Mechanism of Significant Improvement of Large Strain Elasticity in Soft

Article Cite This: Ind. Eng. Chem. Res. XXXX, XXX, XXX−XXX

pubs.acs.org/IECR

Mechanism of Significant Improvement of Large Strain Elasticity in Soft Propylene−Ethylene Random Copolymer via Blending with Hard Propylene−Ethylene Coplymer Jiayi Zhao,†,‡ Xiao Yang,†,§ Yingying Sun,∥ and Yongfeng Men*,†,‡ †

State Key Laboratory of Polymer Physics and Chemistry, Changchun Institute of Applied Chemistry, Chinese Academy of Sciences, Renmin Street 5625, 130022 Changchun, P.R. China ‡ University of Chinese Academy of Sciences, No. 19(A) Yuquan Road, Shijingshan District, 100049 Beijing, P.R. China § University of Science and Technology of China, No. 96, JinZhai Road, Baohe District, 230026 Hefei, P.R. China ∥ ExxonMobil Asia Pacific Research & Development Co., Ltd., 1099 Zixing Road, Minhang District, 200241 Shanghai, P.R. China ABSTRACT: The deformation behavior of propylene−ethylene copolymers with 23 mol % of ethylene (iPPcoE23) and its blend (90/10 w/w) (iPPcoE23b) with propylene−ethylene copolymer (6 mol % ethylene) (iPPcoE6) at room temperature was studied by in situ small-angle X-ray scattering and wide-angle X-ray diffraction techniques. The 10 wt % copolymer with 6 mol % ethylene counits dispersed as small domains in the matrix of iPPcoE23. During stretching, longer and thinner fibrils and better elasticity properties were shown in the blend iPPcoE23b at large strain regime. The higher elasticity for iPPcoE23b at large strain regime was ascribed to the weak ability of strain-induced crystallization due to cocrystallization of the crystallizable long chain sequences of iPPcoE23 with the high crystalline iPPcoE6 during cooling, leaving fewer long crystallizable sequences in the matrix in iPPcoE23b. In addition, the phase-separated domains contributed stronger network modulus in the late stage of stretching compared with that of the recrystallized network of iPPcoE23.

■

INTRODUCTION

The crystallization and mechanical behavior of random propylene−ethylene copolymers, which are mostly used as thermoplastic elastomers, have been investigated in previous research. The higher concentration of ethylene counits results in the decrease of melting and crystallization temperature.16 Especially the γ-form, which can be produced only under high pressure for isotactic polypropylene (iPP),17−19 occupies much higher percentage for propylene−ethylene copolymer in the atmospheric pressure especially at high crystallization temperature20 or slow cooling rate.21 The high content of defects, including ethylene sequences22 and stereo- and regio-defects,23 make producing the nonparallel structure with low energy at an advantage during crystallization.20 The morphology of random propylene−ethylene copolymer with rich γ-form24 is composed of lamellae grown in the radical axis of the spherulite and short lamellae tilted to the main axis. The research on the deformation behavior of random propylene−ethylene copolymers is concentrated on both the recovery properties and the polymorphism transition during tensile tests. Hsiao et al.25

General plastics such as polyethylene and polypropylene are widely utilized in our daily life due to the high yields and low price. An understanding of the deformation mechanism of the semicrystalline polymers is considered as one of the significant subjects for industry. The isotropic lamellar structure transformed into highly oriented fibrils as well as newly generated oriented lamellae along the stretching direction under the stress, as observed by various techniques.1−3 A large amount of research has concluded that both the slips, including crystallographic fine slip and intralamellar mosaic blocks slip,4,5 and melting (disaggregation)−recrystallization6−8 are considered to be the deformation mechanism during tensile tests. The slippage of blocks occurred at small deformation, followed by melting (disaggregation)−recrystallization process which was controlled by the interaction between the crystalline lamellae and the entangled network.7,9−11 In addition, step-cycle tests,9 which can effectively separate the total strain into plastic and elastic parts, assist in judging the critical point C at which the disaggregation−recrystallization process for lamellar blocks takes place. Step-cycle tests are also widely used12−15 in investigating the recovery properties of polymer with high elasticity. © XXXX American Chemical Society

Received: Revised: Accepted: Published: A

January 15, 2018 March 20, 2018 March 22, 2018 March 22, 2018 DOI: 10.1021/acs.iecr.8b00194 Ind. Eng. Chem. Res. XXXX, XXX, XXX−XXX

Article

Industrial & Engineering Chemistry Research Table 1. Samples Information sample iPPcoE23 iPPcoE23b iPPcoE6 a

MFR (g/10 min)

Mw (g/mol)a

Mn (g/mol)a

Mw/Mna

stereo (mol %)

regio (mol %)

total (mol %)

3 3 8

2.2 × 10 2.3 × 105 2.1 × 105

8.8 × 10 8.9 × 104 9.5 × 104

2.56 2.59 2.20

0.3 0.3 0.3

0.2 0.2 0.2

23.5 22.5 6.5

5

4

The molecular weight and Mw/Mn of samples were obtained by gel permeation chromatography (GPC) in 1,2,4-trichlorobenzene.

immiscible if the difference in ethylene content was larger than 20 mol %. The upper critical solution temperature (UCST) behavior of the blend was clearly proved by the temperature dependence of the morphology. McNally et al.42 investigated the properties of blends of iPP and ethylene−octene copolymer (EOC). The scanning electron microscopy (SEM) and dynamic mechanical thermal analysis (DMTA) results both showed PP and EOC to be partially miscible for blends having 10 wt % EOC or less. Janani et al.43 studied the miscibility of blends of iPP and homogeneous iso-propylene-1-hexene copolymers with 11 and 21 mol % of 1-hexene. The critical value of 1-hexene for melt miscibility was 11 mol % at 0.1 MPa according to experimental PVT data. Thus, phase separation of the blend cannot be ignored in the investigation of the structure−property relationship. In this work, the structure evolution of random propylene− ethylene copolymer with ethylene counits content of 23 mol % (iPPcoE23) and its blend (90/10 w/w) (iPPcoE23b) with propylene−ethylene copolymer (6 mol % ethylene) (iPPcoE6) was studied by in situ small-angle X-ray scattering (SAXS) and (WAXD) wide-angle X-ray diffraction techniques at room temperature. The blend sample showed significant improvement in large strain elasticity, which can be attributed to the structural features presented during stretching.

found that strain-induced network was the reason for elastic behavior in cyclic deformation for propylene−ethylene copolymers, which was much like the deformation mechanism of the vulcanized rubber. Hild et al.26 proposed that small crystalline domains, which were regarded as physical cross-links for the amorphous matrix, were the reason for the high elasticity for the low crystalline polypropylene. In addition to the above study on the elasticity mechanism, polymorphism transition during stretching of crystals in γ-form has been investigated in recent years. During stretching, the c-axis of nonparallel structure of γ-form tilted at an angle to the fiber axis, which was called cross-β orientation mode.27−29 The stretching temperature29−31 and concentration of defects32,33 both had a significant influence on the deformation behavior during stretching. The research on random propylene−ethylene copolymers with high content of ethylene counits is rarely observed. De Rosa et al.34 proposed that the origin of the elasticity for poorly isotactic polypropylene (iPP) with very low crystallinity was the physical network by the high degree of entanglements with defects. Alamo et al.22 investigated the crystallization behavior of random propylene−ethylene copolymers in a range of ethylene up to 21 mol % and concluded that segmental mobility was necessary in developing γ-form. Besides ethylene comonomers, the influence of high content of other kinds of comonomers35 has been studied. According to the previous research, the participation of 1-butylene comonomer into propylene unit cell is in the highest level among all the comonomers. In samples with butene concentration of 25−30 mol %, chains crystallized only into α-form, which transformed into trigonal form of iPP during stretching.36 Ten mol % 1octene counits in random propylene-1-octene copolymers led to high elasticity properties because the stress was concentrated in the amorphous matrix and lateral size of lamellar decreased slightly.37 Increasing content of 1-hexene counits in propylene1-hexene copolymers resulted in the decrease of modulus, yield stress, and draw ratios.38 The microstructure of polyolefin blends, including the morphology and miscibility, has been studied by various techniques. Nitta et al.39 observed the significant difference between the morphology of the blend composed of iPP and ethylene-1-butene copolymers with different content of ethylene-1-butene, and the incompatible system displayed segregation at the interface between the two phases in the low strain region. Seki et al.40 studied the miscibility of blends of iPP and ethylene−propylene random copolymers by small-angle neutron scattering (SANS). SANS results proved that whether the ethylene counits were 19 or 47 mol %, the blends were in homogeneous one-phase mixture. The Flory−Huggins interaction parameter did not change within the temperature range. Kamdar et al.41 investigated the miscibility of homogeneous propylene−ethylene blends. Polymers with up to 30 mol % ethylene were blended in pairs to investigate the difference of comonomer content. Copolymers were miscible if the difference in ethylene content was less than 18 mol % and

■

EXPERIMENTAL SECTION The samples iPPcoE23 and iPPcoE23b produced using metallocene catalyst was kindly provided by ExxonMobil Asia Pacific Research and Development Co., Ltd. iPPcoE23 is a random propylene−ethylene copolymer with 23 mol % ethylene counits. iPPcoE23b is the blend of iPPcoE23 and another kind of propylene−ethylene copolymer with 6 mol % of ethylene (90/10 w/w). The melt flow rate (MFR) and molecular weight of all the samples are shown in Table 1. The pellets were first compression-molded at 180 °C into films with a thickness of 0.8 mm and held for 5 min to erase the thermal history. The films were stored at room temperature for two weeks before tensile tests. The tensile bars with dimensions of 10 × 5 × 0.8 mm3 were cut off by a punch. The tensile tests were carried out using a portable tensile testing machine (TST 350, Linkam, UK) at room temperature. Optical photo images of the sample were obtained to measure the strain of the sample during tensile tests. The Hencky strain εH was used as a basic quantity of the extension, which was defined as εH = 2ln

b0 b

where b0 and b are the widths of undeformed and deformed area, respectively. In step-cycle experiment, the sample was stretched step-by-step at a constant speed. After each step, the moving direction of the clamps was inversed to contract the sample until a stress of zero was achieved. Then, the sample was stretched again at this given speed, until it rereached the point at which it left the initial curve. A stepwise stretching of B

DOI: 10.1021/acs.iecr.8b00194 Ind. Eng. Chem. Res. XXXX, XXX, XXX−XXX

Article

Industrial & Engineering Chemistry Research

Figure 1. SEM images of cross-section of iPPcoE23 (a), iPPcoE23b (b), and iPPcoE23b (c) after being heated to 70 °C. All of the samples were etched before experiments.

Figure 2. Left: DSC heating curves for iPPcoE6, iPPcoE23, and iPPcoE23b stored at room temperature for two weeks (heating rate: 10 K/min). Right: tan δ as a function of temperature by DMTA experiments.

acid and then washed in hydrogen peroxide, distilled water, and acetone, respectively. The etchant helps remove the amorphous phase in semicrystalline polymers to reveal lamellar detail in surfaces. The cross section of the etched samples were sputtered with gold and examined on a field emission scanning electron microscope (FE-SEM, S-4800, Hitachi).

the sample with loading−unloading cycles was combined in the process. The imposed strains were therefore decomposed into a quasi-elastic (recoverable) part and a quasi-plastic (irreversible) part. In situ SAXS measurements during step-cycle experiments at 20 μm/s were carried out at the beamline BL16B1, SSRF, Shanghai. The sample-to-detector distance was 1905 mm, and the wavelength of X-ray was 0.124 nm. A 2D MAR CCD X-ray detector with a resolution of 2048 × 2048 (pixel size = 80 μm) was used to acquire 2D-SAXS patterns. The X-ray beam was positioned at the middle of the horizontally placed sample strip. The SAXS pattern was collected within 80 s and background subtracted. In situ WAXD measurements were performed at the beamline 1W2A, BSRF, Beijing. The wavelength of X-ray radiation was 0.154 nm, and the sample-to-detector distance was 86.5 mm. 2D WAXD images were recorded by a MAR CCD detector (2048 pixels ×2048 pixels, pixel size = 80 μm). The WAXD patterns were collected within 15 s and background subtracted. The DSC measurements were carried out with a DSC1 Star System (Mettler Toledo Swiss) which had been calibrated for temperature and melting enthalpy by using indium as a standard under a N2 atmosphere with a heating rate of 10 K/ min. The samples were heated at a constant rate of 10 K/min. Dynamic mechanical thermal analysis (DMTA) was performed on a Q800 DMTA (TA Instruments) in a tension mode with a frequency of 1 Hz and a heating rate of 3 K/min from −60 to 100 °C. The morphologies of microstructures were measured by SEM. The films were given a sharp notch, immersed in liquid nitrogen for 5 min, and immediately broken by hand. The fractured surface was etched for an hour in a 7% per volume solution of potassium permanganate in concentrated sulfuric

■

RESULTS AND DISCUSSION The SEM images in Figure 1 indicated a phase-separated morphology in iPPcoE23b due to the immiscibility between iPPcoE23 and iPPcoE6. All the components in iPPcoE23 are homogeneous, as was shown in Figure 1a. Figure 1b shows that some small domains with the size of about 50 μm dispersed in the matrix. After the iPPcoE23b was heated to 70 °C, part of matrix was melted, as shown in Figure 1c. Combined with Figures 1b and c, the components of the domain and matrix in the blend are iPPcoE6 and iPPcoE23, respectively. The high content of ethylene counits for samples iPPcoE23 and iPPcoE23b affects the crystallization and melting behavior. The differential scanning calorimetry (DSC) and dynamic mechanical thermal analysis (DMTA) curves of iPPcoE6, iPPcoE23, and their blend iPPcoE23b were shown in Figure 2. As shown in the melting curve of the blend iPPcoE23b, the peak located at around 50 °C corresponds to the crystals in copolymer iPPcoE23, while the other one at 110 °C corresponds to the crystals of iPPcoE6. In addition, the melting temperature of iPPcoE23b at low temperature side is lower than that of the iPPcoE23, indicating a cocrystallization of long regular propylene sequences in iPPcoE23 with the iPPcoE6. Only in such a case can the iPPcoE23b system show a crystallization and melting temperature lower than those of iPPcoE23. However, there is only one peak in the curve of tan δ as a function of temperature, shown in the right image of Figure 2. The glass transition temperature, Tg, for iPPcoE23 and C


Article


Figure 3. (a) True stress−strain curves of iPPcoE23 and iPPcoE23b at room temperature with error bar. (b) Variation of plastic (εH,p) and elastic (εH,e) parts as a function of total strain (εH) of iPPcoE23 and iPPcoE23b.

Figure 4. Selected 2D SAXS patterns of both stretched and recovered state during step-cycle tests. The stress was applied in the horizontal direction.

iPPcoE23b is −27.4 and −23.8 °C, respectively. The enhancement of Tg of the blend iPPcoE23b with respect to the Tg of the copolymer iPPcoE23 is ascribed to the thermally induced internal stress resulting from differential volume contraction of the two phases during cooling from the melt.44,45 The possible reason for one Tg in the phase-separated blend is that the amorphous phase of the two components is partly compatible. The DMTA curves of iPPcoE6 included two peaks. The Tg of iPPcoE6 was located at −10 °C, which was due to the higher crystallinity. The other peak which was located at 30 °C represented the α-relaxation. In this work, all the tensile tests were performed at room temperature. Figure 3a shows the true stress−strain curves. Both samples show similar mechanical response upon stretching with nearly overlapping behavior at small to moderate strain and only slight stronger strain-hardening for the blend iPPcoE23b at large strain regime. The resultant similar mechanical behavior for the two samples seems surprising but has its structural origin. As was discussed above, after a higher crystallinity iPPcoE6 was blended to the iPPcoE23, long crystallizable sequences previously in the iPPcoE23 system can be sucked into the dispersed iPPcoE6 domain, forming a kind of cocrystallization structure. As a result, the matrix became less crystalline and thus softer. Therefore, the inclusion of iPPcoE6 in the system of iPPcoE23 has two opposite effects: (1) reinforcement due to high

crystallinity, and (2) reducing matrix crystallinity. The two effects work together, resulting in the current situation of apparently similar mechanical behavior, especially at small strain regime. The relationship between plastic (εH,p) strain, elastic (εH,e) strain, and total strain (εH) of two samples is shown in Figure 3b. Similar to other semicrystalline polymers with higher crystallinity,13,15,46,47 the elastic strain increased rapidly at the beginning of the stretching, while the plastic strain remained low at small deformation. For iPPcoE23, the plastic strain increased significantly at a critical strain of 0.8. At large strain, the recovery property for iPPcoE23b was much better than that of iPPcoE23. Thus, it is necessary to investigate the relationship between the microstructure of the blend iPPcoE23b and the mechanical behavior, especially better elasticity properties at large strain. Figure 4 shows the evolution of 2D SAXS patterns for both stretched and recovered state. The stretching direction is horizontal. Before stretching, the initial SAXS pattern was isotropic. A streak in the vertical direction was observed at large strain for both samples. The presence of cavitation, which led to very strong scattering,48,49 was excluded from the origins of the long and thin streak in the vertical direction. Actually, it was caused by the scattering of fibrils12,50 along the stretching direction. Besides the scattering of fibrils in the vertical direction, there were also scattering from lamellar structure along the stretching direction and a kind of X-shaped structure D


Article


Figure 5. Selected integrated 1D log−log intensity distribution profiles as a function of scattering vector q along the stretching direction for iPPcoE23.

Figure 6. Selected integrated 1D log−log intensity distribution profiles as a function of scattering vector q along the stretching direction for iPPcoE23b.

the decrease of the thickness of films during stretching and decrease of density contrast between crystalline and amorphous region. At large strain, the lamellae structure became more ordered due to stress-induced crystallization. As shown from the broad scattering peak, the oriented structure was composed of two parts with different lamellar long spacing either stretched or recovered for iPPcoE23 and iPPcoE23b. The intensity of both the two scattering peaks continued increasing with the strain, which was very different from the situation that the intensity of original lamellae decreased and newly generated lamellae increased, shown in many other cases.7,52,53 The most probable reason was that two kinds of new lamellae with different long spacing were generated due to the stress-induced crystallization during stretching. The length of fibrils was calculated using Ruland’s method,54 and the radius was calculated from the length of the streak.12 Ruland’s method was used to calculate the length of the fibrils, which was described as follows:

distributed symmetrically with respect to the fibrils. The lateral size of the lamellae is inversely proportional to the intensity distribution Δq of the scattering pattern perpendicular to the stretching direction. As shown in Figure 4, intensity distribution Δq of the scattering pattern perpendicular to the stretching direction was larger for iPPcoE23b. Clearly, the lateral size of lamellae for iPPcoE23b was smaller than that of iPPcoE23. The X-shaped structure usually showed up accompanied by fibrils at the critical strain. It always tilted at around 17° to the vertical direction in stretched state. After the stress was removed, the tilted structure moved closer to the horizontal direction because the angle between the tilted structure and fibrils was around 45°. At large strain, the angle between the tilted structure and fibrils slightly increased both for stretched and recovered state. Unlike the stable fibrils which existed until the end of stretching, the scattering intensity of tilted structure decreased as well as the boundary between the tilted structure and lamellae along the horizontal direction became blurring in stretched state. All the data confirmed that the imperfect oriented structure was mostly destructed, and more sequences in the structure were forced to arrange along the stretching direction under stress. The tilted structure was probably caused by the aggregation of defects in the network.51 The integrated 1D SAXS curves along the horizontal direction at different strain for iPPcoE23 and iPPcoE23b are shown in Figures 5 and 6. The sample consisted of weakly coupled crystalline blocks as there was only a very weak scattering peak in the 1D SAXS profiles at small strain. The intensity decreased at small deformation, which was caused by

Bobs =

1 2π + BΦ lc q

where Bobs, lc, and BΦ represent the integral breadth, the length, and the misorientation of the fibril structure. The length of fibril refers to the long axis size of the long and thin streak. First, the azimuthal integrated curves at several q values were obtained. The integral breadth Bobs at each q equals to the integral area of one azimuthal intensity curve divided by the corresponding q. Finally, the fitted slope of the plotted line was 1 . For the calculation of the radius of fibrils, the peak intensity lc

E


Article


Figure 7. Evolution of length (a) and radius (b) of the crystalline fibrils as a function of total strain (εH).

as a function of q of azimuthal scans done at increasing q values was plotted. In the log−log scale, the intensity drops suddenly at q*. The radius of fibrils equals 2π . The evolution of length q*

and radius of the crystalline fibrils is shown in Figure 7. The arrow makes a connection between stretched state and corresponding recovered state at much lower strain. The sample was isotropic before stretching. The fibrils formed only at large strain during stretching. The streak in the vertical direction of 2D SAXS patterns, which represented the scattering of fibrils, appeared only at strain of 0.8; thus, the data at small strain were not used. The length of fibrils increased and radius decreased with strain under stress. The radius returned to the initial level, and the length of fibrils can hardly be calculated due to the low orientation in recovered state. The data confirmed that the stability for fibrils in this case was very poor. Previous research50 found that fibrils were composed of tightly aggregated polymer chain bundles with entanglements, and lamellar blocks played a role in restricting the fibrils in a confined place. In our case, both the evolution of lamellar structure oriented along the horizontal direction and the tilted structure near the fibrils during stretching indicated that the lamellar blocks were easily recovered, which cannot effectively protect the stability of the fibrils. Compared with iPPcoE23, the fibrils in the blend iPPcoE23b were longer and thinner because the structure composed of long crystallizable chains was not destroyed at small strain, and only the limited short chains were stretched to a large extent. To investigate the structure−properties relationship on the length scale of the crystal unit cell, selected 2D WAXD patterns during the stretching process are given in Figure 8. The patterns revealed that the amorphous chains accounted for a large part in the network even in highly oriented state at large strain. In addition, the diffraction of the branching structure was found near the horizontal direction resulting from the “crosshatch” lamellar structure55,56 generated during stretching. The patterns were almost isotropic after removing the load even at high strain. Detailed discussions on the diffractions of crystal planes are shown in the following section. Figures 9 and 10 show the evolution of WAXD curves as a function of total strain for iPPcoE23 and iPPcoE23b. It is difficult to recognize the crystal polymorph of the sample because the crystallinity is relatively low and the crystallites are full of defects due to the large amount of included ethylene sequences.22 The crystallinity for isotropic iPPcoE23b was a little higher than that of iPPcoE23 because the 10 wt %

Figure 8. Selected 2D WAXD patterns of both stretched and recovered state during step-cycle tests. The stress is applied in the horizontal direction.

copolymer with 6 mol % ethylene counits has contributed part of crystallinity. The azimuthal integrated intensity of crystallographic plane located at 2θ of 14.25° and 16.97° (marked in Figures 9 and 10) is plotted in Figures 11 and 12. For both samples at high strain, the cross-hatched daughter−parent structure was shown in the azimuthal distribution curves at 2θ of 14.25°, indicating the diffraction was from the oriented (110) crystal plane. The cross-hatch structure was rather stable as the azimuthal angle of the branching at 172° remained the same even after recovery. Three-peak profiles were observed in Figures 11c and d, which was similar to the azimuthal intensity distribution of (008) plane of γ-form29,57 with a second preferred orientation at an angle to the fiber axis but not the orientation of (040) plane of α-form even though the diffractions of the two crystal planes almost located at the same position. Thus, it is possible to conclude that the few original crystals in iPPcoE23 before stretching are “γ-form-like” crystals. On the other hand, the crystals in iPPcoE23b were “α-form-like” crystals because the orientation mode of crystal plane at 16.97° in Figure 12c proved that this plane was (040) plane of α-form. The polypropylene copolymers with high content of counits always developed into α-form during stretching at large strain.32 The ratio of daughter lamellae in the meridian is calculated using the method suggested by Fujiyama,58 which is defined as ratio = F

2[B] 2[B] + [P] DOI: 10.1021/acs.iecr.8b00194 Ind. Eng. Chem. Res. XXXX, XXX, XXX−XXX

Article


Figure 9. Selected 1D integrated WAXD intensity curves of iPPcoE23.

Figure 10. Selected 1D integrated WAXD intensity curves of iPPcoE23b.

induced crystallization. After releasing the load, some of the crystals melted while others persisted in recovered state. In the case of partly compatible blend iPPcoE23b, long crystallizable chains in 10 wt % iPPcoE6 contributed part of the crystallinity, as shown in DSC curves of Figure 2. The total crystallinity remained almost the same at small strain and increased slowly even after further stretching. Figure 15 shows the orientation degree parameter of the amorphous halo for two samples during stretching. The orientation degree parameter would be −0.5 if all the chains orientated along the stretching direction as we take the stretching direction as the reference direction and the scattering of the amorphous phase concentrates perpendicularly to the stretching direction when perfectly oriented. The data indicated that it was very difficult for large amount of amorphous phase to be oriented for both samples. The continuous increase in crystallinity for iPPcoE23 was due to the stronger ability of strain-induced crystallization, while the orientation proceeded much slower for iPPcoE23b at large strain which was caused by less-uniform distribution of stress in the partly compatible system.60 Moreover, as was shown by the DSC data in Figure 2, long crystallizable chain sequences in iPPcoE23b were mostly incorporated into the high crystallinity iPPcoE6 domains during the crystallization at high temperature. Thus, the matrix of iPPcoE23b presents ability for strain-induced crystallization behavior much lower than that for iPPcoE23. The weak ability of strain-induced crystallization of iPPcoE23b effectively maintained the elasticity even at large strain. However, the stress−strain curve for iPPcoE23b in Figure 3 displayed higher

where [P] is the area around the azimuthal angle of 90° and [B] represents the area of branching structure around 8° and 172° for crystal plane located at 14.25°. Because the azimuthal intensity curve was symmetric, only the data from 0° to 180° were used in calculation. As shown in Figure 13, the ratio of daughter lamellae was much lower in iPPcoE23b at the same strain even though the crystals behaved like α-form during stretching. The reason was that the smaller lateral size of lamellae was not extensive enough for the formation of crosshatching structure for the blend. The enhancement of ratio of daughter lamellae after strain of 0.9 confirmed the growth of daughter lamellae after fibrillation during stretching. However, the fragmentation of daughter lamellae at small strain for iPP with low stereoregularity was reported in previous research.59 The high content of ethylene counits regarded as the defects in the amorphous phase of the network protected the crystals during stretching to some extent. Figure 14 shows the change tendency of crystallinity during tensile tests for the two samples. Actually, the calculated crystallinity from WAXD data was the sum of oriented crystals and unoriented crystals. Clearly, the crystallinity of iPPcoE23 increased significantly during stretching compared with the blend iPPcoE23b. In addition, the decreased crystallinity in the recovered state described that this kind of crystallization was a reversible process. Because the sequence of crystallizable chains was very short and the numbers were limited in iPPcoE23 with high content of ethylene counits, chains with more defects were forced to arrange into the lattice at large strain. The short chains were highly stretched and acted as nuclei in the strainG


Article


Figure 11. Selected azimuthal integrated intensity of two planes at 14.25° (a, b) and 16.97° (c, d) for iPPcoE23.

Figure 12. Selected azimuthal integrated intensity of two planes at 14.25° (a, b) and 16.97° (c, d) for iPPcoE23b.

H


Article


tests, the initial modulus was almost the same for both samples with similar crystallinity. With the increasing of strain, the chains gradually crystallized for the copolymer iPPcoE23 under the load, while the network of iPPcoE23b was not destroyed at the same strain. The intact hard domains can provide modulus stronger than that of the recrystallized network at large strain. The schematic model of stretching behavior is shown in Figure 16.

Figure 13. Ratio of daughter lamellae as a function of total strain (εH).

Figure 16. Schematic model of stretching behavior for iPPcoE23 and iPPcoE23b.

■

CONCLUSIONS In summary, the structure evolution of the copolymer iPPcoE23 and the blend iPPcoE23b with 10 wt % of copolymer with 6 mol % ethylene counits during step-cycle tests at room temperature was investigated by in situ SAXS and WAXD measurements. The ability of stress-induced crystallization during stretching for iPPcoE23b was lower due to the uneven distribution of the stress and fewer long crystallizable chain sequences in the matrix. The elasticity properties for iPPcoE23b were much better than those for iPPcoE23.

■

Figure 14. Crystallinity as a function of total strain εH calculated from the WAXD data.

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. ORCID

Yongfeng Men: 0000-0003-3277-2227 Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS This work was supported by the National Natural Science Foundation of China (Grants 51525305 and 21134006) and ExxonMobil. We thank Dr. Robert Wittenbrink and Dr. Liang Li at ExxonMobil Asia Pacific Research and Development Company, Ltd. for helpful discussions, Prof. Zhonghua Wu and Dr. Guang Mo in BSRF, and Dr. Xiuhong Li and Dr. Jinyou Lin in SSRF for assistance during synchrotron X-ray scattering measurements.

■

Figure 15. Orientation degree parameter of the amorphous halo as a function of total strain εH.

REFERENCES

(1) Peterlin, A. Plastic Deformation of Polymers with Fibrous Structure. Colloid Polym. Sci. 1975, 253 (10), 809−823. (2) Corneliussen, R.; Peterlin, A. The Influence of Temperature on the Plastic Deformation of Polyethylene. Makromol. Chem. 1967, 105 (1), 193−203. (3) Peterlin, A.; Balta-Calleja, F. J. Diffraction Studies of Plastically Deformed Polyethylene. Colloid Polym. Sci. 1970, 242 (1−2), 1093− 1102.

strain-hardening modulus at large strain. Even though the phase-separated small domains cannot be used as additional physical cross-linkings between the amorphous phase and crystal blocks, the system of blend was still regarded as a composite filled by rigid particles.61 In the beginning of tensile I


Article

Industrial & Engineering Chemistry Research (4) Bowden, P. B.; Young, R. J. Deformation Mechanisms in Crystalline Polymers. J. Mater. Sci. 1974, 9 (12), 2034−2051. (5) Song, H.; Argon, A. S.; Cohen, R. E. Morphology of Highly Textured High-Density Polyethylene. Macromolecules 1990, 23 (3), 870−876. (6) Flory, P. J. Molecular Morphology in Semicrystalline Polymers. Nature 1978, 272 (5650), 226−229. (7) Men, Y.; Rieger, J.; Strobl, G. Role of the Entangled Amorphous Network in Tensile Deformation of Semicrystalline Polymers. Phys. Rev. Lett. 2003, 91 (9), 95502. (8) Wang, Y.; Jiang, Z.; Wu, Z.; Men, Y. Tensile Deformation of Polybutene-1 with Stable Form I at Elevated Temperature. Macromolecules 2013, 46 (2), 518−522. (9) 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 (13), 4390−4403. (10) Men, Y.; Strobl, G. CRITICAL STRAINS DETERMINING THE YIELD BEHAVIOR OF S-PP. J. Macromol. Sci., Part B: Phys. 2001, 40 (5), 775−796. (11) Men, Y.; Rieger, J.; Hong, K. Critical Strains in Tensile Deformed Polyamide 6 and 6/66 Copolymer. J. Polym. Sci., Part B: Polym. Phys. 2005, 43 (1), 87−96. (12) Deplace, F.; Scholz, A. K.; Fredrickson, G. H.; Kramer, E. J.; Shin, Y.-W.; Shimizu, F.; Zuo, F.; Rong, L.; Hsiao, B. S.; Coates, G. W. Tough and Elastic Thermoplastic Organogels and Elastomers Made of Semicrystalline Polyolefin-Based Block Copolymers. Macromolecules 2012, 45 (13), 5604−5618. (13) Sun, Y.; Fu, L.; Wu, Z.; Men, Y. Structural Evolution of Ethylene−Octene Copolymers upon Stretching and Unloading. Macromolecules 2013, 46 (3), 971−976. (14) Zhao, Y.; Zhu, Y.; Sui, G.; Chen, F.; Zhang, Q.; Fu, Q. The Effect of Hard Block Content on the Orientation and Mechanical Properties of Olefin Block Copolymer Films as Obtained via Melt Stretching. RSC Adv. 2015, 5 (100), 82535−82543. (15) Zhao, J.; Sun, Y.; Men, Y. Elasticity Reinforcement in Propylene−Ethylene Random Copolymer Stretched at Elevated Temperature in Large Deformation Regime. Macromolecules 2016, 49 (2), 609−615. (16) van der Burgt, F. P. T. J. Crystallization of Isotactic Polypropylene: The Influence of Stereo-Defects; Technische Universiteit Eindhoven: Eindhoven, Netherlands, 2002. (17) Jones, A. T.; Aizlewood, J. M.; Beckett, D. R. Crystalline Forms of Isotactic Polypropylene. Makromol. Chem. 1964, 75 (1), 134−158. (18) Kardos, J. L.; Christiansen, A. W.; Baer, E. Structure of PressureCrystallized Polypropylene. J. Polym. Sci., Part B: Polym. Phys. 1966, 4 (5), 777−788. (19) Angelloz, C.; Fulchiron, R.; Douillard, A.; Chabert, B.; Fillit, R.; Vautrin, A.; David, L. Crystallization of Isotactic Polypropylene under High Pressure (γ Phase). Macromolecules 2000, 33 (11), 4138−4145. (20) Alamo, R. G.; Kim, M.-H.; Galante, M. J.; Isasi, J. R.; Mandelkern, L. Structural and Kinetic Factors Governing the Formation of the γ Polymorph of Isotactic Polypropylene. Macromolecules 1999, 32 (12), 4050−4064. (21) Foresta, T.; Piccarolo, S.; Goldbeck-Wood, G. Competition between α and γ Phases in Isotactic Polypropylene: Effects of Ethylene Content and Nucleating Agents at Different Cooling Rates. Polymer 2001, 42 (3), 1167−1176. (22) Jeon, K.; Chiari, Y. L.; Alamo, R. G. Maximum Rate of Crystallization and Morphology of Random Propylene Ethylene Copolymers as a Function of Comonomer Content up to 21 Mol %. Macromolecules 2008, 41 (1), 95−108. (23) De Rosa, C.; Auriemma, F.; Resconi, L. Influence of Chain Microstructure on the Crystallization Kinetics of Metallocene-Made Isotactic Polypropylene. Macromolecules 2005, 38 (24), 10080−10088. (24) Hosier, I. L.; Alamo, R. G.; Lin, J. S. Lamellar Morphology of Random Metallocene Propylene Copolymers Studied by Atomic Force Microscopy. Polymer 2004, 45 (10), 3441−3455.

(25) Toki, S.; Sics, I.; Burger, C.; Fang, D.; Liu, L.; Hsiao, B. S.; Datta, S.; Tsou, A. H. Structure Evolution during Cyclic Deformation of an Elastic Propylene-Based Ethylene-Propylene Copolymer. Macromolecules 2006, 39 (10), 3588−3597. (26) Boger, A.; Imhof, C.; Heise, B.; Marti, O.; Troll, C.; Rieger, B.; Hild, S. Stress-Induced Changes in Microstructure of a LowCrystalline Polypropylene Investigated at Uniaxial Stretching. J. Appl. Polym. Sci. 2009, 112 (1), 188−199. (27) Geddes, A. J.; Parker, K. D.; Atkins, E. D. T.; Beighton, E. CrossB” Conformation in Proteins. J. Mol. Biol. 1968, 32 (2), 343−358. (28) Auriemma, F.; De Rosa, C. Stretching Isotactic Polypropylene: From “cross-B” to Crosshatches, from γ Form to α Form. Macromolecules 2006, 39 (22), 7635−7647. (29) Mao, Y.; Li, X.; Burger, C.; Hsiao, B. S.; Tsou, A. H. 2D WAXS/ SAXS Study on Isotactic Propylene-1-Butylene Random Copolymer Subjected to Uniaxial Stretching: The Influence of Temperature. Polymer 2013, 54 (4), 1432−1439. (30) Mao, Y.; Burger, C.; Li, X.; Hsiao, B. S.; Mehta, A. K.; Tsou, A. H. Time-Resolved Synchrotron X-Ray Scattering Study on Propylene−1-Butylene Random Copolymer Subjected to Uniaxial Stretching at High Temperatures. Macromolecules 2012, 45 (2), 951−961. (31) Auriemma, F.; De Rosa, C.; Di Girolamo, R.; Malafronte, A.; Scoti, M.; Mitchell, G. R.; Esposito, S. Deformation of Stereoirregular Isotactic Polypropylene across Length Scales. Influence of Temperature. Macromolecules 2017, 50 (7), 2856−2870. (32) De Rosa, C.; Auriemma, F. Structural−Mechanical Phase Diagram of Isotactic Polypropylene. J. Am. Chem. Soc. 2006, 128 (34), 11024−11025. (33) De Rosa, C.; Auriemma, F.; Ruiz de Ballesteros, O.; Resconi, L.; Camurati, I. Tailoring the Physical Properties of Isotactic Polypropylene through Incorporation of Comonomers and the Precise Control of Stereo- and Regioregularity by Metallocene Catalysts. Chem. Mater. 2007, 19 (21), 5122−5130. (34) De Rosa, C.; Auriemma, F.; Perretta, C. Structure and Properties of Elastomeric Polypropylene from C 2 and C 2 v -Symmetric Zirconocenes. The Origin of Crystallinity and Elastic Properties in Poorly Isotactic Polypropylene. Macromolecules 2004, 37 (18), 6843−6855. (35) Jeon, K.; Palza, H.; Quijada, R.; Alamo, R. G. Effect of Comonomer Type on the Crystallization Kinetics and Crystalline Structure of Random Isotactic Propylene 1-Alkene Copolymers. Polymer 2009, 50 (3), 832−844. (36) De Rosa, C.; Auriemma, F.; Vollaro, P.; Resconi, L.; Guidotti, S.; Camurati, I. Crystallization Behavior of Propylene−Butene Copolymers: The Trigonal Form of Isotactic Polypropylene and Form I of Isotactic Poly(1-Butene). Macromolecules 2011, 44 (3), 540−549. (37) Li, X.; Mao, Y.; Burger, C.; Che, J.; Hsiao, B. S.; Kulkarni, R. R.; Tsou, A. H. The Effect of Comonomer Content on Structure and Property Relationship of Propylene-1-Octene Copolymer during Uniaxial Stretching. Polymer 2013, 54 (17), 4545−4554. (38) Janani, H.; Alamo, R. G. Mechanical Properties and StressInduced Polymorphic Transformations of Blends of Random Isotactic Propylene-1-Hexene Copolymers. Polymer 2016, 102, 21−32. (39) Nitta, K.; Okamoto, K.; Yamaguchi, M. Mechanical Properties of Binary Blends of Polypropylene with Ethylene-E-Olefin Copolymer. Polymer 1998, 39 (1), 53−58. (40) Seki, M.; Nakano, H.; Yamauchi, S.; Suzuki, J.; Matsushita, Y. Miscibility of Isotactic Polypropylene/Ethylene−Propylene Random Copolymer Binary Blends. Macromolecules 1999, 32 (10), 3227−3234. (41) Kamdar, A. R.; Hu, Y. S.; Ansems, P.; Chum, S. P.; Hiltner, A.; Baer, E. Miscibility of Propylene - Ethylene Copolymer Blends. Macromolecules 2006, 39 (4), 1496−1506. (42) McNally, T.; McShane, P.; Nally, G. M.; Murphy, W. R.; Cook, M.; Miller, A. Rheology, Phase Morphology, Mechanical, Impact and Thermal Properties of Polypropylene/metallocene Catalysed Ethylene 1-Octene Copolymer Blends. Polymer 2002, 43 (13), 3785−3793. (43) Janani, H.; Alamo, R. G. Melt Miscibility of Blends of Isotactic Polypropylene and Homogeneous Iso-Propylene-1-Hexene Copolymers. J. Therm. Anal. Calorim. 2014, 116 (3), 1497−1508. J


Article

Industrial & Engineering Chemistry Research (44) Mäder, D.; Bruch, M.; Maier, R.; Stricker, F.; Mülhaupt, R. Glass Transition Temperature Depression of Elastomers Blended with Poly(propene)s of Different Stereoregularities. Macromolecules 1999, 32 (4), 1252−1259. (45) Mäder, D.; Thomann, Y.; Suhm, J.; Mülhaupt, R. Influence of Comonomer Incorporation on Morphology and Thermal and Mechanical Properties of Blends Based upon Isotactic MetallocenePolypropene and Random ethene/1-Butene Copolymers. J. Appl. Polym. Sci. 1999, 74 (4), 838−848. (46) Al-Hussein, M.; Strobl, G. Strain-Controlled Tensile Deformation Behavior of Isotactic Poly(1-Butene) and Its Ethylene Copolymers. Macromolecules 2002, 35 (22), 8515−8520. (47) Jiang, F.; Fang, C.; Zhang, J.; Wang, W.; Wang, Z. Triblock Copolymer Elastomers with Enhanced Mechanical Properties Synthesized by RAFT Polymerization and Subsequent Quaternization through Incorporation of a Comonomer with Imidazole Groups of about 2.0 Mass Percentage. Macromolecules 2017, 50 (16), 6218− 6226. (48) Liu, Y.; Truss, R. W. A Study of Tensile Yielding of Isotactic Polypropylene. J. Polym. Sci., Part B: Polym. Phys. 1994, 32 (12), 2037−2047. (49) Lu, Y.; Wang, Y.; Chen, R.; Zhao, J.; Jiang, Z.; Men, Y. Cavitation in Isotactic Polypropylene at Large Strains during Tensile Deformation at Elevated Temperatures. Macromolecules 2015, 48 (16), 5799−5806. (50) Mao, Y.; Li, X.; Burger, C.; Hsiao, B. S.; Mehta, A. K.; Tsou, A. H. Structure Development during Stretching and Heating of Isotactic Propylene−1-Butylene Random Copolymer: From Unit Cells to Lamellae. Macromolecules 2012, 45 (17), 7061−7071. (51) Tian, N.; Zhao, B.; Li, X.; Liu, Y.; Zhou, W.; Cui, K.; Liu, D.; Li, L. Confined Crystallization in End-Linked PEO Network under Uniaxial Extension. Polymer 2013, 54 (26), 7088−7093. (52) 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 (10), 4727−4732. (53) Lu, Y.; Chen, R.; Zhao, J.; Jiang, Z.; Men, Y. Stretching Temperature Dependency of Fibrillation Process in Isotactic Polypropylene. J. Phys. Chem. B 2017, 121 (28), 6969−6978. (54) Ruland, W. Small-Angle Scattering Studies on Carbonized Cellulose Fibers. J. Polym. Sci., Part C: Polym. Symp. 1969, 28 (1), 143−151. (55) Lotz, B.; Wittmann, J. C. The Molecular Origin of Lamellar Branching in the α (Monoclinic) Form of Isotactic Polypropylene. J. Polym. Sci., Part B: Polym. Phys. 1986, 24 (7), 1541−1558. (56) Bassett, D. C.; Olley, R. H. On the Lamellar Morphology of Isotactic Polypropylene Spherulites. Polymer 1984, 25 (7), 935−943. (57) Zhao, J.; Sun, Y.; Men, Y. Counits Content and Stretching Temperature-Dependent Critical Stress for Destruction of γ Crystals in Propylene−Ethylene Random Copolymers. ACS Omega 2017, 2 (10), 6896−6905. (58) Fujiyama, M.; Wakino, T.; Kawasaki, Y. Structure of Skin Layer in Injection-molded Polypropylene. J. Appl. Polym. Sci. 1988, 35 (1), 29−49. (59) Liu, Y.; Hong, Z.; Bai, L.; Tian, N.; Ma, Z.; Li, X.; Chen, L.; Hsiao, B. S.; Li, L. A Novel Way to Monitor the Sequential Destruction of Parent-Daughter Crystals in Isotactic Polypropylene under Uniaxial Tension. J. Mater. Sci. 2014, 49 (8), 3016−3024. (60) Liu, G.; Zhang, X.; Liu, Y.; Li, X.; Chen, H.; Walton, K.; Marchand, G.; Wang, D. Effect of Elastomer on Crystalline Transition and Deformation Behavior of Isotactic Polypropylene. Polymer 2013, 54 (4), 1440−1447. (61) Jiang, F.; Zhang, Y.; Fang, C.; Wang, Z.; Wang, Z. From Soft to Strong Elastomers: The Role of Additional Crosslinkings in Copolymer-Grafted Multiwalled Carbon Nanotube Composite Thermoplastic Elastomers. RSC Adv. 2014, 4 (104), 60079−60085.

K


Mechanism of Significant Improvement of Large Strain Elasticity in Soft

Recommend Documents