard-segment aggregates with their long axes oriented perpendicular to the chain direction from λ = 1.73 onward. At the fragmentation point of λ = 1.73, the negative orientation reaches the minimum value, and the SAXS pattern
drawing direction when f H = −0.5, while parallel to the drawing direction when f H = 1. f H equals to zero when there is random orientation in the sample. The original calculated f H.PTMO(010) turns negative, indicating that the lattice of PTMO (010) turns gradually perpendicular to the drawing direction; in other words, the lattice of PTMO (010) turns gradually parallel to the drawing direction. During the lamellae−fibril transition from λ = 1.73 to λ = 3.8, the lattice of PTMO (010) orients quickly, but after the turning point, the trend slows down. For the sample A600G2000, the evolution of crystal thickness during the deformation is complicated, since at least four phases coexist when λ is higher than 1.73, e.g., strain-induced PTMO crystals, amorphous PTMO, crystalline PA1012, and amorphous PA1012. The crystal thickness of the hard segment in A600G2000 can be determined neither by the conventional electron density distribution function K(z) for 1D model nor by the form factor method, since multiply phases coexist or the morphology is totally different from that of oxalamide−PTMO copolymers.24,28,41,42 The PTMO crystal sizes (D010) can be estimated from 1D WAXD diffraction of lattice of PTMO (010) by using Scherrer’s formula, which also shows two turning points at λ = 1.73 and λ = 3.8 (Figure S13). The Herman’s orientation functions of the crystalline PA1012 components in the two representative samples are H
DOI: 10.1021/acs.macromol.6b02747 Macromolecules XXXX, XXX, XXX−XXX
Article
Macromolecules
Figure 14. Illustration of the microscopic deformation of the PA1012 HS and PTMO SS in A600G2000 at three critical points during uniaxial stretching in the horizontal direction. Fibrils of PA1012 HS can be observed from the TEM image (Figure S7b). L and F represent the lamellae and the fibrils forms with the direction of their long axes (both in red dashed arrows) drawn differently in the original state and then in the negative orientation (N) and positive orientation (P) in the interval 1 ≤ λ ≤ 1.73. When λ reaches 1.73, B denotes the breakdown of the lamellar form into fibrillar form. O denotes the orientation of amorphous PTMO and fibrillar PA1012 toward the stretching direction. SIC means the strain-induced crystallization of PTMO. When λ reaches 3.8, the lamellae−-fibril transition is completed, and SIC of the PTMO starts to dominate the increase of stress, marking the onset of the failure of the Gaussian model. The amorphous PA1012 is omitted.
changes dramatically; soon after the fibrillar signal is found in the SAXS pattern. The hard segment orientation function turns positive at λ = 3.8, indicating the positive orientation of the fibrillar PA1012 started to dominate. The fragmentation from lamellae to fibrils morphology is completed. Regarding the in situ WAXD, the prominent signals with d-spacings of 0.44 and 0.38 nm in the vertical direction emerge at λ2 − 1/λ = 14.21 (λ = 3.8) (Figure 5a). The trend of Herman’s orientation function of PTMO (010) lattice shows a turning point (Figure 12). The PTMO chain orientation and regular arrangement reach a critical point that makes SIC remarkably evident. Before this critical point is reached, some PTMO paracrystals are formed, and then the macroscopic stresses are mainly dependent on the PTMO SIC after the point. Under the same analysis protocol, λ = 3.3 and λ = 3 mark the onset points of deviation from the Gaussian equation for A1330G2000 and A1700G2000, respectively (Figure S15). The deviation of the A1700G2000 curve is not as evident as that of A600G2000. The Gaussian equation is found to fail at smaller strains for elastomers with higher WPA values. The question arises as whether our proposal can explain the deformation of the other PEBA samples. As shown in Figure S16, the plot of σ against λ2 − λ−1 for the commercial elastomer Pebax 2533, a PA12−PTMO copolymer, is found to deviate from the Gaussian model because its PTMO phase crystallize at large strains.17,34,43,44 By contrast, the deformation behavior of plastic Pebax 7033 can be treated by the Gaussian model, for its PTMO phase never crystallize at large strains. Its 1D integrated WAXD intensity trends show no PTMO SIC signals (Figure S16 inset) and instead indicate the transient phase of PA12, induced from the original γ-phase.18,45,46 The reason for the abidance and the deviation is the same as for our lab-made samples; i.e., PTMO SCI in Pebax 2533 results in deviation
from the Gaussian model, while Pebax 7033 without SIC in SS domains follows the classic model. Departure from linearity at high strains has also been observed from the curve of the σ against λ2 − λ−1 of ultrahigh molecular weight linear polyethylene. This deviation can be ascribed as a “Langevin effect” with SIC occurring due to highly entangled chains.10,47 Therefore, the present research clearly confirms that strain-induced crystallization results in deviation from the classic deformation mechanism for segmented copolymers and explains the microscopic mechanism. Another important finding is that WPA plays a key role in determining whether SIC can occur or not during deformation of PA1012− PTMO copolymers and further justifies or negates the validity of the Gaussian model.
■
CONCLUSIONS In the present investigation, lab-made, additive-free, segmented copolymers consisting of crystalline blocks of PA1012 and amorphous blocks of PTMO provide an opportunity to systematically explore the effect of the chain composition on the condensed state and on the macroscopic mechanical properties via in situ FTIR/WAXD/SAXS experiments. The Gaussian model of Haward and Thackray was applicable to the plastic PEBA without SIC at large strains, which was regarded as the necessary precondition. The strain-hardening modulus was calculated, bearing a linear relationship with WPA. Elastomeric PEBA samples with SIC at large strains negated the Gaussian assumption of amorphous coils. The onset of the Gaussian model failure was determined at the point, where the true stress linearly correlated with the PTMO regular packing absorption band intensity at 1009 cm−1 when λ reached 3.8. At this point, FTIR results suggested that the chain orientation of the hard segments just turned positive after reaching the I
DOI: 10.1021/acs.macromol.6b02747 Macromolecules XXXX, XXX, XXX−XXX
Article
Macromolecules
(9) Haward, R.; Thackray, G. The Use of a Mathematical Model to Describe Isothermal Stress-Strain Curves in Glassy Thermoplastics. Proc. R. Soc. London, Ser. A 1968, 302, 453−472. (10) Haward, R. Strain Hardening of Thermoplastics. Macromolecules 1993, 26, 5860−5869. (11) Haward, R. N. The Derivation of a Strain Hardening Modulus from True Stress-strain Curves for Thermoplastics. Polymer 1994, 35, 3858−3862. (12) Vorselaars, B.; Lyulin, A. V.; Michels, M. A. J. Microscopic Mechanisms of Strain Hardening in Glassy Polymers. Macromolecules 2009, 42, 5829−5842. (13) Men, Y.; Rieger, J.; Strobl, G. Role of the Entangled Amorphous Network in Tensile Deformation of Semicrystalline Polymers. Phys. Rev. Lett. 2003, 91, 955021−955024. (14) Hong, K.; Rastogi, A.; Strobl, G. A Model Treating Tensile Deformation of Semicrystalline Polymers: Quasi-Static Stress-Strain Relationship and Viscous Stress Determined for a Sample of Polyethylene. Macromolecules 2004, 37, 10165−10173. (15) Boulares, A.; Tessier, M.; Maréchal, E. Synthesis and Characterization of Poly(copolyethers-block-polyamides) II. Characterization and Properties of the Multiblock Copolymers. Polymer 2000, 41, 3561−3580. (16) Dong, X.; Zhu, P.; Wang, L.; Gao, Y.; Wang, D.; Liu, X. Polyamide 1012-Based Thermoplastic Elastomer and its Preparation Method. China Patent Application No. 201610069156.4, 2016. (17) Sheth, J.; Xu, J.; Wilkes, G. Solid State Structure−property Behavior of Semicrystalline Poly(ether-block-amide) PEBAX@ Thermoplastic Elastomers. Polymer 2003, 44, 743−756. (18) Bai, L.; Hong, Z.; Wang, D.; Li, J.; Wang, X.; Pan, G.; Li, L.; Li, X. Deformation-induced Phase Transitions of Polyamide 12 in its Elastomer Segmented Copolymers. Polymer 2010, 51, 5604−5611. (19) Kucera, L. R.; Brei, M. R.; Storey, R. F. Synthesis and Characterization of Polyisobutylene-b-polyamide Multi-block Copolymer Thermoplastic Elastomers. Polymer 2013, 54, 3796−3805. (20) Song, Y.; Yamamoto, H.; Nemoto, N. Segmental Orientations and Deformation Mechanism of Poly(ether-block-amide) Films. Macromolecules 2004, 37, 6219−6226. (21) McLean, R. S.; Sauer, B. B. Tapping-Mode AFM Studies Using Phase Detection for Resolution of Nanophases in Segmented Polyurethanes and Other Block Copolymers. Macromolecules 1997, 30, 8314−8317. (22) Hiss, R.; Hobeika, S.; Lynn, C.; Strobl, G. Network Stretching, Slip Processes, and Fragmentation of Crystallites during Uniaxial Drawing of Polyethylene and Related Copolymers. A Comparative Study. Macromolecules 1999, 32, 4390−4403. (23) Men, Y.; Strobl, G. Critical Strains in Poly(ε-caprolactone) and Blends with Poly(vinyl methyl ether) and Poly(styrene-co-acrylonitrile). Macromolecules 2003, 36, 1889−1898. (24) Liu, G.; Zheng, L.; Zhang, X.; Li, C.; Wang, D. Critical Stress for Crystal Transition in Poly(butylene succinate)-Based CrystallineAmorphous Multiblock Copolymers. Macromolecules 2014, 47, 7533− 7539. (25) Al-Hussein, M.; Strobl, G. Strain-Controlled Tensile Deformation Behavior of Isotactic Poly(1-butene) and its Ethylene Copolymers. Macromolecules 2002, 35, 8515−8520. (26) Na, B.; Lv, R.; Xu, W.; Yu, P.; Wang, K.; Fu, Q. Inverse Temperature Dependence of Strain Hardening in Ultrahigh Molecular Weight Polyethylene: Role of Lamellar Coupling and Entanglement Density. J. Phys. Chem. B 2007, 111, 13206−13210. (27) Niesten, M.; Harkema, S.; van der Heide, E.; Gaymans, R. J. Structural Changes of Segmented Copolyetheresteramides with Uniform Aramid Units Induced by Melting and Deformation. Polymer 2001, 42, 1131−1142. (28) Odarchenko, Y. I.; Anokhin, A. D.; Doblas, D.; Rosenthal, M.; Hernandez, J. J.; Vidal, L.; Sijbrandi, N. J.; Kimenai, A. J.; Mes, E.; Broos, R.; Bar, G.; Dijkstra, P. J.; Feijen, J.; Soloviev, M.; Ivanov, D. A. Primary Chemical Sequence Ultimately Determines Crystal Thickness in Segmented All-Aliphatic Copolymers. Macromolecules 2014, 47, 7890−7899.
negative minimum, and WAXD pattern showed that two pronounced spots ascribed to PTMO SIC in the vertical direction emerged. Meanwhile, Herman’s orientation function of the PTMO SIC exhibited a turning point, and the SAXS pattern advised a highly oriented fibrillar morphology formed after the spherulitic−fibrillar transition was completed. To the best of our knowledge, the present work is the first one to describe the preconditions for the application of the Gaussian model to the stress−strain curves of segmented copolymers and to explain why this classic model fails to treat the deformation of elastomeric PEBA from multiscale viewpoints explicitly and definitively. Therefore, our research provides deep insights into the deformation mechanism of semicrystalline segmented copolymers and into the complex interplay between two segments on the deformation behaviors of such copolymers.
■
ASSOCIATED CONTENT
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.6b02747. Figures S1−S16 and Table S1 (PDF)
■
AUTHOR INFORMATION
Corresponding Author
*(X.D.) Email
[email protected]. ORCID
Xia Dong: 0000-0002-6409-7011 Notes
The authors declare no competing financial interest.
■
ACKNOWLEDGMENTS This work was financially supported by the National Natural Science Foundation of China (Grants 21574140 and U1510207). The SSRF beamlines BL14B1 and BL16B1 are acknowledged for kindly providing the beam time and assistance. Detailed information about beamline BL14B1 can be found in ref 48.
■
REFERENCES
(1) Strobl, G. In The Physics of Polymers, 3rd ed.; Springer: Berlin, 2007; Chapter 10, p 443. (2) Peterlin, A. Molecular Model of Drawing Polyethylene and Polypropylene. J. Mater. Sci. 1971, 6, 490−508. (3) Jiang, Z.; Tang, Y.; Rieger, J.; Enderle, H. F.; Lilge, D.; Roth, S. V.; Gehrke, R.; Heckmann, W.; Men, Y. Two Lamellar to Fibrillar Transitions in the Tensile Deformation of High-density Polyethylene. Macromolecules 2010, 43, 4727−4732. (4) Farge, L.; Andre, S.; Meneau, F.; Dillet, J.; Cunat, C. A Common Multiscale Feature of the Deformation Mechanisms of a Semicrystalline Polymer. Macromolecules 2013, 46, 9659−9668. (5) Jiang, Z.; Tang, Y.; Men, Y.; Enderle, H.; Lilge, D.; Roth, S. V.; Gehrke, R.; Rieger, J. Structural Evolution of Tensile-deformed Highdensity Polyethylene during Annealing: Scanning Synchrotron Smallangle X-ray Scattering Study. Macromolecules 2007, 40, 7263−7269. (6) Bartczak, Z. Effect of Chain Entanglements on Plastic Deformation Behavior of Linear Polyethylene. Macromolecules 2005, 38, 7702−7713. (7) G’Sell, C.; Jonas, J. J. Yield and Transient Effects during the Plastic Deformation of Solid Polymers. J. Mater. Sci. 1981, 16, 1956− 1974. (8) G’Sell, C.; Hiver, J. M.; Dahoun, A.; Souahi, A. Video-controlled Tensile Testing of Polymers and Metals beyond the Necking Point. J. Mater. Sci. 1992, 27, 5031−5039. J
DOI: 10.1021/acs.macromol.6b02747 Macromolecules XXXX, XXX, XXX−XXX
Article
Macromolecules (29) Wang, L.; Dong, D.; Huang, M.; Wang, D. Transient Microstructure in Long Alkane Segment Polyamide: Deformation Mechanism and its Temperature Dependence. Polymer 2016, 97, 217− 225. (30) Jones, N. A.; Atkins, E. D. T.; Hill, M. J.; Cooper, S. J.; Franco, L. Polyamides with a Choice of Structure and Crystal Surface Chemistry. Studies of Chain-Folded Lamellae of Nylons 8 10 and 10 12 and Comparison with the Other 2N 2(N + 1) Nylons 4 6 and 6 8. Macromolecules 1997, 30, 3569−3578. (31) Yeh, F.; Hsiao, B. S.; Sauer, B. B.; Michel, S.; Siesler, H. W. InSitu Studies of Structure Development during Deformation of a Segmented Poly(urethane-urea) Elastomer. Macromolecules 2003, 36, 1940−1954. (32) Armstrong, S.; Freeman, B.; Hiltner, A.; Baer, E. Gas Permeability of Melt-processed Poly(ether-block-amide) Copolymers and the Effects of Orientation. Polymer 2012, 53, 1383−1392. (33) Sauer, B. B.; Mclean, R. S.; Brill, D. J.; Londono, D. J. Morphology and Orientation during the Deformation of Segmented Elastomers Studied with Small-Angle X-Ray Scattering and Atomic Force Microscopy. J. Polym. Sci., Part B: Polym. Phys. 2002, 40, 1727− 1740. (34) Bonart, R.; Mueller-Riederer, G. Modellvorstellungen zur Molekü lorientierung in Gedehnten Segmentierten Polyurethan Elastomeren. Colloid Polym. Sci. 1981, 259, 926−936. (35) Lin, S. B.; Hwang, K. S.; Tsay, S. Y.; Cooper, S. L. Segmental Orientation Studies of Polyether Polyurethane Block Copolymers with Different Hard Segment Lengths and Distributions. Colloid Polym. Sci. 1985, 263, 128−140. (36) Lee, H. S.; Park, H. D.; Cho, C. Domain and Segment Orientation Behavior of PBS-PTMG Segmented Block Copolymers. J. Appl. Polym. Sci. 2000, 77, 699−709. (37) Fakirov, S.; Denchev, Z.; Apostolov, A. A.; Stamm, M.; Fakirov, C. Morphological Characterization during Deformation of a Poly(ether ester) Thermoplastic Elastomer by Small-angle X-ray Scattering. Colloid Polym. Sci. 1994, 272, 1363−1372. (38) Cai, Z.; Bao, H.; Zhu, C.; Zhu, S.; Huang, F.; Shi, J.; Hu, J.; Zhou, Q. Structure Evolution of Polyamide 1212 during the Uniaxial Stretching Process: In Situ Synchrotron Wide-Angle X-ray Diffraction and Small-Angle X-ray Scattering Analysis. Ind. Eng. Chem. Res. 2016, 55, 7621−7627. (39) Finnigan, B.; Jack, K.; Campbell, K.; Halley, P.; Truss, R.; Casey, P.; Cookson, D.; King, S.; Martin, D. Segmented Polyurethane Nanocomposites: Impact of Controlled Particle Size Nanofillers on the Morphological Response to Uniaxial Deformation. Macromolecules 2005, 38, 7386−7396. (40) Ran, S.; Fang, D.; Zong, X.; Hsiao, B. S.; Chu, B.; Cunniff, P. M. Structural Changes during Deformation of Kevlar Fibers via On-line Synchrotron SAXS/WAXD Techniques. Polymer 2001, 42, 1601− 1612. (41) Strobl, G.; Schneider, M. Direct Evaluation of the Electron Density Correlation Function of Partially Crystalline Polymers. J. Polym. Sci., Polym. Phys. Ed. 1980, 18, 1343−1359. (42) Sijbrandi, N. J.; Kimenai, A. J.; Mes, E.; Broos, R.; Bar, G.; Rosenthal, M.; Odarchenko, Y. I.; Ivanov, D. A.; Dijkstra, P. J.; Feijen, J. Synthesis, Morphology, and Properties of Segmented Poly(ether amide)s with Uniform Oxalamide-Based Hard Segments. Macromolecules 2012, 45, 3948−3961. (43) Kamal, T.; Park, S.; Park, J.; Chang, Y. Structural Evolution of Poly(ether-b-amide12) Elastomers During the Uniaxial Stretching: An In situ Wide-Angle X-Ray Scattering Study. Macromol. Res. 2012, 20, 725−731. (44) Konyukhova, E. V.; Buzin, A. I.; Godovsky, Y. K. Melting of Polyether Block Amide (Pebax): the Effect of Stretching. Thermochim. Acta 2002, 391, 271−277. (45) Wang, D.; Shao, C.; Zhao, B.; Bai, L.; Wang, X.; Yan, T.; Li, J.; Pan, G.; Li, L. Deformation-Induced Phase Transitions of Polyamide 12 at Different Temperatures: An In Situ Wide-Angle X-ray Scattering Study. Macromolecules 2010, 43, 2406−2412.
(46) Kamal, T.; Park, S.; Choi, M.; Chang, Y.; Chuang, W.; Jeng, U. An In-situ Simultaneous SAXS and WAXS Survey of PEBAX Nanocomposites Reinforced with Organoclay and POSS during Uniaxial Deformation. Polymer 2012, 53, 3360−3367. (47) Coates, P. D.; Ward, I. M. Neck Profiles in Drawn Linear Polyethylene. J. Mater. Sci. 1980, 15, 2897−2914. (48) Yang, T.; Wen, W.; Yin, G.; Li, X.; Gao, M.; Gu, Y.; Li, L. Introduction of the X-ray Diffraction Beamline of SSRF. Nucl. Sci. Technol. 2015, 2, 1−5.
K
DOI: 10.1021/acs.macromol.6b02747 Macromolecules XXXX, XXX, XXX−XXX
