Extension-Induced Crystallization of Poly(ethylene oxide) Bidisperse

The extensional rheometer used in this study can apply well-defined extension .... With a constant strain rate (ε̇) of 3.1 s–1, the strain was var...
0 downloads 0 Views 2MB Size
Download PDF
Article pubs.acs.org/Macromolecules

Extension-Induced Crystallization of Poly(ethylene oxide) Bidisperse Blends: An Entanglement Network Perspective Kunpeng Cui, Lingpu Meng, Youxin Ji, Jing Li, Shanshan Zhu, Xiangyang Li, Nan Tian, Dong Liu, and Liangbin Li* National Synchrotron Radiation Lab and College of Nuclear Science and Technology, CAS Key Laboratory of Soft Matter Chemistry, University of Science and Technology of China, Hefei, China S Supporting Information *

ABSTRACT: The role of long chains in extension flow-induced crystallization was studied with a combination of extension rheological and in situ small-angle X-ray scattering (SAXS) measurements at 52 °C. To elucidate the effects of long chains, bidisperse blends of poly(ethylene oxide) (PEO) with the long-chain concentration above the overlap concentration were prepared, constructing long-chain entanglement network in short-chain matrix. Rheological data of step extension on PEO melt are divided into two regions with fracture strain of pure shortchain sample as a boundary. Distinctly different features of crystallization kinetics and crystal morphologies are observed in these two regions, exactly corresponding to rheological behavior. A new mechanism based on entanglement network perspective is proposed, in which the second entanglement network constructed by long chains has three effects: (i) helping flow to change the free energy of polymer melt more effectively; (ii) ensuring the specific work can impose on the system; (iii) favoring the formation of precursors. This mechanism captures both rheological observation and crystallization behavior successfully and offers a new viewpoint for FIC study.

■

INTRODUCTION Flow-induced crystallization (FIC) is not only an interesting scientific topic but also an important industrial issue. The degree of crystallinity and crystal orientation of semicrystalline polymer materials strongly affects their physical properties, which is associated with FIC during polymer processing.1−4 The macroscopic flow forces polymer chains to be oriented or stretched and further controls the structural hierarchy from subnanometer to micrometer length scales.5 It is generally accepted that flow profoundly alters the rate of crystallization6−8 and induces the formation of the so-called shish-kebab structure, which significantly increases stiffness and strength.9−13 Nevertheless, despite decades of research, no satisfactory molecular theory of FIC has been achieved yet. The roles of molecular and flow field parameters and the structures of nuclei or precursors have not been illuminated well. Great attention has been paid to the influence of molecular parameters on FIC in recent years, focusing on the roles of short and long chains.14−18 According to coil−stretch transition (CST), for specified flow conditions only chains longer than the critical weights (M*) can be stretched and the rest shorter ones remain in the coil state.19,20 The stretched chains transform into shish and the coiled chains crystallize as kebab. Although CST enjoys great success in analyzing the molecular mechanism for shish-kebab,20−25 it is still a controversial viewpoint. Pennings et al. proposed that the stretched network rather than the stretched individual chains is © 2014 American Chemical Society

responsible for shish formation, since full extension of individual chains is unlikely to happen in the entangled polymer melt.26 Commonly, only portions of the chain segments between topological constraints are thought to be extended at typical experimental conditions.27,28 Although the viewpoint of stretched network has been demonstrated by rheo-SAXS experiment in polyethylene (PE) recently,29 it is still required confirmation in other polymer systems. For isotatic polypropylene (iPP), we found that the formation of precursors during flow is responsible for shish formation and proposed a “ghost nucleation” model.30 Here, the appearance of heterogeneities in melt structure further complicates the interpretation on the effect of long chains in FIC. To explore the role of long chains, bidisperse polymer blends, comprised of a tiny fraction of long chains in shortchain matrix, were widely used in FIC study.25,27,28,31−47 A brief summary of the effects of long chains on FIC is given herein. (i) Changing the morphologies of samples. It was suggested that the long chains play an important role in formation of shish-kebab morphology. Concerning flow-induced nuclei (shish) for shish-kebab morphology, Hsiao et al. emphasized that long chains dominate the formation of shish, where CST may occur in section of long chains.27,39,42 However, Kornfield Received: October 2, 2013 Revised: January 2, 2014 Published: January 15, 2014 677

dx.doi.org/10.1021/ma402031m | Macromolecules 2014, 47, 677−686

Macromolecules

Article

Bidisperse blends with different long-chain concentrations were prepared. Solution blending procedures were used to ensure that two components were intimately mixed at the molecular level. The blending procedure was listed as following. The mixture of LMW and HMW components was first dissolved in ethanol to form a homogeneous solution and then held at 70 °C for 60 min with continuous stirring under a nitrogen atmosphere. Precipitates were obtained by keeping the solution at 0 °C for 12 h. The precipitates were dried in a vacuum at 50 °C for 3 days to remove the residual ethanol. Control sample of pure LMW PEO was also prepared as reference material using the same procedure as for the blends. The precipitates were molded to plates with thickness of 1 mm by a vulcanizing press at 80 °C and then were cut into rectangular shape with dimensions of 30 × 18 × 1 mm3 for study. The concentrations of HMW component in the blends (0.5, 1.5, 5, and 10 wt %) designed were all higher than the estimated c* (0.004 g/ cm3), which aim to construct a HMW PEO entanglement network in LMW PEO matrix. The overlap concentration (c*) was estimated based on the equation c*= 3Mw/4π[⟨Rg2⟩1/2]3Na, where Rg is the radius of gyration and Na is Avogadro’s number. The value of ⟨Rg2⟩1/2/ Mw1/2 is 0.39 Å/(g/mol)1/2 based on small-angle neutron scattering measurements.48 Rheological Measurements. The linear viscoelastic properties of the LWM PEO melt were determined by small amplitude oscillatory shear (SAOS) with 25 mm plate−plate geometry (TA-AR2000EX, TA Instruments). The chosen strain of 2% fell well within the linear viscoelastic regime. The experimental temperature was varied from 80 to 120 °C. All the rheological experiments were performed in a nitrogen atmosphere to avoid possible sample degradation. The terminal relaxation time (τd) of LMW PEO for molecular disengagement is about 3 s at 80 °C (Figure 1). Through extrapolating linear

et al. argued that shish is consisted of both long and short chains, in which the long chains play a catalytic role and recruit other chains into formation of shish.28,44,45 The results of Kanaya et al. implied that the shish nuclei are derived from the deformed network of long chains.33 (ii) Enhancing the crystallization kinetics. It is believed that the long chains undergo a significant degree of orientation or stretch under flow. Kornfield et al. reported that long chains moderately enhance the formation of pointlike precursors but greatly enhance the formation of threadlike precursors.28 (iii) Altering the external field for crystallization. Peters et al. investigated the nonisothermal crystallization of PE under shear flow and observed that the incorporation of long chains shifts the onset temperature for crystallization.38 Mykhaylyk et al. suggested that the critical work required for flow-induced nucleation is affected by long-chain concentration.36 It was reported recently that the critical stress to induce threadlike nuclei decreases with the increase of long-chain concentration.45 The concept of overlap concentration (c*) is introduced to manifest the role of entanglement network of long chains.28,33,39 As polymer melt is a transient network constructed with highly entangled chains, an orientated network of entangled chains should be produced under the flow field. At low long-chain concentration, heterogeneities of long chains are easily to occur under flow field due to the low density of long-chain entanglements. But when long-chain concentration is high enough, the scenario of FIC may be different. Because of the high long-chain entanglement densities, a homogeneous deformed long-chain network similar to network material containing cross-links will be presented under flow. The topology of flow-induced morphology thus should be determined by the interplay between the flow field and the dynamics of entangled network. Therefore, systematical study on the influence of long-chain concentration on crystallization may gain deeper insight into the physics of the effect of long chains. The present study focuses on influences of the long chain and its concentration on crystallization behavior under shortterm extension flow. For this purpose, a combination of homemade extensional rheometer and in situ SAXS measurement was used to investigate FIC of PEO. A suitable PEO bidisperse blends system was chosen with long chains far longer than short ones, to unambiguously identify the effects due to stretch of long chains. The extensional rheometer used in this study can apply well-defined extension flow field and provide stress information, which is necessary to derive flow-induced conformational changes. Two regions exist in the rheological behavior, which exhibit distinctly different features of crystallization kinetics and crystal morphologies. The crystallization behavior coincides exactly well with the rheological observation. On the basis of the perspective of polymer entanglement network, we propose a new mechanism which can capture both the rheological and crystallization behavior successfully. The effects of the long chain and its concentration on crystallization in different stages of strain are also discussed.

■

Figure 1. Storage and loss modulus (G′ and G″) curves vs the oscillation frequency for LMW PEO melts at 80 °C. viscoelasticity data with time−temperature superposition,49 the τd is 8 s at 52 °C for LMW PEO. For HMW PEO, the estimated relaxation times for orientation (τd) and for stretch (τS) are about 1326 and 3 s, respectively, at 52 °C according to the relations τs ∼ Z2 and τd ∼ Z3.31,50 Z is the number of entanglements per chain. Note that the above calculation is a rough estimation and average result because of the polydispersity of the sample. Instrumentation. A homemade extensional rheometer was used in this study, which can apply well-defined thermal history and impose extension flow field. The details of this apparatus have been described elsewhere.51 In brief, the ends of sample were secured to two geared drums, and an extension flow field is obtained by reversing rotations of two drums. With this apparatus, the strain and strain rate can be controlled independently. In situ SAXS measurements were carried out at BL16B of the Shanghai Synchrotron Radiation Facility (SSRF). The X-ray wavelength was 0.124 nm. A Mar165 CCD detector (2048 × 2048 pixels with pixel size of 80 μm) was employed to collect time-resolved twodimensional (2D) patterns, which placed 5320 mm from the sample. The 2D scattering images were analyzed with Fit2D software from the European Synchrotron Radiation Facility.52 Experimental Procedures. Each sample was initially held at 80 °C for 10 min to erase the thermal history. Then the sample was cooled to the crystallization temperature (Tc). Immediately after Tc was reached, a step strain with constant strain rate was imposed to the

EXPERIMENTAL SECTION

Materials and Sample Preparation. High molecular weight (HMW) PEO with weight-average molecular weight (Mw) of 1017.2 kDa and low molecular weight (LMW) with Mw of 185.2 kDa, which have polydispersity of 1.3 and 9.0, respectively, were used in this study. The molecular weights of PEO were determined by GPC-MALLS. Both samples were provided by Liansheng Chemical Co. of Shanghai. 678

dx.doi.org/10.1021/ma402031m | Macromolecules 2014, 47, 677−686

Macromolecules

Article

sample. Temperature stability was maintained within ±0.5 °C, and a nitrogen gas flow was used to homogenize the system temperature and prevent sample from degradation. Tc was 52 °C in this work. With a constant strain rate (ε̇) of 3.1 s−1, the strain was varied from 0 to 3.5. HMW 0, HMW 0.5, HMW 1.5, HMW 5, and HMW 10 are used to represent samples with long-chain concentration of 0, 0.5, 1.5, 5, and 10 wt %, respectively. Date Analysis. SAXS data were corrected for background scattering, detector spatial distortion, and X-ray beam fluctuation first. Then the 2D SAXS patterns were sector integrated to obtain the scattered intensity I(q) as a function of q = (4π/λ) sin θ (Figure 4b), where q is the module of scattering vector, λ is the wavelength of Xray, and 2θ is the scattering angle. The integrated intensity (Iint) was obtained by summing up the I(q) over the whole accessible q range: Iint = ∫ qminqmaxI(q) dq. The normalized intensity is defined as [Iint(t) − Iint(min)]/[Iint(max) − Iint(min)], where Iint(t) is the integrated intensity at time t, and Iint(min) and Iint(max) are the minimum and maximum integrated intensity, respectively. The orientation of lamellar crystals is characterized by Herman’s orientation parameter ( f) which is defined as f = [3⟨cos2 ϕ⟩ − 1]/2, where ϕ is the angle between the reference direction (extensional direction) and the normal direction of the lamellae. Thus, f has a value of 1 when lamellar normal parallels to the flow direction, and lamellae with no preferred orientation give an f of 0. The full width at half-maximum (fwhm), extracted from azimuthal intensity of the SAXS at the peak value, was used to calculate f. The azimuthal distributions are fitted with Lorentz functions well in our study (Figure S1, Supporting Information).

ensure that no structure generates during extension. Weak strain hardening is observed for both HMW 0.5 and 10 at large strains, suggesting indeed long-chain stretch occurs during extension. This is also supported by a large Deborah number for stretch (DeS = ε̇ × τS > 1), where ε̇ is 3.1 s−1 and τS of HMW PEO is 3 s. The increment of stress during strain hardening at 52 °C is far steeper than that at 80 °C. By rough estimation, the slopes of strain hardening at 52 and 80 °C are 0.96 and 0.06, respectively, for HMW 10. Photographs of PEO melts after being subjected to different strains at 52 °C are shown in Figure 3. For pure short chains

Figure 3. Photographs of PEO samples after subjected to different strains at 52 °C.

■

(HMW 0), samples exhibit homogeneous deformation with strain smaller than 0.9 but delayed cohesive fracture with strain larger than 0.9 after the cessation of step extension. However, once long chains were added, all samples keep integrality without fracture in the explored range of strain. Thus, two flow regions, termed as region I and region II, are divided by a boundary of strain 0.9 based on the rheological behavior. Here, region I represents weak or moderate strains and region II represents strong strains. Interestingly, the corresponding crystallization behaviors also differ in region I and region II. The crystallization behaviors of HMW 0.5, 1.5, and 5 are almost identical in those two regions. For conciseness, we only give the results of HMW 0, 0.5, and 10 in the following. SAXS data at strain 0.9 are selected to represent the crystallization behavior in region I. A series of 2D SAXS patterns at different crystallization time are presented in Figure 4a with the flow along horizontal (meridional) direction. The time t = 0 s corresponds to the first scattering pattern obtained after cessation of extension. At a short delay of 80 s after cessation of extension, both HMW 0.5 and 10 samples clearly exhibit streak scattering features along flow direction. The meridional streak suggests the presence of poor periodicity of lamellae at the early stage of crystallization. Increasing the crystallization time, the meridional streak gradually transforms into fan-shaped scattering, indicating the formation of welldefined lamellar stacks with lamellar normal oriented along the flow direction. While without adding of long chains (HMW 0), SAXS pattern shows a very weak diffuse scattering of melt at t = 80 s, indicating the absence of any detectable structure. (Note that the ellipse scattering close to the beamstop may come from the residual catalyst of PEO.) The weak scattering maximum appears at crystallization time of about 480 s, which concentrates slightly in the meridional direction, suggesting the formation of weakly oriented lamellar stacks. On the basis of 2D SAXS patterns, the evolution of normalized intensity of lamellar signal during the crystallization process is further analyzed in Figure 4b. The half-time of crystallization (t1/2) can be quantitatively extracted, which is defined as the time for normalized intensity reaches 0.5. A decreasing process is

RESULTS The extensional rheological behavior of PEO melts is exhibited first. The engineering stress (σengr)−Hencky strain curves during step extension at 52 °C are presented in Figure 2a. For

Figure 2. Engineering stress−Hencky strain curves for PEO blends with three long-chain concentrations (a) at 52 °C and (b) at 80 °C.

pure short chains (HMW 0), the stress increases almost linearly with strain initially. Following the linear deformation region, a stress overshoot is observed, which leads to a stress maximum. After exceeding the stress maximum, the evolution of stress with strain undergoes a decline process. This nonlinear rheological behavior is typical for polymer melts subjected to step extension flow with a constant strain rate.53 Upon adding long chains, the trend of stress evolution is rather different. As blends show similar stress evolution with strain at various longchain concentrations, only HMW 0.5 and HMW 10 are demonstrated for conciseness. After the stress overshoot, strain hardening is observed rather than decline process of stress. The critical strain for the onset of strain hardening maintains at 2.3, while the stress is significantly enhanced with the increase of long-chain concentration. Such a phenomenon of strain hardening indicates that viscosity changes during flow due to structure formation or/and stretch of long chains. In order to check whether stretch of long chains occurs during extension, the same extension experiments are conducted at 80 °C (Figure 2b) above the equilibrium melting point of PEO (69 °C) to 679

dx.doi.org/10.1021/ma402031m | Macromolecules 2014, 47, 677−686

Macromolecules

Article

extension (t = 0 s), both scattering patterns with strains of 1.5 and 3.5 show diffuse scattering, which can be attributed to the absence of any detectable structure. While for a moderate strain of 2.5, SAXS pattern clearly shows the onset of crystallization with lamellar scattering right after cessation of extension (t = 0 s), as demonstrated by the obvious meridional streaks. The other is that, with increasing the crystallization time, two scattering streaks appear perpendicular to the extension direction at the later stage of crystallization with strains of 2.5 and 3.5. This equatorial streak is different from the common equatorial streak resulted from the flow-induced shish structure, which has been reported in plenty of SAXS studies of FIC at the early stage of crystallization.11,30,54 The unique feature of the equatorial streak suggests that, in addition to density fluctuation in lamellar structure, some kind of extra large electron density contrast is detected in the later stage of crystallization. The normalized intensity scattered in the meridional and equatorial regions (see Figure 4b) for HMW 0.5 is illustrated in Figure 6. With strain of 1.5 (Figure 6a), the equatorial

Figure 4. (a) 2D SAXS patterns of the PEO samples at selected time intervals after extension at 52 °C with strain of 0.9. (b) 1D SAXS normalized intensity profiles of lamellar (meridional) scattering.

observed after the normalized intensity reaches the maximum, which attributes to volume fraction of lamellae larger than 50%. This is because the normalized intensity is proportional to the product of volume fractions of lamellae and surroundings. Thus, the definition of t1/2 used here may overestimate the overall crystallization kinetics. At current experimental conditions, the t1/2 are 900 and 240 s for HMW 0 and HMW 0.5, respectively, which indicates the crystallization kinetics is largely promoted by the addition of long chains. From the t1/2 values for HMW 0.5 (240 s) and HMW 10 (320 s), it appears that the concentration of long chains has a weak influence on crystallization kinetics in region I. For region II, the crystallization behavior depends on the long-chain concentration. Figure 5 presents a series of 2D SAXS patterns of HMW 0.5 during crystallization after step extension at 52 °C. From Figure 5, we notice two interesting phenomena. One is that, immediately after cessation of

Figure 6. SAXS meridional and equatorial intensities of the HMW 0.5 samples at 52 °C as a function of time with strain of (a) 1.5 and (b) 3.5.

scattering intensity decreases with the crystallization time first, which also corresponds to the increase process of lamellar (meridional) scattering intensity. A leveling off is exhibited in both meridional and equatorial intensity after the crystallization time of about 500 s. The negative correlation between the evolution of meridional and equatorial scattering intensity is because they are resulted from different structures. When crystallization starts, there are overall increasing volumes of regions that have contrast between crystallites and amorphous due to the growth of lamellae, while the electron density contrast decreases between residual catalysts and surroundings with crystallization, because the contrast between residual catalysts and amorphous is larger than the contrast between residual catalyst and crystallites. Different from strain of 1.5, the equatorial scattering intensity increases rather than levels off after the crystallization time of about 400 s (as indicated by red line in Figure 6b) for strain of 3.5 due to formation of equatorial streaks. Figure 7 displays 2D SAXS patterns during the crystallization process for HMW 10 after extension at 52 °C. Similar to the results of HMW 0.5, both the scattering patterns with strains of 1.5 and 3.5 present diffuse scattering of melt immediately after cessation of extension (t = 0 s). For the moderate strain of 2.5, instead of meridional streaks, two discrete lobes emerge, suggesting the formation of well-defined lamellar stacks right after cessation of extension (t = 0 s). Different from results of HMW 0.5, no equatorial streak forms for HMW 10 at the later

Figure 5. 2D SAXS patterns of the HMW 0.5 samples at selected time intervals after extension at 52 °C with three strains. 680

dx.doi.org/10.1021/ma402031m | Macromolecules 2014, 47, 677−686

Macromolecules

Article

hardening in stress−strain plots. While for HMW 10, just two lobes scattering appear for the final crystallized samples. In order to explore the effect of strain on the initial stage of crystallization, 1D SAXS intensity profiles of meridional (lamellar) signal just after cessation of extension (t = 0 s) is given in Figure 9. For HMW 10 (Figure 9b), with strain smaller

Figure 7. 2D SAXS patterns of the HMW 10 samples at selected time intervals after extension at 52 °C with three strains. Figure 9. 1D SAXS meridional integrations I(q) immediately after cessation of extension (t = 0 s) with different strains at 52 °C for (a) HMW 0.5 and (b) HMW 10 samples.

stage of crystallization, suggesting unexpected absence of discernible extra large electron density contrast. Figure 8 summarizes 2D SAXS patterns for crystallized samples with different strains and long-chain concentrations. Five types of scattering patterns were registered for PEO blends after extension-induced crystallization: diffraction ring, scattering arcs, fan-shaped diffraction, lobes with and without equatorial streaks. The diffraction ring indicates the morphology with no preferred orientation of lamellae. The second and the third types of SAXS patterns are deemed as oriented structures, which are commonly observed in FIC study at moderate flow conditions.35,46 The last two types of scattering patterns, two lobes with and without equatorial streaks, will be discussed in detail in the following. For HMW 0, with strain smaller than 0.9, the isotropic ring in scattering patterns indicates the formation of lamellar structure with random orientation. With incorporation of long chains, there exists a transition of scattering patterns from arcs to fan-shaped scattering with the increase of strain. For HMW 0.5, 1.5 and 5, the critical strain of the transition is 0.9, which decrease to 0.6 for HMW 10. With strain from 2.5 to 3.5, two lobes with equatorial streaks scattering generate for HMW 0.5, 1.5, and 5 at the later stage of crystallization, which corresponds to strain

than 1.5, the 1D intensity profiles show the typical feature of an amorphous polymer melt. Increasing strain from 2.0 to 2.5, obvious scattering maximum around q of 0.15 nm−1 is observed, indicating the formation of lamellar stacks. By further increasing strain from 3.0 to 3.5, no obvious peak appears. This indicates that the time for the onset of crystallization increases with increasing strain beyond a critical value (e.g., 2.5 for HMW 10). This unusual crystallization behavior is also observed in HMW 0.5 (Figure 9a). The orientation parameter of lamellar crystals is employed to illustrate the structural evolution in the early stage of crystallization. In Figure 10b for HMW 10, f shows a monotonic decrease process with a strain of 1.5. By increasing strain to 2.5, f follows a nonmonotonic evolution, first increasing and then decreasing with crystallization time. By further increasing strain to 3.5, contrary to the results with strain of 1.5, f exhibits a rapid increase first and then approaches a steady-state level. The HMW 0.5 also show similar results (Figure 10a). With constant long-chain concentration, taking HMW 10 as an example, the orientation parameters are 0.96,

Figure 8. 2D SAXS patterns for crystallized samples with different strains and long-chain concentrations at 52 °C; the number in the parentheses at top left corner of each pattern represents the value of strain. 681

dx.doi.org/10.1021/ma402031m | Macromolecules 2014, 47, 677−686

Macromolecules

Article

Figure 10. Evolution of lamellar crystals orientation parameter of three strains in the early stage of crystallization at 52 °C for (a) HMW 0.5 and (b) HMW 10 samples.

Figure 12. Plots of (a) half-time of crystallization and (b) orientation parameter of final lamellar structure with strain for different long-chain concentrations at 52 °C.

0.95, and 0.84 respectively for strains of 1.5, 2.5, and 3.5 (t = 16 s), which decreases with the increase of strain. For the same strain of 3.5 and at same crystallization time of 16 s, the orientation parameter of HMW 0.5 is 0.95, which is much larger than HMW 10 ( f = 0.84). Based on Figure 10, it is obvious that the degree of lamellar order decreases with the increase of either strain or long-chain concentration in the early stage of crystallization under strong strains. The final long period (Lp= 2π/q) as a function of strain for crystallized samples with different long-chain concentrations is presented in Figure 11. Different from the viewpoint that the

with the increase of either strain or long-chain concentration. Both the half-time of crystallization and the orientation parameter exhibit a two-stage process. The transition of these two stages occurs at strain of 0.9, coinciding with the boundary of region I and region II defined by rheological behavior. With strain larger than 0.9, the decay of half-time of crystallization and the increase of the orientation parameter with strain are faint.

■

DISCUSSION Combining the results of extension rheological and in situ SAXS measurements, some interesting findings can be extracted. (i) The rheological behavior of step extension can be defined as two regions with a strain of 0.9 as boundary, namely region I and region II. For pure short-chain matrix, the sample may result into a delayed fracture with strain larger than 0.9. Once long chains were added, no fracture of the sample is observed in the whole explored range of strain. Moreover, strain hardening appears in region II with addition of long chains. (ii) A twostage process is presented in both the half-time of crystallization and orientation parameter of final lamellar crystals with the transition strain corresponding to the boundary of region I and region II. (iii) When long-chain concentration varies from 0.5 to 5 wt %, the critical strain to induce the fan-shaped scattering is 0.9. Equatorial streak arises with strain larger than 2.5, which corresponds exactly to the strain hardening in rheological behavior. However, by increasing long-chain concentration to 10 wt %, the critical strain to induce fan-shaped scattering reduces to 0.6, and no equatorial streak appears during the whole crystallization process. (iv) The final long period of sample depends on strain and long-chain concentration. Furthermore, the shapes of the final long period (Figure 11) and the stress (Figure 2) evolve with strain in a similar way. From the above results, it is obvious that an intrinsic correlation exists between rheological observation and crystallization behavior. As seen above, the incorporation of a small amount of long chains significantly alters the crystallization behavior of polymer. In region I, the crystallization kinetic is enhanced by addition of long chains. According to microrheological model which based on the Doi−Edwards theory and Lauritzen− Hoffman theory, the effect of the macroscopic flow on the free energy of polymer melt (ΔGf) is proportional to DeO2.56 Here, DeO = ε̇ × τd is the Deborah number for chain orientation. For LMW PEO, the DeO is about 25, while for HMW PEO, the DeO is about 4111. The huge difference in DeO implies that the long chains can enhance the crystallization kinetics more remarkable than short ones under flow. Additionally, the long chains possess a long relaxation time, being able to retain more

Figure 11. Final long period as a function of strain for crystallized samples with different long-chain concentrations at 52 °C.

final long period is a thermodynamic result determined by temperature, the long period shows strain and long-chain concentration-dependent behavior in this study. For HMW 0, the long periods are almost the same (around 39 nm) with increasing strain from 0.3 to 0.9. With addition of long chains, for all of blends studied here, the final long period first shows a rapid increase with strain smaller than 0.9 and then enters a faint increase region with strain from 0.9 to 2.0, followed by a second speed-up with strain up to 3.5. For example, the long periods of HMW 5 are 37.6 and 50.6 nm for strain of 0.3 and 3.5, respectively, which increase 13 nm by raising strain. Coincidently, the shape of stress−strain (Figure 2) is analogous to evolution of the long period with strain, indicating an intrinsic correlation between rheological and crystallization behavior in FIC. The half-time of crystallization and the orientation parameter of final lamellar structure are employed to study the influences of strain and long-chain concentration on FIC. Figure 12a shows the half-time of crystallization at different strains and long-chain concentrations. The overall trend of half-time of crystallization decreases with the increase of strain, which lies in the general expectation of reported experiments and some existing theory.55,56 At current experimental conditions, the concentration of long chains has a weak influence on the halftime of crystallization. Figure 12b shows the orientation parameter of final lamellar crystals at different strains and long-chain concentrations. The orientation parameter increases 682

dx.doi.org/10.1021/ma402031m | Macromolecules 2014, 47, 677−686

Macromolecules

Article

Figure 13. Schematic illustrations of nucleation and crystallization model for PEO blends. Red (thick) and blue (thin) curves represent HMW and LMW PEO chains. (A) Pure short chain matrix, (B) with low long-chain concentration, (C) with high long-chain concentration, and (D) structural morphologies formed after crystallization at different degree of deformation of polymer chains and corresponding SAXS patterns. Precursors or nuclei formed during extension are omitted here for concision.

further increasing the strain even entering into strain hardening. One possible explanation is that the long chains possess long relaxation time. The delay time for fracture is increased with the addition of long chains,53 which may result into the situation that crystallization occurs before fracture. Another possible explanation is that interchain structures have already been induced by flow during extension, which may be precursors or nuclei as generally assigned by the community. The occurrence of precursors in large strains is also supported by much larger slope of strain hardening at 52 °C than that at 80 °C. At 80 °C, sample exhibits delayed fracture after cessation of extension. At 52 °C, the precursors act as physical cross-links to ensure sample to sustain further deformation and even enter into strain hardening. This inference of structure formation during flow could be supported by the recent fast X-ray scattering results conducted by Peters et al.31 and Matsuba et al.58 The concentration of long chains seemingly has small influence to crystallization kinetics but large influence to crystal morphology. The crystal morphology is often related to chain deformation in FIC study. Nevertheless, the relationship between the intensity of flow field and the degree of flowinduced chain deformation is still not well established. A roundabout method to deduce chain conformation is based on final morphology of the sample, e.g., shish-kebab structure results from strong chain stretch. In our study, five types of scattering patterns are observed, which correspond to five morphologies because of the different degree of chain deformation during flow (Figure 13D). The diffraction ring (Figure 13D1), scattering arcs (Figure 13D2), and fan-shaped diffraction (Figure 13D3) are commonly observed in polymers after FIC. Here, we focus our attention on the morphologies formed at the strong strains. In previous studies, fan-shaped diffraction pattern is observed at moderate flow conditions, which may correspond to distorted spherulites proposed by Mykhaylyk et al.35 or sausages structures proposed by Kornfield et al.45 Further increasing the intensity of flow, shish-kebab morphology will be formed. While in our study, further increasing the intensity of flow, a pseudo shish-kebab scattering

deformation after extension. Thus, the degree of orientation is also increased with the addition of long chains. To disclose the role of long chains on FIC in region II, the effects of long chains on rheological behavior need to be clarified. For pure short-chain sample, delayed fracture occurs with strain larger than 0.9. A similar phenomenon has been observed and explained in our previous FIC study of iPP.30 In brief, PEO melts is considered as an entanglement network, which follows a viscoelastic deformation with DeO > 1. With strain larger than a critical value, which is 0.9 for PEO short chains in this study, the flow may lead to unbalance of forces and finally result into fracture of the sample (Figure 13A). An intriguing question is why no fracture happens with the addition of long chains in the explored flow conditions. The relaxation times of the two individual components in the PEO blends are widely separated. Based on the relations τS ∼ Z2 and τd ∼ Z3, the estimated τS of HMW and LMW chains are different by a factor of 30, whereas the difference in their τd reaches 166. The wide separation of relaxation times makes the HMW and LMW components to stay in different deformation regions under a given strain rate of 3.1 s−1. That is to say, the HMW component formed a second entanglement network, which is highly reasonable due to the large separation of its relaxation time from that of LMW chains. During flow, even though the LMW network has disintegrated, the HMW network did not give in because it involves much greater entanglement spacing.53,57 Thus, the blends can keep integrity with strain larger than 0.9 (Figure 13B,C). The observation that the onset of failure is significantly postponed by the addition of long chains is also reported in extension rheological experiment of styrene−butadiene random copolymers system.53 At this stage, the effect of long chains can be considered as the mediator that keeps the sample out from destructive dissipation due to the instability. This effect ensures that the specific work derived from flow can impose on the system and thus influence the crystallization behavior. The incorporation of long chains just postpones the onset of fracture strain. However, samples can keep integrity with 683

dx.doi.org/10.1021/ma402031m | Macromolecules 2014, 47, 677−686

Macromolecules

Article

question is why the long-chain entanglement network can be deformed homogeneously under flow field, as the straininduced concentration fluctuation may lead to local concentration enrichment of short and long chains. A possible explanation is the intermolecular force of PEO is tough due to hydrogen bond, which makes it difficult to generate concentration fluctuation. The precursors formed during flow also act as physical cross-links, preventing the homogeneous deformed network from the separation of short and long chains. A subsequent question is why the onset time of crystallization increases with the increase of strain for PEO blends under strong strains. Intuitively, strong strain is in favor of formation of precursors which assists the following crystallization process. However, the precursors already formed may be fragmented by further increasing the intensity of strain, which leads to an increase of onset time for crystallization. The cross-link network constructed by precursors is harder than surrounding chains. Thus, the tension is mainly transmitted through the cross-link joints under flow field. Even in polymer solids, crystal is easily destroyed under tensile deformation.59 Therefore, it is plausible that the precursors will be fragmented if the strain is strong enough. The evolution of the orientation parameter of initial lamellar crystals is another evidence to support the fragment process of precursors during flow. Two reasons may be responsible for that the degree of initial lamellar order decreases with the increase of strain or long-chain concentration under strong strains. One is that, in the fragment process during flow, the precursors may be forced to rotate or tilt, which decreases the initial degree of lamellar order. The other is that tilted segments of entanglement network are increased with increasing the long-chain concentration, which induce nucleation and also reduce the initial orientation of lamellae. Summarizing the above discussion, a mechanism for FIC of PEO blends is proposed in Figure 13. For low long-chain concentration entanglement network (from 1 to 11 times c*) at strong strains, despite the generation of precursors, the density of stretched chains is still low. Thus, fibrillar lamellar stacks develop on the nuclei that generated mainly from stretched chains. At the later stage of crystallization, the distance between neighboring fibrillar lamellar stacks becomes smaller with crystallization time. The large electron density contrast at the interspaces between neighboring fibrillar stacks leads to the emergence of equatorial streak when they get close to each other. This is supported by the formation of equatorial streak when the lamellar normalized intensity levels off (Figure 6b). The materials in the interspaces may be voids or amorphous. One possible situation is that voids form at the interspaces due to shrinkage of crystallization. The other possible situation is that the interspaces between lamellar structures (i.e., the growth front) are rich of entanglements and conformation defects especially at strong flow field, which make them to stay as amorphous. Additionally, dimensional parameters of the interspace are extracted from equatorial streak (see Figure S2 and Table S1, Supporting Information). For high long-chain concentration entanglement network, such as 22 times c* for PEO blends (HMW 10) in this study, with the assistance of precursors, high density and homogeneously spatial distribution of nuclei is presented at strong strains. As a consequence, small lamellar stacks with high orientation decorate the network homogeneously. Thus, no equatorial streak is observed even with strain entering strain hardening.

pattern, with two lobes and equatorial streaks that formed at the later stage of crystallization (Figure 13D4), appears with long-chain concentration varied from 0.5 to 5 wt %. The first needs to be clarified is that whether the equatorial streak formed at later stage of crystallization is attributed to the shish structure. This is because the shish scattering signal may behind in the beamstop in the early stage of crystallization and then the shish reach dimensions become able to scatter in the space probed by the detector. The normalized intensity of this equatorial streak increases with crystallization time when the normalized intensity of lamellae levels off (Figure 6b). If this equatorial streak is attributed to shish structure, it should be weakened and even disappear when the crystallization is completed as there is no longer sufficient electron density contrast. The shish-kebab structure is unlikely formed in our study because of the following reasons. No equatorial streak forms in the early stage of crystallization at all investigated strains and long-chain concentrations, implying the absence of shish structure. Meanwhile, the crystallization kinetics also gives the same implication. Peters et al. investigated shear-induced crystallization of iPP with a combination of the rheometer and optical microscope.8 They found that the half-time of crystallization first decreases with the increase of shear time and then reaches a plateau, where the crystal morphology stays as spherulite. Further increasing shear time, a second decrease in the half-time of crystallization is observed, which corresponds to shish-kebab structure. The second speed-up of crystallization kinetics is attributed to the formation of shish nuclei. This demonstrates that the morphology change of nuclei can be reflected by the crystallization kinetics. Thus, the crystallization kinetics also can be used as a criterion for shish formation. In our study, the half-time of crystallization does not show a plateau or a second decrease around the fan-shaped diffraction, which further confirms the absence of shish structure. One may wonder whether the absence of shish-kebab is caused by the insufficient stretching of the long chains. The occurrence of equatorial streak in scattering pattern corresponds to strain-hardening in stress−strain plots exactly, which is a hint of strong stretch of chains. Additionally, the formation of precursors enhances the relaxation time and makes chains to be stretched more easily. For long-chain concentration between 0.5 and 5 wt % (from 1 to 11 times c*), the sample can be considered as a dilute long-chain entanglement network (Figure 13B). Because of the density of long-chain entanglement is low, inhomogeneous deformation of long-chain network may happen as the stress transfer between adjacent molecules mainly through entanglements. Even though the long-chain entanglement network can be deformed homogeneously under flow field, taking HMM 0.5 as an example, one segment of long-chain entanglement network is surrounded by about 199 short-chain segments. Although the formation of precursors enhances the amount of stretched chains, the density of stretched chains is still low. Thus, fibrillar-like lamellar stacks form with nucleation mainly caused by stretched chains (Figure 13D4). For high long-chain concentration entanglement network (approaching 22 times c*, as indicated in Figure 13C), with the assistance of precursors, homogeneous deformation, and high density of stretched chains are expected to take place. All stretched chains enjoy equal opportunity to form nuclei. Thereby, the high nucleation density due to high density of stretched chains and homogeneously spatial distribution of nuclei are expected (Figure 13D5). One raised 684

dx.doi.org/10.1021/ma402031m | Macromolecules 2014, 47, 677−686

Macromolecules

Article

The evolution of the long period of final lamellar stacks with strain further supports the network view that we proposed. In region I, the long period increases with the increase of strain. This attributes to the orientation and stretch of polymer chains, which results in thicker interlamellar amorphous layer with high entanglement densities as constraint to crystallization. A similar result has also been observed in our previous FIC study of cross-linked network and free chains.54 By increasing the strain to region II, disentanglement of short chains occurs, which partially balances out strain effect and leads to a faint increase of final long period. However, precursors generate by further increasing the strain. Those precursors act as physical crosslinks, which freeze some entanglements between them. In other words, the formation of precursors enhances the effect of strain, exhibiting as a second speed-up in final long period.

Synchrotron Radiation Facility (SSRF). We thank Prof. Xulin Jiang (Wuhan) for assistance on GPC-MALLS measurements, Prof. Zhigang Wang (Hefei) for assistance on SAOS measurements, and Dr. Ningdong Huang for assistance with English corrections.

■

■

CONCLUSION In situ SAXS was carried out to investigate extension-induced crystallization of PEO blends, where the long-chain concentrations are all higher than the overlap concentration to construct long-chain entanglement network in short-chain matrix. To our knowledge, this is the first set of extensioninduced crystallization experiments in which the strain rate reaches the threshold for chain stretch for long chains of welldefined length and concentration. Two regions are identified by rheological behavior, which exhibit different crystallization kinetics and final crystal morphologies. On the basis of the perspective of polymer entanglement network, we offer a mechanism that captures both the rheological and crystallization behavior successfully. In this mechanism, the effect of the long chain is analyzed in different stage of strains. In region I, the enhancement of crystallization kinetics is ascribed to that the long chains can help flow to change the free energy of polymer melt more effectively. In region II, the long chains first act as mediator to keep the sample out from destructive dissipation due to fracture and ensure the specific work can impose on the system. By further increasing the intensity of flow, the long chains favor the formation of precursors and influence the density of stretched chains. This mechanism gives a new viewpoint for FIC study and may provide underpinnings for theoretical models.

■

ASSOCIATED CONTENT

S Supporting Information *

Methods for fitting azimuthal intensity and SAXS equatorial streak and the fitting results. This material is available free of charge via the Internet at http://pubs.acs.org.

■

REFERENCES

(1) Ward, I. M. Makromol. Chem., Macromol. Symp. 1988, 22, 59−82. (2) Perepelk, Ke. Angew. Makromol. Chem. 1972, 22 (Mar 30), 181. (3) Janeschitzkriegl, H.; Eder, G.; Krobath, G.; Liedauer, S. J. NonNewtonian Fluid Mech. 1987, 23, 107−122. (4) Yang, H. R.; Lei, J.; Li, L. B.; Fu, Q.; Li, Z. M. Macromolecules 2012, 45 (16), 6600−6610. (5) Tian, N.; Zhou, W.; Cui, K.; Liu, Y.; Fang, Y.; Wang, X.; Liu, L.; Li, L. Macromolecules 2011, 44 (19), 7704−7712. (6) Bashir, Z.; Hill, M. J.; Keller, A. J. Mater. Sci. Lett. 1986, 5 (9), 876−878. (7) Li, L. B.; de Jeu, W. H. Interphases Mesophases Polym. Cryst. II 2005, 181, 75−120. (8) Housmans, J. W.; Steenbakkers, R. J. A.; Roozemond, P. C.; Peters, G. W. M.; Meijer, H. E. H. Macromolecules 2009, 42 (15), 5728−5740. (9) Zhang, C. G.; Hu, H. Q.; Wang, D. J.; Yan, S.; Han, C. C. Polymer 2005, 46 (19), 8157−8161. (10) Cavallo, D.; Azzurri, F.; Balzano, L.; Funari, S. R. S.; Alfonso, G. C. Macromolecules 2010, 43 (22), 9394−9400. (11) Somani, R. H.; Yang, L.; Zhu, L.; Hsiao, B. S. Polymer 2005, 46 (20), 8587−8623. (12) Murase, H.; Ohta, Y.; Hashimoto, T. Macromolecules 2011, 44 (18), 7335−7350. (13) Chen, Y. H.; Zhong, G. J.; Lei, J.; Li, Z. M.; Hsiao, B. S. Macromolecules 2011, 44 (20), 8080−8092. (14) Azzurri, F.; Alfonso, G. C. Macromolecules 2005, 38 (5), 1723− 1728. (15) Vleeshouwers, S.; Meijer, H. E. H. Rheol. Acta 1996, 35 (5), 391−399. (16) Jerschow, P.; JaneschitzKriegl, H. Int. Polym. Proc. 1997, 12 (1), 72−77. (17) Graham, R. S.; Olmsted, P. D. Phys. Rev. Lett. 2009, 103 (11), 115702. (18) van Meerveld, J.; Peters, G. W. M.; Hutter, M. Rheol. Acta 2004, 44 (2), 119−134. (19) Degennes, P. G. J. Chem. Phys. 1974, 60 (12), 5030−5042. (20) Keller, A.; Kolnaar, J. W. H. Prog. Colloid Polym. Sci. 1993, 92, 81−102. (21) Steenbakkers, R. J. A.; Peters, G. W. M. J. Rheol. 2011, 55 (2), 401. (22) Dukovski, I.; Muthukumar, M. J. Chem. Phys. 2003, 118 (14), 6648. (23) Wang, M. X.; Hu, W. B.; Ma, Y.; Ma, Y. Q. Macromolecules 2005, 38 (7), 2806−2812. (24) Zhu, P. W.; Edward, G. Macromolecules 2004, 37 (7), 2658− 2660. (25) Kanaya, T.; Matsuba, G.; Ogino, Y.; Nishida, K. Abstr. Pap. Am. Chem. Soc. 2007, 234. (26) Smook, J.; Pennings, A. J. J. Mater. Sci. 1984, 19 (1), 31−43. (27) Keum, J. K.; Zuo, F.; Hsiao, B. S. Macromolecules 2008, 41 (13), 4766−4776. (28) Seki, M.; Thurman, D. W.; Oberhauser, J. P.; Kornfield, J. A. Macromolecules 2002, 35 (7), 2583−2594. (29) Yan, T.; Zhao, B.; Cong, Y.; Fang, Y.; Cheng, S.; Li, L.; Pan, G.; Wang, Z.; Li, X.; Bian, F. Macromolecules 2010, 43 (2), 602−605. (30) Cui, K.; Meng, L.; Tian, N.; Zhou, W.; Liu, Y.; Wang, Z.; He, J.; Li, L. Macromolecules 2012, 45 (13), 5477−5486. (31) Balzano, L.; Rastogi, S.; Peters, G. W. M. Macromolecules 2009, 42 (6), 2088−2092. (32) Balzano, L.; Kukalyekar, N.; Rastogi, S.; Peters, G. W.; Chadwick, J. Phys. Rev. Lett. 2008, 100 (4), 048302.

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (L.L.). Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS This work is supported by the National Natural Science Foundation of China (51033004, 51120135002, 51227801, and 51303166), the Fundamental Research Funds for the Central Universities (WK2310000031), the China Postdoctoral Science Foundation (2012M521233), and the 973 program of MOST (2010CB934504). The experiment is partially carried out in National Synchrotron Radiation Lab (NSRL) and Shanghai 685

dx.doi.org/10.1021/ma402031m | Macromolecules 2014, 47, 677−686

Macromolecules

Article

(33) Ogino, Y.; Fukushima, H.; Matsuba, G.; Takahashi, N.; Nishida, K.; Kanaya, T. Polymer 2006, 47 (15), 5669−5677. (34) Matsuba, G.; Sakamoto, S.; Ogino, Y.; Nishida, K.; Kanaya, T. Macromolecules 2007, 40 (20), 7270−7275. (35) Mykhaylyk, O. O.; Fernyhough, C. M.; Okura, M.; Fairclough, J. P. A.; Ryan, A. J.; Graham, R. Eur. Polym. J. 2011, 47 (4), 447−464. (36) Mykhaylyk, O. O.; Chambon, P.; Graham, R. S.; Fairclough, J. P. A.; Olmsted, P. D.; Ryan, A. J. Macromolecules 2008, 41 (6), 1901− 1904. (37) Okura, M.; Mykhaylyk, O. O.; Ryan, A. J. Phys. Rev. Lett. 2013, 110 (8), 087801. (38) Balzano, L.; Rastogi, S.; Peters, G. Macromolecules 2011, 44 (8), 2926−2933. (39) Yang, L.; Somani, R. H.; Sics, I.; Hsiao, B. S.; Kolb, R.; Fruitwala, H.; Ong, C. Macromolecules 2004, 37 (13), 4845−4859. (40) Somani, R. H.; Yang, L.; Hsiao, B. S. Polymer 2006, 47 (15), 5657−5668. (41) Zuo, F.; Keum, J. K.; Yang, L.; Somani, R. H.; Hsiao, B. S. Macromolecules 2006, 39 (6), 2209−2218. (42) Hsiao, B.; Yang, L.; Somani, R.; Avila-Orta, C.; Zhu, L. Phys. Rev. Lett. 2005, 94 (11), 117802. (43) Kumaraswamy, G.; Kornfield, J. A.; Yeh, F. J.; Hsiao, B. S. Macromolecules 2002, 35 (5), 1762−1769. (44) Kimata, S.; Sakurai, T.; Nozue, Y.; Kasahara, T.; Yamaguchi, N.; Karino, T.; Shibayama, M.; Kornfield, J. A. Science 2007, 316 (5827), 1014−7. (45) Fernandez-Ballester, L.; Thurman, D. W.; Zhou, W.; Kornfield, J. A. Macromolecules 2012, 45, 6557−6570. (46) Heeley, E. L.; Fernyhough, C. M.; Graham, R. S.; Olmsted, P. D.; Inkson, N. J.; Embery, J.; Groves, D. J.; McLeish, T. C. B.; Morgovan, A. C.; Meneau, F.; Bras, W.; Ryan, A. J. Macromolecules 2006, 39 (15), 5058−5071. (47) Zhao, B. J.; Li, X. Y.; Huang, Y. J.; Cong, Y. H.; Ma, Z.; Shao, C. G.; An, H. N.; Yan, T. Z.; Li, L. B. Macromolecules 2009, 42 (5), 1428− 1432. (48) Kugler, J.; Fischer, E. W.; Peuscher, M.; Eisenbach, C. D. Makromol. Chem. 1983, 184 (11), 2325−2334. (49) Vaidyanathan, T. K.; Vaidyanathan, J.; Cherian, Z. Dent Mater. 2003, 19 (1), 46−53. (50) Doi, M.; Edwards, S. F. The Theory of Polymer Dynamics; Clarendon Press: Oxford, 1986. (51) Liu, Y. P.; Zhou, W. Q.; Cui, K. P.; Tian, N.; Wang, X.; Liu, L. B.; Li, L. B.; Zhou, Y. G. Rev. Sci. Instrum. 2011, 82 (4). (52) Hammersley, A. P. FIT2D V9.129 Reference manual v3.1, ESRF report 98HA01T, 1998. (53) Wang, Y. Y.; Cheng, S. W.; Wang, S. Q. J. Rheol. 2011, 55 (6), 1247−1270. (54) Liu, D.; Tian, N.; Cui, K.; Zhou, W.; Li, X.; Li, L. Macromolecules 2013, 46 (9), 3435−3443. (55) Zuidema, H.; Peters, G. W. M.; Meijer, H. E. H. Macromol. Theory Simul. 2001, 10 (5), 447−460. (56) Coppola, S.; Grizzuti, N.; Maffettone, P. L. Macromolecules 2001, 34 (14), 5030−5036. (57) Wang, Y.; Boukany, P.; Wang, S.-Q.; Wang, X. Phys. Rev. Lett. 2007, 99 (23), 237801. (58) Zhao, Y.; Hayasaka, K.; Matsuba, G.; Ito, H. Macromolecules 2013, 46 (1), 172−178. (59) Bowden, P. B.; Young, R. J. J. Mater. Sci. 1974, 9 (12), 2034− 2051.

686

dx.doi.org/10.1021/ma402031m | Macromolecules 2014, 47, 677−686