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Oct 7, 2015 - Propylene (PP) is a semicrystalline polymer, which has become one of the most used polymers to produce porous membranes.1−3...
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Rheology of PP/PP‑g‑MA and PP/PP‑g‑AA Blends and Incidence on Orientation and Crystalline Structure of Their Cast Films Amir Saffar,† Seyed H. Tabatabaei,† Pierre J. Carreau,† Abdellah Ajji,*,† and Musa R. Kamal‡ †

Research Center for High Performance Polymer and Composite Systems (CREPEC), Chemical Engineering Department, Polytechnique Montreal, P.O. Box 6079, Stn Centre-Ville, Montreal, Quebec H3C 3A7, Canada ‡ CREPEC, Department of Chemical Engineering, McGill University, 3610 University Street, Montreal, Quebec H3A 2B2, Canada ABSTRACT: The intrinsic hydrophobic character of polypropylene limits its performance in many porous membrane applications. This can be improved by hydrophilic modification of the film surface through blending with hydrophilic polymers. For producing the precursor films and, consequently, microporous membranes with the appropriate crystalline lamellar morphology, the polypropylene chains should preserve their elongated form to serve as initial nuclei for the later development of lamellar crystals. The relaxation time of the chains is the most important factor for this stage. Commercial maleic anhydride and acrylic acid grafted polypropylenes were melt blended with a polypropylene at different weight ratios. The results showed that the modifiers lowered the crystalline orientation of the blends as compared to neat polypropylene films. The effect of the modifier on the melt relaxation spectra of the blends was investigated, and a linear relationship was found between the characteristic relaxation time of the blends and their crystalline orientation function.

1. INTRODUCTION Propylene (PP) is a semicrystalline polymer, which has become one of the most used polymers to produce porous membranes.1−3 Solution/phase separation process is usually used to produce membranes. Porous membranes can be prepared using a dry process by the formation of a specific crystalline morphology in the films, followed by annealing and stretching (lamellae separation).2,4,5 If this method was developed 30 years ago for some semicrystalline polymers, most of the detailed information regarding processing remained proprietary and was not available to the scientific community. Three consecutive stages need to be carried out to obtain a satisfactory porous membrane using the stretching method: (1) producing a film containing an oriented shish-kebab structure, (2) annealing the precursor film to improve the crystalline structure, and, finally, (3) stretching the film at room temperature and then near the melting point of the polymer to create and enlarge the pores by lamellae separation, respectively.2,5−8 Despite its popularity as a membrane material, PP does not adhere to hydrophilic substrates because of its intrinsic hydrophobic nature that limits its applications in areas where this property is required.3,9−11 In our previous work,12 to produce microporous hydrophilic membranes, different contents of maleic anhydride grafted polypropylene (PP-g-MA) and acrylic acid grafted polypropylene (PP-g-AA) were melt blended with PP. It was shown that the addition of the modifiers enhanced the surface hydrophilicity of the membranes. Furthermore, it was indicated that the hydrophilic modifiers on the surface could interact with other hydrophilic groups.13 It was shown that the new secondary crystalline lamellae are formed in the PP-based precursor films during the annealing process.14 Following these, the pore structure as a function of stretching in the third step of the membrane formation was evaluated, and its impact on the membrane performance was reported.14 © 2015 American Chemical Society

In the stretching method, the important factor for producing microporous membranes is controlling crystalline structure of the films.2 In fact, the crystalline orientation factor, Hermans crystalline orientation, should be larger than 0.3.15 Applying the suitable processing conditions as well as selecting the right materials is the first step to control and generate the desired lamellar morphology. This structure is obtained during melt stretching by crystallization of the chains, which is commonly called flow-induced crystallization (FIC).5,16 The basic mechanisms that lead to FIC, from a physical viewpoint, are relatively well understood. The strain under flow markedly increases the nucleation rate and the crystal growth rate of polymer chains. Moreover, a different morphology is formed as compared to quiescent conditions.17,18 FIC is particularly relevant in the processes where high velocity gradients are involved, such as in blow molding, film casting, injection molding, and fiber spinning.19,20 In cast film, both elongational and shear fields are applied on the polymer melt.21,22 In the last few decades, there have been several investigations focused on FIC, most of which focused on shear flow. Many of the results have been performed using rheometer or shear cells in the low shear rate region.23−27 However, the effect of intense shear flow was studied by the research group of JaneschitzKriegl28−30 by using an extrusion die setup. The nonhomogeneous character of the flow field is a major drawback of this setup. Baert et al.18 used a sandwich-type shear cell, which imposed a homogeneous flow, and studied the transition to highly anisotropic structures in isotactic poly-1-butene (PB-1) crystallization for flow close to industrial processing conditions. Although numerous studies have been carried out on the effect Received: Revised: Accepted: Published: 10787

June 13, 2015 September 30, 2015 October 7, 2015 October 7, 2015 DOI: 10.1021/acs.iecr.5b02141 Ind. Eng. Chem. Res. 2015, 54, 10787−10795

Article

Industrial & Engineering Chemistry Research Table 1. Main Characteristics of the Neat Materials resin code

supplier

MFR (g/10 min)

nomenclature

ηοa (kPa s)

Tmb (°C)

Tcb (°C)

polypropylene PP5341E1 polypropylene PP4712E1 PP-g-MA (Bynel50E739) PP-g-AA (Polybond1002)

Exxon Mobil Exxon Mobil DuPont Chemtura

0.8 2.8 6 20

PP0.8 PP2.8 PPMA PPAA

45.6 16.3 7.5 3.1

161 160 141 161

115 114 102 123

Zero-shear viscosity values calculated using the area under the weighted relaxation spectrum curves, T = 190 °C. bThe melting point and the crystallization temperature were obtained using differential scanning calorimetry (DSC). a

of shear flow on the crystallization, very few works have systematically studied the role of pure extensional flows.31−33 This lack of investigation is due to the difficulties to measure the extensional material functions. The shear flow can be represented as the superposition of elongational and rotational flow. The rotational and elongational components in a simple shear flow are equal, such that the polymer cannot attain a stable fully stretched condition: the extensional component orients and stretches the molecules, whereas the rotational component causes a fluctuation in the extension.34 Therefore, shear flow is considered a weak flow that is less efficient than elongational flow to obtain high orientation or stretching.22,32,35 Termonia36 reported that the coil−stretch transition induced by elongational or shear flows in dilute polymer solutions occurred at a much lower flow rate in elongational flow. In this study, in preparing the precursor film via extrusion, the polymer chains are oriented by shear inside the die, then by elongation after the die through drawing in the melt state followed by rapid crystallization. The extended chains have a higher free energy, which can crystallize faster and form rownucleated fibril crystals to serve as nuclei for the later development of chain-folding lamellar crystals.16,37−39 The combination of the above two structures is commonly referred to as shish-kebab morphology, which was first reported by Pennings and Kiel40 for a polymeric solution. The processing parameters as well as the resin characteristics are the key factors for controlling the final morphological features of the extruded precursor films.19,21,26,38 Also, draw ratio, die temperature, die gap, rate and position of the air knife, and chill roll temperature are the major processing parameters that need to be optimized depending on the resin. It is obvious that during processing, the polymer chains should preserve their elongated form to be able to act as initial nuclei.37 On the basis of experimental evidence, a critical deformation or elongation rate needs to be exceeded to change the spherulitic structure into an oriented crystallization regime where fibrillar growth occurs.24,39,41 This critical deformation may depend on the characteristics of the polymer and may be lowered by addition of long and/or branched chains.42 It is worth mentioning that for highly oriented precursor films, the stresses involved for the formation of the membranes are mostly associated with the stretching of tie chains. However, for precursor films with low orientation levels, the stresses are spent mostly for reorienting the crystal blocks in the stretching direction.21,43 Hence, the purpose of this study is to determine how the blending of modifiers (PP-g-MA or PP-gAA) changes the rheological characteristics of PP films including their main characteristic time and, consequently, their crystalline orientation. Understanding the influence of the relaxation time of polymers on the crystalline structure can be used to apply suitable processing conditions and, consequently, prepare precursor films and porous membranes with desired properties.

2. EXPERIMENTAL SECTION 2.1. Materials. Two commercial linear polypropylenes (MFR values of 0.8 and 2.8 g/10 min under ASTM conditions of 230 °C and 2.16 kg) were selected as the matrixes for the microporous membranes. Both polymers were supplied by ExxonMobil Co. Two types of commercial amphiphilic polymers consisting of PP backbones and hydrophilic side chains, that is, PP-g-MA (with MA content lower than 1 wt %) and PP-g-AA (with AA content of 6 wt %, which was confirmed using X-ray photoelectron spectroscopic (XPS) measurements12), were used as modifiers. The PP-g-MA was supplied by DuPont Co., and the PP-g-AA was obtained from Chemtura Corp. Table 1 shows the main characteristics of the polymers. 2.2. Film Preparation. Precursor films from blends of PP and modifiers at different weight ratios were created using a 45 mm diameter Killion single screw extruder. The temperature profile was set in the range of 165−220 °C, and the screw speed was 12.5 rpm. The film was cooled by an air knife right at the exit of the die. To have the precursor films having a thickness of around 30 μm, the ratio of the roll speed to the extrudate velocity at die exit, or draw-down ratio, was set at 25. Also, the cast rolls temperature was constant at 50 °C, while its distance from the die exit was around 15 cm. 2.3. Rheological Characterization. Rheological measurements were performed using a parallel plate rheometer (AntonPaar MCR 301 stress controlled rheometer) with a gap size of 1 mm and plate diameter of 25 mm. The measurements were carried out at 190 °C under nitrogen flow to avoid thermal degradation. Time sweep tests were first done at a frequency of 0.628 rad/s for 1 h, to check for degradation during the tests. The complex viscosity and storage modulus were determined in the frequency range from 0.0628 to 628 rad/s, while the strain amplitude was set as 10% (in the linear viscoelastic regime). 2.4. Fourier Transform Infrared Spectroscopy (FTIR). The infrared spectra in the range from 4000 to 600 cm−1 were collected using a spectrum 65 FTIR spectrometer (Perkin Elmer) with a spectral resolution of 4 cm−1. The surface composition of the samples was characterized using the attenuated total reflectance (ATR) mode. The crystalline orientation was measured in the transmissionFTIR mode using a polarized radiation parallel and perpendicular to the machine direction (MD). The Hermans orientation function was determined as follows:44 D−1 f= (1) D+2 where D is defined as the ratio of the absorbencies parallel and perpendicular to machine direction at the specific vibration used. The absorption at 998 cm−1 can be used for calculating the crystalline orientation function, fc, of polypropylene. 2.5. X-ray Diffraction and Scattering. A Bruker D8 Discover equipped with a Hi-STAR two-dimensional area detector (Cu Kα radiation, λ = 1.542 Å), operating at a voltage of 40 kV 10788

DOI: 10.1021/acs.iecr.5b02141 Ind. Eng. Chem. Res. 2015, 54, 10787−10795

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could be related to the degradation of the PPAA during the tests, whereas the increase of G′ with time for the PPMA might be due to cross-linking and reaction in the sample as a result of the presence of the residual free maleic anhydride. As all of the following rheological data were collected within 20 min, the 3% increases of G′ for the neat PPs and PPMA and the 10% reduction for PPAA during that time period are considered acceptable. The normalized complex viscosity (i.e., η*/η0, where η0 is the zero-shear viscosity) of the neat polymers as a function of frequency is presented in Figure 2a. The zero-shear viscosities of the polymers obtained from the Carreau−Yasuda model (not shown here) show good agreement with those obtained using the area under the weighted relaxation spectrum curves (see Table 1 and discussed below). Those values are consistent with their MFRs, as presented in Table 1. As the molecular weight increases, polymer behavior becomes more shear-thinning, and the transition of Newtonian to the power-law region is observed at lower frequencies.2,6 Figure 2b shows the loss angle, δ (tan δ = G″/G′), as a function of frequency. It is clear that the loss angle decreases with increasing Mw. The larger elasticity of the PPs (lower loss angle) as compared to the modifiers could be attributed to the presence of longer chains. 3.2. Rheological Properties of the Blends. Figure 3 shows the complex shear viscosity as a function of frequency for the neat polymers and their blends with PPMA and PPAA. It can be seen that the presence of low molecular weight modifier decreases the complex shear viscosity and causes a slight reduction of the shear-thinning behavior as well. In addition, it is clearly observed that the viscosities of the blends fall in the intermediate range between those of their components. Relaxation of the chain happens when macromolecules reach conformations of lower energetic state, which depends on the interactions between the species.45 The weighted relaxation spectra were obtained from the dynamic moduli data and using the NLREG (nonlinear regularization) software, for quantitatively analyzing the change of melt relaxation by adding of modifiers.46 The relaxation spectra can show even small differences in molecular architecture, because it is sensitive to the molecular structure of polymers.45 Figure 4 presents the weighted relaxation spectra, while the range of frequencies covered during the test is indicated by the vertical dash lines. It can be observed that increasing modifier concentration decreases the number of entanglements and the weight-average molecular weight. Hence, the spectrum shape becomes narrower, and the characteristic (mean) relaxation time corresponding to the peaks in the curves, τc, is shifted to shorter times. The area under the curves, which is attributed to the zero-shear viscosity of the melt,2,47 enhances with molecular weight. Also, the positions of the peaks for the blends are intermediate to those for the neat components, while PP exhibits the largest characteristic relaxation time due to its larger molecular weight and broader area as compared to PPMA and PPAA. Furthermore, the results demonstrate a reasonable negative linear correlation between the mean relaxation time and modifier content (not shown). It should also be mentioned that the same trends were found for PP2.8 blends, and the data are not presented here for the sake of brevity. 3.3. Crystalline Orientation of the Precursor Films. SAXS was used to identify the morphologies of the precursor films made of the neat PP0.8 and its blends with PPMA (shown in our previous publication14). An equatorial streak and meridian maxima in the SAXS patterns, which is related to the

and current of 40 mA, was used for wide-angle X-ray diffraction (WAXD) measurements. The detector was fixed at a distance of 9.95 cm from the samples. The diffraction of a monochromatic X-ray beam by the crystallographic planes (hkl) of the crystalline phase is the basic principle of WAXD. As the sample is rotated with respect to the beam, the probability distribution of the orientation of the normal to the hkl plane can be found. The Hermans orientation function was calculated according to5 f=

(3 cos2 ϕ − 1) 2

(2)

where ϕ is the angle between the reference axes and the unit cell axes (a, b, and c). A Bruker AXS Nanostar system equipped with a microfocus copper anode at 45 kV/0.65 mA was used for small-angle X-ray scattering (SAXS) measurements. The distance between MONTAL OPTICS and a VANTEC 2000 2D detector and the samples was set at 107.2 mm. 2.6. Differential Scanning Calorimetry (DSC). A TA Instruments differential scanning calorimeter (DSC) Q 1000 was used to study the thermal properties of the specimens in the range of 50−220 °C with a heating/cooling rate of 10 °C/min.

3. RESULTS AND DISCUSSION We first present the rheological results regarding the neat polymers and their blends. Subsequently, the experimental results regarding the crystalline structure of the films are reported. The results clearly demonstrate the effect of the mean relaxation time of the polymers on the crystalline structure of the films. The correlation between the mean relaxation time and the crystalline orientation is discussed. Because the relaxation times of the polymer chains are affected by temperature,21 all of the rheological tests were done at the same temperature (190 °C). Also, the same processing condition was used for preparing all of the precursor films. 3.1. Characterization of the Neat Materials. Time sweep tests were first performed for 1 h at a frequency of 0.628 rad/s, to evaluate the thermal stability of the neat polymers. Figure 1 presents the normalized storage modulus (G′(t)/G′(t = 0)). For PPAA the normalized G′ decreases with time, while for the other polymers it increases. The reduction

Figure 1. Normalized storage modulus versus time for the neat polymers; T = 190 °C and ω = 0.628 rad/s. 10789

DOI: 10.1021/acs.iecr.5b02141 Ind. Eng. Chem. Res. 2015, 54, 10787−10795

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Figure 2. (a) Normalized complex viscosity and (b) loss angle versus frequency for the neat polymers; T = 190 °C.

Figure 3. Complex viscosity versus angular frequency for the neat PP0.8 and its blends with (a) PPMA and (b) PPAA; T = 190 °C.

anhydride and acrylic acid groups at different positions on the polypropylene chains of the neat modifiers were identified in ATR-FTIR tests.12 Furthermore, in comparison with the PP2.8 matrix, the PP0.8 blend shows larger crystalline orientation because of its higher molecular weight (Mw), which increases the number of fibrils or nuclei sites.12 The difference in the crystalline structure of the blends could be related to the configuration of tie chains and the orientation of the crystal blocks.4 So, to clearly describe the change of the crystalline structure, blends of PP0.8 with PPMA were checked using WAXD. The diffraction intensity profiles as well as WAXD patterns for the blends presented in Figure 6 reveal four different diffraction patterns. The results suggest the presence of the α-phase only, while no peaks for beta (β) or any new crystals are detected. In the pole figures, the first and second rings represent the patterns of the 110 and 040 crystalline planes, respectively.2,21 It is known that the normal to the 110 plane is the bisector of the a and b axes, while the normal

shishes aligned in MD and kebabs perpendicular to MD, confirm that our system has a shish-kebab structure.2,43 Using FTIR, the measured Hermans crystalline orientation functions for various blends are plotted in Figure 5. It is obvious that the addition of modifiers lowers the orientation of the crystalline phases, due to their lower molecular weight.2 Lower crystalline orientation levels were expected for the blends of PPAA in comparison with the blends containing PPMA, due to the lower molecular weight of the former.2 Yet, the crystalline orientation of the blends remained approximately the same in the range of low modifier contents. However, PP/PPMA blends showed lower crystalline orientation as compared to the PP/ PPAA blends for large modifier contents, which could be related to the presence of ethylene groups in PPMA. As reported previously,12 the presence of ethylene groups was found in the PPMA polymer only. Thus, the presence of ethylene groups could reduce the compatibility and even the miscibility between the PP matrix and modifiers.48,49 Also, traces of maleic 10790

DOI: 10.1021/acs.iecr.5b02141 Ind. Eng. Chem. Res. 2015, 54, 10787−10795

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Figure 4. Weighted relaxation spectra for the neat PP0.8 and its blends with (a) PPMA and (b) PPAA; T = 190 °C.

Figure 5. Hermans crystalline orientation parameter (obtained from FTIR) for the neat PP0.8 precursor film and its blends with (a) PPMA and (b) PPAA.

to the 040 is along the b-axis of crystal unit cells.7 When the crystal orientation is random, circular rings are observed. Otherwise, in the case of the large crystal orientation, the arcs with strong intensities are observed.25 Broader and less intense arcs in the equatorial zone are observed as the amount of PPMA increases, implying lower orientation for the crystal lamellae, which is in agreement with the FTIR results.2,21 The pole figures of the 110 and 040 planes can also be used to quantitatively analyze the crystalline orientation, as shown in Figure 7. The pole figures of the neat PPMA precursor film show slight orientations of these two planes in MD and ND, respectively. However, by decreasing the PPMA content in PP0.8/PPMA, a significant orientation and alignment of the 110 and 040 planes (b-axis) crystallographic planes is observed along both TD and ND. This indicates less orientation with increasing PPMA content as observed above. Furthermore, the pole figures of the sample obtained from blending 20 wt %

PPMA with PP0.8 is approximately similar to the neat PP0.8 with slightly lower orientation intensities. The terms of cos2(ϕ) of the crystalline axes along MD, TD, and ND can be used to show the orientation features of the precursor films. The triangular graph of Figure 8 presents the results. Blending with PPMA causes movement of the c-axis of the crystals from MD (0.5−0.8) and the a-axis from the TD and ND. However, the orientation of the b-axis along MD is very small for all blends and does not take a position farther away from the TD and ND planes. So, the results show the significant decreases in the crystalline orientation of the films by blending with PPMA, in agreement with the FTIR results (see Figure 5). The results for the Hermans crystalline orientation parameter of the neat PP0.8 precursor film and its blends with PPMA are presented in Figure 9. It is clear that the results of WAXD are slightly larger than the orientation obtained using the FTIR technique, but the same trend is observed.5 The results of WAXD 10791

DOI: 10.1021/acs.iecr.5b02141 Ind. Eng. Chem. Res. 2015, 54, 10787−10795

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are mainly due to the information on the crystalline phase.2 So, the discrepancies in the orientation values are probably caused by the peak deconvolution and contribution of the amorphous part in FTIR. However, these results prove that FTIR is a reliable characterization technique to determine the crystalline orientation factors. The relaxation of polymer chains is the key parameter, which is controlled by the structure of the polymer and the cooling process. Quenching prevents relaxation of the chains from their stretched state, which produces shishs upon crystallization.37 However, the lamellae growth, secondary crystallization, might be affected by rapid cooling. Therefore, optimum cooling conditions and enough time should be employed to have the appropriate crystalline structure. As all of the blends were prepared under the same processing conditions, the results help to find the relationship between the material structure and morphology. The results suggest a correlation between the crystalline orientation (using FTIR values) and the characteristic relaxation time, τc, as presented in Figure 10 for the PP0.8/ PPMA, PP0.8/PPAA, PP2.8/PPMA, and PP2.8/PPAA blends. A good linear correlation is observed: the longer is the relaxation time, the higher is the crystalline orientation. The high molecular weight species (longer relaxation time),50 due to their underlying relaxation behavior, have the important effect on the formation of fibrillar structure.25 So, the results confirm

Figure 6. 2D WAXD patterns of the films: (a) neat PP0.8, (b) PP0.8+PPMA (20 wt %), (c) PP0.8+PPMA (50%), and (d) neat PPMA; and (e) diffraction spectra with integration through the circles of the blends.

Figure 7. Pole figures from WAXD for the neat PP0.8 precursor film and its blends with PPMA; schematics showing the assumed crystal orientation as well as the crystal block coordinates and the film production axes, respectively. 10792

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Figure 8. Orientation characteristics as cos2(ϕ) of the crystal axes (a, b, and c) along MD, TD, and ND for the neat PP0.8 precursor film and its blends with PPMA; the schematic on the left shows the crystal block coordinates and the film production axes.

Figure 9. Hermans crystalline orientation parameter using FTIR and WAXD techniques for the neat PP0.8 precursor film and its blends with PPMA.

Figure 10. Hermans crystalline orientation parameter versus the characteristic relaxation time, τc, for the PP0.8/PPMA, PP0.8/PPAA, PP2.8/PPMA, and PP2.8/PPAA blends.

the dependence of the crystalline orientation (crystalline structure) on the relaxation time of materials (molecular structure).

conditions, which were kept constant. A linear correlation was obtained between the characteristic relaxation time and the crystalline orientation function. Because of relaxation behavior, the high molecular weight species have the important effect on the formation of the fibrillar structure or polymer nuclei. It is also shown that rheology could be a very useful tool to optimize the processing conditions for obtaining appropriate precursor films for the preparation of porous membranes. Increasing the content of modifier results in shifting the characteristic relaxation time to shorter times. So, for preparing precursor films with large enough crystalline orientation from blends or lower Mw polymers (lower relaxation time), applying a larger draw ratio or changing the position of air cooling after die slit is required. In addition, a lower crystalline orientation was observed for the PPMA blends as compared to the PPAA blends, which

4. CONCLUDING REMARKS In this study, the rheological properties and crystalline structures of blends of polypropylene (PP) with maleic anhydride and acrylic acid grafted PP (PP-g-MA and PP-gAA) were investigated. For producing the precursor films and, consequently, microporous membranes with the appropriate crystalline lamellar morphology, the polypropylene chains should preserve their elongated form to serve as initial nuclei for the later development of lamellar crystals. The relaxation of the chains is the strongest factor for this stage. This factor is controlled by the material structure and the processing 10793

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(16) Balzano, L.; Kukalyekar, N.; Rastogi, S.; Peters, G. W. M.; Chadwick, J. C. Crystallization and Dissolution of Flow-Induced Precursors. Phys. Rev. Lett. 2008, 100, 048302. (17) Lamberti, G. Flow-induced crystallization during isotactic polypropylene film casting. Polym. Eng. Sci. 2011, 51, 851. (18) Baert, J.; Puyvelde, P. V.; Langouche, F. Flow-Induced Crystallization of PB-1: From the Low Shear Rate Region up to Processing Rates. Macromolecules 2006, 39, 9215. (19) Bischoff White, E.; Henning Winter, H.; Rothstein, J. Extensional-flow-induced crystallization of isotactic polypropylene. Rheol. Acta 2012, 51, 303. (20) Coppola, S.; Grizzuti, N. Micro-Rheological Modeling of FlowInduced Crystallization in Mixed Shear/Extensional Flows. Macromol. Symp. 2001, 169, 137. (21) Tabatabaei, S. H.; Carreau, P. J.; Ajji, A. Effect of processing on the crystalline orientation, morphology, and mechanical properties of polypropylene cast films and microporous membrane formation. Polymer 2009, 50, 4228. (22) Nogales, A.; Hsiao, B. S.; Somani, R. H.; Srinivas, S.; Tsou, A. H.; F.J, B.-C.; Ezquerra, T. A. Shear-induced crystallization of isotactic polypropylene with different molecular weight distributions: in situ small- and wide-angle X-ray scattering studies. Polymer 2001, 42, 5247. (23) Somani, R. H.; Yang, L.; Zhu, L.; Hsiao, B. S. Flow-induced shish-kebab precursor structures in entangled polymer melts. Polymer 2005, 46, 8587. (24) Somani, R. H.; Hsiao, B. S.; Nogales, A.; Srinivas, S.; Tsou, A. H.; Sics, I.; Balta-Calleja, F. J.; Ezquerra, T. A. Structure Development during Shear Flow-Induced Crystallization of i-PP: In-Situ Small-Angle X-ray Scattering Study. Macromolecules 2000, 33, 9385. (25) Agarwal, P. K.; Somani, R. H.; Weng, W.; Mehta, A.; Yang, L.; Ran, S.; Liu, L.; Hsiao, B. S. Shear-Induced Crystallization in Novel Long Chain Branched Polypropylenes by in Situ Rheo-SAXS and − WAXD. Macromolecules 2003, 36, 5226. (26) Pogodina, N. V.; Lavrenko, V. P.; Srinivas, S.; Winter, H. H. Rheology and structure of isotactic polypropylene near the gel point: quiescent and shear-induced crystallization. Polymer 2001, 42, 9031. (27) Vleeshouwers, S.; Meijer, H. H. A rheological study of shear induced crystallization. Rheol. Acta 1996, 35, 391. (28) Liedauer, S.; Eder, G.; Janeschitz-Kriegl, H.; Jerschow, P.; Geymayer, W.; Ingolic, E. On the Kinetics of Shear Induced Crystallization in Polypropylene. Int. Polym. Process. 1993, 8, 236. (29) Eder, G.; Janeschitz-Kriegl, H.; Krobath, G. Shear induced crystallization, a relaxation phenomenon in polymer melts. Prog. Colloid Polym. Sci. 1989, 80, 1. (30) Jerschow, P.; Janeschitz-Kriegl, H. The Role of Long Molecules and Nucleating Agents in Shear Induced Crystallization of Isotactic Polypropylenes. Int. Polym. Process. 1997, 12, 72. (31) Chellamuthu, M.; Arora, D.; Winter, H. H.; Rothstein, J. P. Extensional flow-induced crystallization of isotactic poly-1-butene using a filament stretching rheometer. J. Rheol. 2011, 55, 901. (32) Hadinata, C.; Boos, D.; Gabriel, C.; Wassner, E.; Rüllmann, M.; Kao, N.; Laun, M. Elongation-induced crystallization of a high molecular weight isotactic polybutene-1 melt compared to shearinduced crystallization. J. Rheol. 2007, 51, 195. (33) Sentmanat, M.; Delgadillo-Velázquez, O.; Hatzikiriakos, S. Crystallization of an ethylene-based butene plastomer: the effect of uniaxial extension. Rheol. Acta 2010, 49, 931. (34) Iervolino, R. Rheology and morphology of the flow induced crystallization in polymers. Ph.D. Thesis; Università degli Studi di Salerno: Salerno, 2006. (35) Derakhshandeh, M.; Hatzikiriakos, S. Flow-induced crystallization of high-density polyethylene: the effects of shear and uniaxial extension. Rheol. Acta 2012, 51, 315. (36) Termonia, Y. Coil-stretch transition in deformation flows. J. Polym. Sci., Part B: Polym. Phys. 2000, 38, 2422. (37) Yang, L.; Somani, R. H.; Sics, I.; Hsiao, B. S.; Kolb, R.; Fruitwala, H.; Ong, C. Shear-Induced Crystallization Precursor Studies in Model Polyethylene Blends by in-Situ Rheo-SAXS and Rheo-WAXD. Macromolecules 2004, 37, 4845.

was attributed to the presence of ethylene groups in PPMA. Also, good agreement with the crystalline orientation results of transmission FTIR and wide-angle X-ray analysis was observed.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We gratefully acknowledge the NSERC Network for Innovative Plastic Materials and Manufacturing Processes (NIPMMP) for the funding of this work.

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REFERENCES

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DOI: 10.1021/acs.iecr.5b02141 Ind. Eng. Chem. Res. 2015, 54, 10787−10795