Article pubs.acs.org/Macromolecules
Interplay between a Strong Memory Effect of Crystallization and Liquid−Liquid Phase Separation in Melts of Broadly Distributed Ethylene−1-Alkene Copolymers Al Mamun, Xuejian Chen, and Rufina G. Alamo* Department of Chemical and Biomedical Engineering, FAMU-FSU College of Engineering, 2525 Pottsdamer St., Tallahassee, Florida 32310-6046, United States S Supporting Information *
ABSTRACT: Narrow metallocene-made random ethylene copolymers display a strong memory effect of crystallization above their equilibrium melting point akin to the melt memory effect observed in model ethylene−1-butene copolymers. The onset temperature for self-nucleation or surviving self-seeds displays a bell shape with increasing comonomer content with a maximum at ∼2 mol % branches. Self-seeds do not survive at temperatures above the equilibrium melting point for homopolymers and copolymers either with very low branching or with a branching content >4.5 mol %. The self-seeds are associated with clusters of ethylene sequences that remain in the melt in close proximity and accelerate a subsequent crystallization, as observed by higher crystallization peak temperatures and higher nucleation density. Contrasting this behavior, commercial ethylene−1-alkene copolymers with a broad, bimodal comonomer distribution display an inversion of the crystallization rate in a range of melt temperatures where narrow copolymers show a continuous acceleration of the rate. The inversion demarcates the onset of a self-seed assisted liquid−liquid phase separation (LLPS) between comonomer-rich and comonomer-poor molecules. The interplay between number of self-seeds and chain diffusion during LLPS causes a decrease in the crystallization rate with decreasing melt temperature. When crystallites remain in the melt at temperatures Tm°copo) was initially attributed to a broader molar mass distribution than for hydrogenated polybutadienes and to the presence of stabilizers and cocatalyst residues.15 The rational is that if additives or catalyst residues act as nucleants, the effect on increasing crystallization rate may be even higher than the melt memory effect on crystallization. In such a case, the memory effect above Tm°copo on a further crystallization may not be resolved experimentally. In the present work we have extended the study of strong melt memory effect on crystallization to two types of ethylene− 1-alkene copolymers (LLDPEs). Within the first type, copolymers were synthesized with a single-site metallocene catalyst leading to a relatively narrow unimodal interchain comonomer content distribution. Studies on this series will demonstrate that the strong melt memory of the crystallizable sequence partitioning that evolves during copolymer crystallization, as described earlier for hydrogenated polybutadienes,15 is a general feature of the fusion and crystallization of narrow random copolymers regardless of the synthetic route. Moreover, the narrow metallocene set allows extending the branching content to values < ∼2 mol %, which are not accessible for hydrogenated polybutadienes. The second type of copolymers are commercial LLDPEs synthesized with either a Ziegler−Natta or with a metallocene catalyst, but with bimodal and broad interchain comonomer content distribution and broad molar mass distribution.18,19 The crystallization from different melt temperatures of the latter are unexpected and are interpreted as resulting from the interplay between the strong melt memory and the onset of liquid−liquid phase separation.
even at temperatures well above their equilibrium melting point.15 The Tc of copolymers with Mw > ∼4000 g/mol increases as the temperature of the initial melt (Tmelt) decreases and only becomes constant when cooling from above a critical Tmelt that could be as high as 40 °C above the equilibrium melting.15 This critical melt temperature is termed Tonset as it demarcates the onset of surviving self-seeds that accelerate a subsequent crystallization. The increased crystallization rate is a consequence of an increase in nucleation density, as shown by polarized optical microscopy, and was explained as a selfnucleation effect that is unique for random copolymers.15 The surviving seeds at such high temperatures were associated with clusters of ethylene sequences with little or no order that remain in the melt as a memory of the crystallizable sequence length partitioning in the initial evolution of the crystallites.15,16 These clusters are most probably molten ethylene sequences that remain in close proximity even at temperatures above equilibrium, as WAXS and SAXS diffractograms collected at these temperatures lack any sign of crystalline order. The increase in crystallization rate was observed even after holding the melt for over 30 min. Hence, the slow diffusion of these sequences back to the randomized copolymer melt was explained as a consequence of the copolymer’s melt topology that builds during crystallization15 or as a weakly segregated melt state due to the sequence partitioning.16 In the first case, the need for selecting and dragging to the crystal front suitable crystallizable sequences within the entangled copolymer molecules, and the rejection of the branches from the crystals was presumed to add topological constraints (ties, loops, knots) in the intercrystalline region of the copolymer phase structure, preventing a fast homogenization of the crystalline sequences upon melting.15 The result is a melt structure with memory (or self-seeds) of the initial sequence segregation needed to build the copolymer crystallites. The surviving self-seeds speed up a further crystallization as the step to rearrange crystalline sequences is partially bypassed. In line with this interpretation, molecular dynamics simulations have recently shown that slow relaxation of entanglements built during crystallization can explain the role of thermal history or melt memory effects in polymer crystallization.17 In our prior work we studied a series of hydrogenated polybutadienes which are models for random ethylene−lbutene copolymers with a very narrow molar mass distribution (Mw/Mn ∼ 1.1) and basically no interchain comonomer content distribution. There are very few studies of effect of melt memory on the crystallization of commercial linear lowdensity polyethylenes (LLDPEs), and in most of the available studies the effect of self-nucleation on increasing crystallization rate was observed at Tmelt below the equilibrium melting temperature of the copolymer.12,13 Although the commercial copolymers studied were synthesized with a single-site metallocene-type catalyst, the lack of strong melt memory
2. EXPERIMENTAL SECTION 2.1. Samples. The copolymers studied and their characterization data are listed in Table 1. Most of the copolymers shown in Table 1 are commercial, and all were provided by the ExxonMobil Co.20 Copolymers type A were synthesized by a metallocene-type catalyst and have narrow interchain distribution of the comonomer content; their molar mass and distribution of molar mass are very similar, and the branching content changes from ∼1 to 2.5 mol %. Three copolymers of type A are ethylene−l-hexene based, and one is an ethylene−l-butene. Copolymers type B have a broad bivariate distribution or a broad distribution of molar mass and broad distribution of the content of lhexene across the molar mass. EM 100 was synthesized in a gas-phase reactor with a metallocene catalyst; its bivariate distribution was characterized in detail in a previous work.18 NTX 141 is a commercial ethylene−1-hexene synthesized with a Ziegler−Natta-type catalyst. The molar mass and the average butyl branch content of both resins are very similar at ∼123 000 g/mol and 1.7 mol %, respectively. 2.2. Additive Removal. As received, all copolymers contained thermal stabilizers (Irganox 1010, Irgafos 168, or both). NTX 141 contains multiple additives, such as talc (an antiblock additive), slip, and melt processing aids in addition to the thermal stabilizers.20 In order to minimize or avoid the effect of the additives in the 7959
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Figure 1. (a) Room temperature FTIR spectra for Enable 2010. The inset represents the content (wt %) of Irganox in the original (solid symbol) and solvent-treated (open symbol) samples. Plotted are data for Enable 2010 (diamond), EX001 (square), NTX 141 (triangle), and Exact 3031 (circle) copolymers. (b) Cooling and subsequent melting runs at 10 °C/min for Enable 2010, original (black) and solvent treated (red).
Figure 2. (a) Room temperature FTIR spectra and (b) cooling and subsequent melting runs at 10 °C/min for NTX 141, original (black) and solvent treated (red). crystallization, all copolymers except EM 100 were subjected to recrystallization from a 0.01% w/v solution in xylene at 120 °C using acetone as the copolymer precipitant (added in a proportion of 10 times the xylene volume). Acetone is miscible with xylene and is a good solvent at room temperature for Irganox and Irgafos; hence, most of these additives remain dissolved in acetone and are separated after filtering. The copolymer crystals were filtered and repeatedly washed with acetone. They were further dried in vacuum at room temperature overnight. The content of Irganox 1010 prior to and after solvent treatment was calculated by transmission FTIR using 100 ± 10 μm thick films that were melt-pressed from the original pellets or from the dried powders. FTIR spectra were collected at room temperature using a Thermo Scientific Nicolet 6700 spectrometer equipped with a TE cooled DTGS detector. The spectral range was 400−4000 cm−1 and the resolution 2 cm−1. The absorbance of the carbonyl stretching from the ester group of Irganox 1010 (1745 cm−1), normalized to one of the CH2 wagging modes of polyethylene for thickness correction (2019 cm−1), was used to quantify the content of Irganox21−23 and to provide evidence of additive removal by the solvent treatment. An example is given in Figure 1 for Enable 2010 in the frequency region of interest. The intensity of the carbonyl absorbance is significantly lower after the solvent treatment. The content of Irganox 1010 is given in the inset of Figure 1a for original (filled symbols) and solvent-treated (open symbols) copolymers. Most open symbols fall well below the original values, indicating an effective dissolution of the additive in acetone. After dissolution and reprecipitation, it was confirmed that the
copolymer’s primary chain structure remains unchanged, as shown by the crystallization and melting thermograms of Figure 1b. While there are obvious differences in nucleation density and therefore in the crystallization peak, the melting behavior remains basically unchanged. Removal of talc, used as antiblock, from NTX 141, required a different procedure than the one discussed above for antioxidants. Talc (magnesium silicate) does not dissolve in acetone and is known to be an efficient nucleator agent for polyethylenes and polypropylenes.24−26 Thin films from the original copolymer were subjected to Soxhlet extraction with xylene using a 0.8 μm pore size thimble. After copolymer extraction, the talc particles remained in the thimble. The copolymer was precipitated from the xylene solution with acetone. Talc extraction was quantitative as shown in Figure 2a by the disappearance of the FTIR band at 1017 cm−1 characteristic of the Si− O−Si asymmetric stretching vibration. The procedure also removed Irganox as seen in the inset by the disappearance of the 1745 cm−1 absorbance. The nucleating activity of talc shifts the crystallization peak by ∼4 °C to higher temperatures while there is little difference in the melting peaks, as expected (Figure 2b). 2.3. Techniques. The crystallization and melting behaviors were followed by differential scanning calorimetry (DSC) using a TAQ2000 instrument and thermal protocols similar to those used in our prior work.15 We recorded crystallization peak temperatures and the onset of crystallization from different melt temperatures. In a standard heating/cooling protocol, the copolymer was heated at 10 °C/min from 40 °C to a temperature above the observed melting (Tmelt), holding at Tmelt for 5 min, cooling at 10 °C/min to 40 °C, holding at 7960
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40 °C for 5 min, and heating again to a different Tmelt to further repeat the same cycle. Exotherms during cooling at rates of 5, 20, and 40 °C/ min were also obtained in one of the copolymers. The temperatures of the initial melts for each cooling trial were varied in a random pattern, above and below the equilibrium melting temperature of the copolymers, so as to avoid any possible systematic error in the measured crystallization temperatures. The actual heating and cooling cycling protocols can be found in the Supporting Information (Figure S1). The DSC was calibrated for static temperature, thermal lags, and heat of fusion with indium. The TA Q2000 is connected to an intracooler to maximize heat transfer and to allow subambient temperature control. To ensure good contact, a single flat piece of polymer of about 4 mg was cut from 0.3 mm thick films and encapsulated in aluminum pans suitable for DSC work. Polarized optical micrographs were recorded using ∼50 μm thick films obtained by melt pressing the copolymer film between two microscope cover glasses. The morphology was recorded using an Olympus BX51 optical microscope equipped with an Olympus DP72 digital CCD camera and a Linkam hot stage TMS94 for temperature control. Heating and cooling cycles followed the same protocol as for the DSC experiments. The Linkam’s stage temperature was calibrated recording the change in light intensity of the copolymers during cooling and during heating in the polarized optical microscope, in reference to the DSC enthalpic change at the same cooling or heating rate. High-resolution AFM images (512 × 512) were obtained using a Bruker Multimode V8 scanning probe microscope in tapping mode with Bruker silicon-coated cantilevers with resonant frequency of ∼75 kHz. Free surface films suitable for room temperature AFM imaging were prepared placing a ∼100 μm thick film between microscope cover glasses and were melted on a Linkam hot stage TMS94. The top glass was then removed by quenching the sandwich in liquid nitrogen. After melting at the required temperature, the free surface film was subsequently cooled at 10 °C/min to 40 °C. In an attempt to extract molecules with high comonomer content that may be segregated from more crystalline components, the free-surface films were subsequently etched in n-hexane at room temperature for 24 h and further dried for 48 h at ambient temperature. For TEM imaging, ∼0.5 mm thick films were melted and crystallized in the Linkam hot stage. The films were removed from the glass holder, cryo-fractured, and stained with ruthenium tetroxide vapor for 3 h.27 They were then thin-sectioned at room temperature. TEM images were taken at 200 eV accelerating voltage in bright field mode using a FEI Tecnai F20 field emission instrument. Images were recorded on a 1K pixel CCD camera. TREF and GPC distribution profiles were carried out in instruments commercialized by PolymerChar, Valencia, Spain. For TREF, the solvent was o-dichlorobenzene (oDCB), the crystallization rate was 0.5 °C/min from the stabilization temperature of 95 to 25 °C (or to 0 °C for copolymers type B), and the subsequent solvent elution rate was at 0.5 mL/min from 25 °C (or from 0 °C) to 110 °C at a rate of 1 °C/min. The solvent for GPC-IR was tetrachlorobenzene (TCB) stabilized with 300 ppm of BHT at a flow rate of 1.0 mL/min.
a relatively narrow interchain distribution of the comonomer content, but broader than for HPBDs. The difference in interchain comonomer content distribution with that of HPBDs is shown in the TREF profiles of Figure 3. The
Figure 3. TREF profiles of type A copolymers with unimodal and relatively narrow interchain comonomer content distribution. The profile of a hydrogenated polybutadiene (P108) with 2.2 mol % ethyl branches is included for comparison. TREF profiles are normalized by equal area.
width in elution temperatures is associated with the interchain comonomer composition. Copolymers listed type A in Table 1 have unimodal, relatively narrow TREF patterns, yet the patterns are significantly broader than the analogue for a hydrogenated polybutadiene with 2.2 mol % ethyl branches (P108). The elution temperature range in the TREF profile of P108 is less than half of the full width at half-height of the commercial narrow copolymers with equivalent comonomer composition. Furthermore, narrow copolymers with a branching content lower than 2 mol %, such as EX 001 and Enable 2010, were investigated to probe the behavior of a copolymer microstructure that approaches the unbranched chain. Although comparative TREF profiles are useful to infer differences in comonomer composition distribution, the combination of molar mass and comonomer content distribution is not resolved in the 2D profiles of Figure 3. The distribution of comonomer (or branches) across the molar mass distribution was elucidated by molar mass fractionation and further collection of GPC and TREF profiles of each of the fractions.18,30 The 3D bivariate distribution was further obtained by combining normalized GPC and TREF profiles of each fraction in proportion to their mass fractions. A representative bivariate distribution of type A copolymers is given in Figure 4 for Enable 2010. In spite of a molar mass distribution (PDI ∼ 4) broader than the ideal value corresponding to a single site copolymerization (PDI = 2), the interchain branching distribution from the 3D bivariate is relatively narrow for all type A copolymers, as also indicated by the surface plot of Figure 4. This detailed branching and molar mass characterization is essential to understand the melt memory effect on crystallization behavior. Prior to additive removal, the crystallization temperatures of the “as received” type A copolymers are independent of the temperature of the initial melt down to melt temperatures near the observed final copolymer melting. However, after additive removal, the exotherms shift to higher values when cooled from Tmelts even above the equilibrium melting point, in consonance
3. RESULTS AND DISCUSSION 3.1. Copolymers with Narrow Comonomer Composition Distribution. Detailed work on melt memory of crystallization in random ethylene copolymers was carried out in a prior study using hydrogenated polybutadienes (HPBDs), which are models for ethylene−l-butene copolymers.15 These copolymers were synthesized by anionic polymerization and have no interchain branching composition heterogeneity. Because intermolecular composition heterogeneity is inherent to all commercial LLDPEs at different degrees,28,29 before presenting data for the broadly distributed copolymers of main interest in the present study, we first study the crystallization of commercial copolymers synthesized with a single-site metallocene catalyst as reference of the behavior of copolymers with 7961
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the free energy of nucleation and the acceleration of crystallization. In Figure 5, a vertical line is drawn over the points with constant crystallization to extract the onset of melt memory (Tonset) or the critical value of Tmelt at which any remaining seeds dissipate and the melt becomes homogeneous. It is clear from the data of Figure 5 that Tonset depends strongly on comonomer content. The copolymer with the lowest branching content (1.08 mol %) displays weak melt memory as the increase in Tc,peak from melts above the equilibrium melting temperature of this copolymer is only ∼0.3 °C (yet, this small increase is distinctive and reproducible, as can be seen in a more expanded scale in Figure S3). The increase in Tc,peak raises up to 3 °C with increasing branching up to about 4 mol %, and melt memory above Tm°copo vanishes in copolymers with >4.5 mol % branches due to low levels of crystallinity, as previously described.15 We recall that the strong melt memory that shifts crystallization temperature to higher values is only observed when Tmelt is approached from below, and hence it is associated with the development of the copolymer’s crystalline state. The shift to higher crystallization temperatures is not observed above Tm°copo when Tmelt is approached from above. The strength of melt memory is defined as the difference between Tonset and Tm°copo. The greater this difference, the higher is the temperature above Tm°copo where seeds still survive. This difference is plotted against mol % branches in Figure 6. Open symbols are data for hydrogenated poly-
Figure 4. 3-D bivariate distribution for Enable 2010.
to the behavior of hydrogenated polybutadienes (Figure S2).15 The crystallization peak temperatures (Tc,peak) plotted as a function of the initial melt temperatures (Tmelt) are given in Figure 5. Here, we observe the distinct shift of Tc,peak toward
Figure 5. Temperature of the initial melt vs crystallization temperature for narrow copolymers. The branching content is indicated. Vertical lines are drawn over data with constant Tc,peak, and horizontal lines demarcate Tmocopo of each copolymer.
Figure 6. Plot of difference between onset temperature for melt memory and equilibrium melting temperature against branching content of narrow ethylene copolymers and hydrogenated polybutadienes (HPBDs), the latter from ref 15.
higher temperatures even at Tmelt > Tm°copo, thus confirming that the apparent lack of strong melt memory of the original copolymers is due to the effect of additives in the crystallization. The observed crystallization temperature upshift is independent of cooling rate in a range of 5−40 °C/min and is attributed to enhanced nucleation since the melting endotherms are independent of the initial Tmelt. Hence, the copolymer crystallization kinetics are strongly affected by the state of the initial melt even above Tm°copo, in analogy to the behavior of HPBDs.15 Some type of self-seeds remain in the melt up to temperatures close to 170 °C and disappear at a critical temperature (Tonset) above which the crystallization kinetics are unchanged. The self-seeds are associated with clusters of ethylene sequences from the initial crystallites that, although they may be molten at temperatures above Tm°copo, remain in close proximity, thus facilitating a further gathering of ethylene sequences on a subsequent cooling. The result is a decrease of
butadienes from our previous work,15 and closed symbols represent data for the narrow metallocene copolymers studied here. There appears to be minimal difference in melt memory between the ethyl and the butyl branch as the ethylene−1butene (EXR 705) and the ethylene−1-hexene data fall on the HPBD line. Moreover, the strength of melt memory shows a bell-shaped dependence with branching content that is explained by two competing structural parameters. Increasing branching up to ∼2 mol %, the number of ethylene sequences increases and so does the buildup of topological constraints in the intercrystalline region that prevent fast diffusion of crystalline sequences upon melting; hence, with increasing number of sequences a higher number of seeds survive at temperatures above Tm°copo, increasing the value of Tonset. 7962
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copolymers dissolve in a range of temperatures of about 20 deg, while for type B the range is 4 times broader and bimodal. These copolymers display two major components with respect to 1-hexene content with complex distribution of 1-hexene over the distribution of chain lengths.18,19 The comonomer-rich component dissolves between 30 and 80 °C and peaks at ∼60 or 75 °C for EM 100 and NTX 141, respectively. The component with a low content of 1-hexene dissolves between 80 and 100 °C and peaks for both copolymers at ∼90 °C. In reference to 13C NMR data of fractions from these copolymers, the average branching content of comonomer-rich and poor components is 1 and 7 mol %, respectively.18 A detailed distribution of the comonomer content across the molar mass distribution is given by the 3-D bivariate plots in Figure 8, which were also constructed from GPC and TREF profiles of molar mass fractions that were obtained by a solvent−nonsolvent procedure from the original copolymers.18 In spite of the difference in catalysis, metallocene type for EM100 and Ziegler−Natta for NTX 141, the shape of the overall bivariate is remarkably similar; the major difference between the resins being in the molar mass of molecules with the highest comonomer content, or molecules eluting in the range of 25−50 °C which are of relatively low molar mass in the ZN-derived copolymer (NTX 141) and of high molar mass in the metallocene (EM 100). For both copolymers, the component eluting at ∼90 °C is made of molecules with about the same content of comonomer (∼1 mol %) and broadly distributed in molar mass. Conversely, the comonomer-rich component is broadly distributed in both comonomer content and molar mass. The content of butyl branches in the comonomer-rich component decreases from ∼7 to 1.5 mol % with increasing molar mass. From the overall distribution of EM100 and NTX 141, the comonomer-rich component accounts for ∼48 mass %, and the rest (52 mass %) are molecules with low comonomer content. It is perceived that the gradual change in 1-hexene content across a broad molecular weight of the comonomer-rich component provides more intimate cocrystallization and more effective intercrystalline links and that both are possibly responsible for the enhanced mechanical properties of these copolymers compared to binary blends.18 With a clear understanding of the bivariate distribution in these complex systems, we can now analyze their crystallization
Moreover, beyond 2 mol % branches, the number of sequences increases but the crystallinity level decreases sharply, reducing the number of intercrystalline topological constraints; the seeds that may remain at and above Tm°copo decrease accordingly and eventually dissipate. Lack of melt memory from crystallization above Tm°copo occurs at branching contents >4.5 mol % or for zero or negative values of Tonset − Tm°copo. The increased nucleation density in narrow metallocene copolymers due to self-nucleation from melts > Tm°copo results in finer spherulitic structures, as observed by polarized optical microscopy, while WAXS and SAXS of self-nucleated melts lack any signature of molecular order. It is then concluded that a small increase in the breadth of the bivariate distribution of metallocene copolymers with respect to the distribution of HPBDs does not alter the memory effect of crystallization in melts above the equilibrium point. Crystallites do not survive above Tm°copo, but clusters of molten long ethylene sequences must remain as long-lived transients. 3.2. Copolymers with Broad Bivariate (Molar Mass− Comonomer Composition) Distribution. Figure 7 displays
Figure 7. TREF profiles of type B copolymers with bimodal and broad interchain comonomer content distribution.
TREF profiles of the copolymers with a broad distribution of the comonomer content (listed as type B in Table 1). The contrast in dissolution temperatures with those of narrow copolymers (Figure 3) is clearly drastic. Narrowly distributed
Figure 8. 3-D bivariate distribution for NTX 141 (left) and EM 100 (right). Elution temperature is inversely proportional to 1-hexene content. 7963
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Figure 9. Temperature of the initial melt (Tmelt) versus crystallization peak temperature for broad, type B copolymers: (a) NTX 141 and (b) EM100. Closed circles: Tmelt is approached from below. Open symbols: Tmelt is approached from above.
Figure 10. Plot of heat flow against temperature for cooling from melt temperatures in a range of 200−126 °C and subsequent heating at 10 °C/min for EM100. The inversion in crystallization rate with lowering the initial melt temperature from region B to C (red to blue exotherms) is emphasized in the more expanded exotherms of the inset.
corresponds to the onset of incomplete melting as shown in the cooling and heating thermograms in region D (Figure 10); therefore, decreasing further Tmelt down from 125 °C, the expected increase in crystallization temperature is observed again. In region D crystallites that remain in the melt speed up a subsequent crystallization, and the Tc,peak again increases substantially. Regions A, B, and D are the expected behavior for random copolymers.15 Region C is unexpected and must be associated with a drastic change in melt structure due to the complex bimodal distribution of comonomer content in these resins. It seems that region C, where Tc,peak undergoes an inversion in the rate with decreasing Tmelt, is a unique feature of bimodal and broadly distributed copolymers. As for narrow copolymers, regions B and C vanish, and only region D prevails when the temperature of the initial melt (Tmelt) is approached from above (open symbols in Figure 9b). Hence, the shifts in crystallization rate (Tc,peak) in regions B and C must be associated with the sequence partitioning that takes place in the development of the copolymer’s crystalline state. Representative thermograms for each region are given in Figure 10. Notice that independent of the drastic differences in crystallization with lowering Tmelt, the melting endotherms are essentially unchanged, thus indicating that the increase (region
behavior in reference to the initial melt topology or melt memory above the equilibrium melting. The nucleating activity of talc is high in the original NTX 141 resin, such that selfnucleation is not resolved experimentally, and data of Tmelt vs Tc fall on a straight line until temperatures well below the equilibrium value (Figure S4). The shift of Tc,peak to higher temperatures is shown from melts above equilibrium when talc is removed from NTX 141, as expected. These data are plotted in Figure 9 for EM100 and additive removed NTX 141. The crystallization behavior of both broad copolymers, especially the metallocene-based EM100, given by the variation of Tc,peak vs Tmelt in Figure 9, is quite remarkable and fully reproducible. As a function of decreasing Tmelt we divide the data of these figures in four regions. With decreasing Tmelt down to 170 °C, region A, Tc,peak is constant and corresponds to crystallizations from a homogeneous, one-phase randomized melt. From Tmelt 170 to 150 °C, region B, the peak crystallization temperature shifts by ∼3 °C to higher values indicative of a strong melt memory effect of crystallization, as found in the model HPBDs and narrow metallocene-made copolymers. However, decreasing Tmelt from 150 to 125 °C, the crystallization temperature decreases by almost 2 deg (region C), which is not expected. Lower Tmelt, for example 123 °C, 7964
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in region C phase separation is weak and slow because there are no partitioned ethylene sequences to aid with the process, or LLPS has not started yet at the relatively fast cooling rate of 10 °C/min used. Indeed, even in the binary model blends studied by Han et al., very long holding times were required to obtain spatially distinguished domains at temperatures deep in the LLPS region (see discussion below).34−37 The implication is that LLPS in type B copolymers is a lot more efficient when the binodal is approached from below. The decrease in crystallization rate with decreasing Tmelt in region C is a feature in line with the decrease in nucleation density observed in model binary blends crystallized from temperatures in the LLPS region.35−37 The decrease of Tc,peak with decreasing Tmelt in region C points to an interplay between the self-seeds at those melt temperatures and the extent of LLPS. To facilitate the discussion, a schematic LLPS diagram is given in Figure 11. For the copolymers of interest with
B) and decrease (region C) in the crystallization rate, associated with Tc,peak, are mainly due to changes in nucleation density. Selected exotherms of regions A, B, and C in a more expanded scale are given in the inset of Figure 10 to point out more clearly the inversion in crystallization rate with decreasing Tmelt from region B to C. To understand the change in melt topology that triggers this sudden inversion in the rate of crystallization, we focus first on the bimodal nature of NTX 141 or EM100both are analogous to a 50/50 blend of comonomer-poor and comonomer-rich components, the latter with a broad range of interchain comonomer content from 1.5 to ∼13 mol % 1hexene18,19and recall extensive studies on the phase structure of melts in binary blends of model random copolymers with comonomer contents covering this range.31−42 Studies using neutron scattering demonstrated that melts of binary mixtures of ethylene copolymers with a difference in branching content >4 mol % are phase separated.32,43−45 Furthermore, the phase diagram of blends of ethylene−1-hexene (1 mol % branches) and ethylene−1-butene (∼7 mol % branches) shows an upper critical solution temperature (UCST) at 146 °C for the 56/44 compositional blend.34−36 Since the UCST coincides with the temperature of the initial melt where the inversion in the crystallization rate is first observed in region C, we attribute the drastic change in crystallization to the onset of liquid−liquid phase separation. In region A the crystallization rate occurs from a single phase homogeneous melt and is dominated by the low comonomer content molecules or those with the longest crystallizable sequences. Since in region A Tmelt is above Tonset, the Tc,peak is constant as expected. In region B, the melt is essentially one phase but heterogeneous; here clusters of ethylene sequences remain in close proximity retaining some of their initial partitioning. Hence, compared to region A, the crystallization rate from region B increases. The rate increases because the clusters facilitate the ethylene sequence selection in a subsequent crystallization. Moreover, in region C the crystallization rate decreases because liquid−liquid phase separation precedes crystallization. We propose that the decrease of crystallization rate in region C is so sharp because in approaching Tmelt from below the seeds (clusters of ethylene sequences) assist with the molecular segregation dynamics in the thermodynamic liquid−liquid phase transition. The thermodynamic drive for liquid−liquid phase separation (LLPS) in region C is stronger than the kinetic diffusion drive of region B; hence, when the UCST boundary is approached from below, two phase-separated domains form quickly in the melt assisted by the clusters of ethylene sequences that remain at that melt temperature. One melt domain is rich in molecules with low comonomer content, and the second is composed mainly by comonomer-rich molecules that segregate from the low comonomer ones. Molecular diffusion during LLPS dissipates some of the ethylene clusters, leaving a lower number of nuclei in the melt. The decrease in number of surviving nuclei compared to the number of nuclei at temperatures just above the UCST explains the decrease in crystallization rate (lower Tc,peak) in region C. That self-seeds play a role in the dynamics of the liquid− liquid phase transition is demonstrated experimentally by the open symbols of Figure 9b. Approaching Tmelt from above gives the Tc,peak observed in region A; i.e., the crystallization rate is unchanged. As no seeds remain in the melt, the crystallization from region B occurs from a homogeneous one phase melt, and
Figure 11. Schematic phase diagram for bimodal, broad ethylene−1hexene type B copolymers. f denotes mass fraction of comonomerpoor molecules in the copolymer. The dotted lines are estimates of the equilibrium melting temperatures from data of binary blends, and the dashed line inside the bimodal demarcates the mass fraction of the bimodal copolymers analyzed.
composition f of 0.5, when the temperature is approached from below to a region inside the binodal but near the UCST (∼150 °C), the number of molecules diffusing toward phase segregated domains is relatively low because the temperature is too close to the binodal; hence, a large number of self-nuclei (crystalline sequences that remain molten at Tmelt but in close proximity) remain in the majority phase, and the effect on overall crystallization rate with respect to region B is relatively small. However, when Tmelt is further from the UCST, for example at 130 °C in region C, heating into the LLPS is deep into this region, and because the number of seeds from crystalline sequences at this lower temperature is so large, segregation into the two phase domains is facilitated by these seeds or by the large number of already partitioned sequences; accordingly, molecular diffusion dissipating the seeds is very efficient. The result is a reduced or minimal number of seeds left in the LLPS domains and, as a consequence, a slower crystallization rate in spite of a decrease in Tmelt. The situation changes radically when seeds with a true crystalline nature remain at Tmelt < 125 °C. At these temperatures the remaining crystallites lower the free energy barrier for nucleation, thus accelerating again crystallization. Accordingly, crystallization rates in region D increase with decreasing Tmelt as shown in Figure 9b. 7965
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Figure 12. Optical micrographs taken at 40 °C of EM100 crystallized during cooling from initial melts of (a) 200, (b) 170, (c) 162, (d) 150, (e) 140, (f) 130, and (g) 125 °C. As shown by the thermal sequence on each micrograph (200−40−Tmelt−40), prior to cooling, the Tmelt was approached from below. The scale bar represents a length of 20 μm.
Figure 13. (a) Change of Tc,peak with annealing time at the indicated Tmelt for EM100, and polarized optical micrographs (POM) taken at 40 °C after annealing at 135 °C for 5 min (b) and for 20 h (c). As in Figure 12, Tmelt was approached from below.
polarized optical microscopy.15 We now focus on the crystalline morphology of bimodal broadly distributed copolymers treated under the same thermal protocol used to record the data of Figure 9. Micrographs for EM 100 recorded at 40 °C after
3.3. Morphology. The increase in nucleation density of narrow copolymers when crystallized from melts that keep memory from the crystalline ethylene sequence partitioning is akin to that observed in hydrogenated polybutadienes by 7966
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cooling at 10 °C/min from different Tmelt are shown in the composite of Figure 12. The difference in brightness of the images of Figure 9a−g indicates that nucleation density first increases (region B) and further decreases (region C), in agreement with the postulated change in self-seeds when the melt crosses the LLPS region. Micrographs a and b display the same coarse spherulitic morphology as they are obtained from homogeneous one phase melts at 200 and 170 °C, respectively. As expected, the number of nuclei increases drastically, and the spherulitic morphology becomes finer and more profuse when seeds remain in the one-phase heterogeneous melt (region B), such as micrographs c and d taken from melt temperatures of 162 and 150 °C, respectively. Examples of the crystalline structure formed from the two-phase region are given in micrographs e and f for Tmelt of 140 and 130 °C, respectively. Here the morphology becomes coarser than in region B, in agreement with the proposed reduction in number of crystal nuclei with decreasing Tmelt in region C and denoting the inversion of the crystallization rate. In region C the crystalline density is uniform, lacking brighter boundaries that may indicate a preferential crystal nucleation at the liquid−liquid interphase as proposed earlier for model binary mixtures.36,37 To gain additional insight into the predicted differences in melt dynamics between regions B and C, time-dependent experiments were carried out in both regions. The change in Tc,peak with time after annealing at 135 °C (region C), 150 and 160 °C (region B) are given in Figure 13a. In the two liquid phase region, Tmelt of 135 °C, Tc,peak decreases exponentially with annealing time, which is consistent with the proposed dissipation of seeds via chain diffusion. Accordingly, after 20 h annealing, a more coarse texture is observed in the POM image of Figure 13c, consistent with less nuclei, compared to the image after 5 min annealing (Figure 13b). Conversely, the decrease in Tc,peak after annealing at 150 or 160 °C in region B is significantly less pronounced, in agreement with slower diffusion kinetics in this region.15 The micrographs of Figures 12 and 13 lend support to the notion that the number of seeds initially present in the melt, as a memory of the sequence partitioning of copolymer crystallization, is the major contributor to crystal nucleation for crystallization from the two-phase region for these broadly distributed systems, rather than molecular fluctuations at the boundary between the two phases. At Tmelt of 145 °C the extent of LLPS is low, and many of the initial seeds remain to accelerate crystallization as seen in the micrographs. Lowering the temperature in the LLPS region, but above the crystal transition, for example 130 °C, the extent of LLPS is much greater according to the binary phase diagram of Figure 11; hence, a larger number of seeds are destroyed via molecular diffusion during spinodal decomposition and concentration fluctuations between the two phases. In this process, some of the partitioned crystalline sequences that remain in close proximity in the melt diffuse and mix with other sequences toward the comonomer-rich domain. This type of molecular diffusion is very slow in region B (one phase melt) due to the accumulation of branches, ties, knots, and other molecular constraints in the intercrystalline region. However, in spite of the topological constraints in the intercrystalline region, molecular diffusion in region C is relatively fast due to the large thermodynamic drive in the two-phase binodal (stronger than the kinetic drive in region B). Hence, in the strong phaseseparated region, the initial number of seeds diminishes, and
crystal nucleation decreases accordingly because such seeds are the most effective to nucleate copolymer crystallites. Fitting the temporal decrease of crystallization data of Figure 13 (Tmelt = 135 °C), after normalization, with a simple exponential decay [Tc,peak = exp(−t/τ)], we obtain a time constant (τ) of 261 min, and for dimensions of the self-seed clusters of the average length of the initial lamellae (∼0.5 μm long), the ethylene sequences would have to diffuse over a surface of ∼(0.25 μm)2 to dissipate the cluster. With these data, the diffusion coefficient of EM 100 clusters at 135 °C during LLPS is ∼4 × 10−14 cm2/s. This value is 3 orders of magnitude smaller than the diffusivity of the single random coil (D ∼ 10−11 cm2/s) at comparable molar mass and temperature.46−48 Nonetheless, a slower diffusion than the value for the isolated random coil is expected in view of the constrained topology of the intercrystalline regions discussed above. All attempts to image the two phase domains in melts of EM 100 or NTX 141 using phase contrast microscopy were unsuccessful, even after long time annealing in the two-phase region. We attribute this lack of phase domain resolution to a very small difference in refraction index between the phases and to the formation of very small comonomer-rich domains (4 mol % butyl branches dissolve in n-hexane. Moreover, the effectiveness of n-hexane extraction is limited by the degree of mixing and cocrystallization between molecules with different comonomer content in the minority phase. It is clear from the TEM of 7967
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the homogeneous melt at 200 °C (Figure 15a), the darker regions increase substantially in the image of Figure 15b, when crystallization takes place from a two-phase melt. The n-hexane etching was even more effective in enhancing contrast between low and high crystallinity regions when the two-phase melt was annealed for a longer time prior to crystallization, as shown in the image of Figure 15c. Together with DSC, POM, and TEM images, the differences in nucleation and crystalline texture observed in the AFM images support the proposed phase diagram for the commercial broadly distributed copolymers and offers strategies to control the melt phase structure of these copolymers, rate of solidification, and properties of the crystalline state that depend on the initial state of the melt.
4. CONCLUSIONS Narrowly distributed metallocene-made random ethylene−1alkene copolymers display memory of the copolymer crystalline sequence partitioning at melt temperatures well above the equilibrium melting point (Tm°copo). The memory is associated with self-seeds made of clusters of ethylene sequences that remain in close proximity at high temperatures. Even above Tm°copo, some crystalline sequences are unable to diffuse back to the initial randomized copolymer melt due to a constrained chain topology in the intercrystalline region. The self-seeds accelerate a subsequent crystallization. The critical melt temperature at which seeds do not survive (Tonset) is the onset of melt homogeneity. Above Tonset the crystallization rate is independent of the temperature of the melt. For comonomers excluded from the crystalline regions, Tonset displays a bell shape with increasing branching content with a maximum at ∼2 mol %. Commercial copolymers with a broad, bimodal bivariate distribution display an inversion of the crystallization rate, associated with the peak crystallization temperature (Tc,peak), in a range of melt temperatures where narrow copolymers show a continuous acceleration of the rate. Approaching Tmelt down to ~170 °C from below, Tc,peak is constant as the crystallization takes place from a one-phase homogeneous melt free of selfseeds. From Tmelt of 170 to ∼150 °C Tc,peak increases, or the crystallization rate accelerates substantially due to self-seeds that remain in the melt akin to the behavior of narrow copolymers. Furthermore, for Tmelts from 150 down to 125 °C, Tc,peak decreases, denoting the inversion in the crystallization rate. The inversion in the rate demarcates the onset of a selfseed assisted liquid−liquid phase separation (LLPS) between comonomer-rich and comonomer-poor molecules. Decreasing Tmelt in the LLPS region, an increasingly larger number of selfseeds dissipate as long ethylene sequences that make up these
Figure 14. Thin section TEM micrographs of EM100 cooled to room temperature from (a) 200 °C (5 min), (b) 135 °C (5 min), and (c) 135 °C (20 h) at two magnifications. The bar scale represents 0.2 and 0.1 μm in (a) and (b) and 0.5 and 0.2 μm in (c). Darker areas in (c) correspond to comonomer-rich phase-separated domains in the melt. For all the images the initial melt was reached from below.
Figure 14c that some lamellae form in this phase, indicating the presence of copolymer molecules with
