Ethylene

Article pubs.acs.org/Macromolecules

Direct Formation of Different Crystalline Forms in Butene-1/Ethylene Copolymer via Manipulating Melt Temperature Yaotao Wang, Ying Lu, Jiayi Zhao, Zhiyong Jiang, and Yongfeng Men* State Key Laboratory of Polymer Physics and Chemistry, Changchun Institute of Applied Chemistry, Chinese Academy of Sciences, University of Chinese Academy of Sciences, Renmin Street 5625, 130022 Changchun, P. R. China ABSTRACT: The crystallization behavior of a butene-1/ethylene random copolymer with 9.88 mol % ethylene counits was investigated by means of differential scanning calorimetry, wideand small-angle X-ray scattering, and polarized optical microscopy. Unlike in its homopolymer counterpart which crystallizes always into a metastable form II from the melt state, the random copolymer was found to crystallize either into form II or stable form I′ directly from its melt state after being cooled down. The occurrence of either crystalline form only depended on the temperature where the crystalline material was molten before cooling down. Even though the material was brought to temperatures higher than the equilibrium melting temperature, heterogeneities of segmental segregation character were preserved which promoted massive nucleation of form I′ crystallites, which makes it possible that the material is able to crystallize into pure form I′. Only if when the melt temperature was high enough where all heterogeneities of the above-mentioned character were erased can the material be crystallized into pure form II. The effect is found independent of the previous crystalline form, meaning that the helical conformation of chains in the heterogeneous melt does not affect the nucleation of stable form I′.

■

INTRODUCTION Polymers are long chain molecules of random coil structure entangled together in the molten state. In most cases, such highly entangled polymer melt can crystallize into a semicrystalline state with layer-like lamellar crystallites stacking together with entangled amorphous chain segments in between.1 The crystallization behavior of polymers strongly depends on thermal history. When a polymer is melted just above the melting temperature or well below the equilibrium melting temperature (Tm°), it may leave small traces of surviving crystallites, oriented segments, or less-entangled polymer chains,2−9 resulting in higher crystallization rates or crystallization temperatures, smaller spherulites, or polymorphic modifications upon followed crystallization.10−16 Such a phenomenon is often addressed as “melt memory” which is an important issue in researching polymer crystallization and controlling the morphology and property of polymer in industrial processes.17,18 The memory effect has been studied as a function of many experimental conditions, such as heating rate, melt temperature, holding time, and others by means of dilatometry, differential scanning calorimetry (DSC), X-ray scattering, infrared spectroscopy (IR), and many other methods. Several interpretations of memory effect were suggested by some researchers. Many studies on bulk polymer systems considered that the memory was due to infusible heterogeneous nuclei which vanished with time at a certain temperature.3,19,20 Ziabicki and Alfonso18 interpreted this nucleus picture in terms of small “atomic © XXXX American Chemical Society

clusters” that may retain from previous crystallites and can relax with time, which determines the initial number of crystal nuclei and the initial rate of thermal nucleation. Recently, the role of melt state in polymer crystallization attracts much attention, whose structure and dynamics may be influenced by a previous crystal structure and then facilitate subsequent nucleation and/ or growth of crystals. Rastogi et al.21−25 have concentrated their efforts on understanding melting and the dynamics of the firstorder transition toward specific distribution of entanglements. Hikosaka and Yamazaki26,27 produced a “disentangled” preordered melt of polyethylene by slow melting of chainextended crystals and observed that the nucleation rate decreased with the increase of apparent entanglement density. They argued that the origin of memory effects basically arose from the change of the state of entanglement during annealing rather than the seeding of primary nuclei. Luo et al.28 reported a direct correlation between the state of entanglement in a supercooled polymer melt and its crystallization behavior and provided a direct evidence for memory effect due to persistence of disentangled regions after melting of single crystals. Alamo and Hu29,30 researched the melt memory of random copolymer by experiments and dynamic Monte Carlo simulations. They attributed the memory effect to the sequence-length Received: September 24, 2014 Revised: November 22, 2014

A

dx.doi.org/10.1021/ma5019796 | Macromolecules XXXX, XXX, XXX−XXX

Macromolecules

Article

number-average (Mn) and weight-average molecule weight (Mw) are 69 and 153 kg/mol as was measured at 150 °C by a PL-GPC 220 type high-temperature gel permeation chromatography. 1,2,4-Trichlorobenzene (TCB) was employed as the solvent at a flow rate of 1.0 mL/ min, and the calibration was made by polystyrene standard Easi-Cal PS-1 (PL Ltd.). The concentration of ethylene counits is 9.88 mol %, and the fraction of isotactic pentad is 87.1%, determined by 13C nuclear magnetic resonance at 90 °C in the 1,2-dichlorobenzene-d4 solution. For simplicity, the material is termed as PBcoE10 in the following discussions. The crystallization and melting behaviors of this PBcoE10 were measured by a DSC1 Stare System (Mettler Toledo Instruments, Swiss) under a nitrogen atmosphere (50 mL/min). The temperature scale was calibrated using indium as a standard to ensure reliability of the data acquired. The main thermal protocols are present in Schemes 1 and 2 where the range of melt temperature is from 130 to 180 °C, and the detailed processes are described in the following part.

segregation in the heterogeneous melt above the equilibrium melting temperature. Polymorphism is a ubiquitous phenomenon in semicrystalline polymers, where polymers can solidify in different crystalline structures. The type of crystalline structure can be strongly influenced by melt memory effect. Cho et al.13,14 investigated the melt memory effect of β-isotactic polypropylene based on the existence of locally ordered α-form crystallites in the melt. Lower hold temperatures and shorter hold time led to samples rich in α-form, whereas higher temperatures and longer time favored β-form during followed crystallization. De Rosa31 and Sorrentino32 investigated the crystallization behavior of syndiotactic polystyrene and obtained the pure β-form only after prolonged heat treatment at high temperatures, ensuring the disappearance of α-form nuclei. Cavallo et al.33 observed that several temperature regions can be distinguished with decreasing the self-seeding temperature in polybutene-1 (PB-1), such as the formation of pure form II crystallites, concomitant crystallization of form II and form I′, and only form I′ crystallized eventually. Other polymorphic samples, such as poly(pivalolactone),34 poly(butylene adipate),35 and poly(vinylidene fluoride),36 were also investigated under different thermal treatments. However, the mechanism of memory effect is still under debate. The polymorphic polymers could be potential candidates aiding to understand the memory effect of semicrystalline polymers further. As a typical polymorphic polymer, PB-1 exhibits four crystal modifications depending on the formation conditions: form I, twined hexagonal with a 31 helix; form I′, untwined hexagonal with a 31 helix; form II, tetragonal with a 113 helix; and form III, orthorhombic with a 41 helix.37−40 Form II is a metastable structure, which forms when the polymer crystallizes from melt at atmospheric pressure. Form II crystallites can transform into the thermodynamically stable form I when held at room temperature.41−44 This transformation process can be accelerated by stretching,45 and the reversible process was also observed upon stretching at the elevated temperature due to a stress-induced melting and recrystallization process.46 Form I′ crystals can be formed directly under peculiar conditions, like in ultrathin film,47 self-seeding,33,48,49 high pressure,50−53 and in polymers with different concentration of stereodefects54 or counits.55−58 Form III and form I′ have been found to form in solution crystallization, depending on the solvent, concentration, and the crystallization temperature59−61 which can be transformed into form II by stretching at higher temperature.62 Up to now, the exact mechanism of direct crystallization of form I′ from molten state of PB-1 and its copolymers remains elusive. In this study, we report an exceptional crystallization behavior in a butene-1/ethylene copolymer from its bulk molten state. Depending on the melt temperature of a previous crystalline sample, the butene-1/ethylene copolymer can crystallize into pure form I′, form II, or a mixture of the two forms, irrespective of the previous crystalline forms. The results can be understood as a consequence of facilitating extensive nucleation of form I′ crystals in low-temperature heterogeneous melt.

■

Scheme 1. Schematic Illustration of the Dynamic Thermal Treatment Applied to PBcoE10 in Stable Form I with Different Melt Temperatures

Scheme 2. Schematic Illustration of the Dynamic Thermal Treatment Applied to the PBcoE10 Sample

To investigate the evolution of crystalline structure during the same thermal protocols, wide-angle X-ray diffraction (WAXD) and polarized optical microscopy (POM) experiments were both carried out. WAXD experiments were measured in a modified Xeuss system of Xenocs France with the aid of a semiconductor detector (Pilatus100 K, DECTRIS, Swiss) attached to a multilayer focused Cu Kα X-ray source (GeniX3D Cu ULD, Xenocs SA, France), generated at 50 kV and 0.6 mA. A piece of PBcoE10 sample was tightly wrapped with a thin aluminum foil in order to promote thermal conductivity, and the temperature was controlled by a portable heating device (TST350, Linkam, UK) installed at the setup. The wavelength of the X-ray radiation was 0.154 nm. The sample-to-detector distance was 178 mm, and the effective range of the scattering angle 2θ was 5°−25°. Each

EXPERIMENTAL SECTION

The random butene-1/ethylene copolymer was produced using Ziegler−Natta catalysts and kindly provided by BASELL Polyolefins. The melt flow rate (MFR) is 40 g/10 min (190 °C/2.16 kg), and the B


Macromolecules

Article

WAXD pattern was collected within 5 min, background corrected, and normalized using the standard procedure. The crystallization process were also followed by using polarized optical microscopy equipped with a 20× objective lens (Zeiss Axio Imager A2m, Carl Zeiss, Germany) and combined with a heating device (THMS 600, Linkam, UK) to control the thermal history. To investigate the microstrucutral information after isothermal crystallization, small-angle X-ray scattering (SAXS) measurements were performed at beamline 1W2A, BSRF, Beijing, China. The distance from sample to detector was 3015 mm, and the wavelength of X-ray was 0.154 nm. The effective range of the scattering vector q (q = (4π sin θ)/λ, where 2θ is the scattering angle and λ is the wavelength) was 0.10−1.00 nm−1. Each pattern was collected within 60 s at room temperature. All the two-dimensional SAXS patterns were then background corrected and normalized using the standard procedure. The scattering patterns after calibration were averaged over all directions at a constant q, resulting in one-dimensional scattering intensity curves. Because of the isotropic distributed stacks of parallel lamellar crystallites in the system, a Lorentz correction (multiplication of I by q2) was performed in order to calculate the long spacing of the lamellar stacks.64 Beside 1D scattering intensity distribution, the correlation function analysis was also used to give detailed structural information on the system. The electron density correlation function K(z) can be derived from the inverse Fourier transformation of the intensity distribution I(q) as follows:63−65

Figure 1. DSC cooling and heating curves of PBcoE10. Cooling and heating rate are 20 and 5 K/min, respectively.

crystallization occurred. The supercooled melt was heated up to 130 °C at the rate of 5 K/min immediately after it reached 0 °C (lower red curve). An exothermic peak at 26 °C and an endothermic peak at 89 °C appeared, meaning that the supercooled melt crystallized first and melted again during heating process. In the second experiment, the sample with stable form I was heated up to 130 °C and held for 10 min followed by cooling and heating processes as mentioned above (blue curves), where a similar phenomenon was observed. There was no crystallization peak during cooling while the crystallization and melting peaks both appeared during heating. The melting temperatures of crystallites generated from the melt at different temperatures were nearly equal, smaller than the one of initial crystallites in form I. However, the crystallization peak at 17 °C from the 130 °C melt was much lower than the one from the 180 °C melt, indicating a higher crystallization rate for the low-temperature melt. The DSC results may provide a clue that the crystallization behavior was strongly influenced by melt temperature in this random butene1/ethylene copolymer. To investigate the effect of melt temperature on the crystallization kinetics further, isothermal crystallization experiments were carried out after the PBcoE10 in stable form I were melted at different temperatures. The crystallization time at conversion fraction of 0.5 from melt to crystalline phase was denoted as the half-time t1/2 of isothermal crystallization process. The top part in Figure 2 presents the effect of melt temperature Tm on the half-time t1/2 in PBcoE10 crystallized at 50 °C. The half-time t1/2 varies from 540 to 1700 s with increasing melt temperature from 130 to 180 °C. The bottom part shows the evolution of half-time t1/2 as a function of crystallization temperature Tc after melting at 130 and 180 °C, respectively. The crystallization rates decrease evidently with increasing crystallization temperature. At the same crystallization temperature, the half-time t1/2 in sample with the melt at 130 °C is always smaller than the one at 180 °C. This is the typical behavior of a faster crystallization kinetics associated with melt memory, even though the melt temperature of this random butene-1/ethylene copolymer is above its equilibrium melting temperature Tm,c ° . Indeed, Alamo et al.29 observed the similar behavior that the crystallization rate of ethylene copolymers increased as the initial melt temperature decreased even above Tm,c ° . They considered that the memory effect on crystallization of copolymers was due to the heterogeneous melt, and the melt above certain temperature was in a

∞

K (z) =

∫0 I(q)q2 cos(qz) dq ∞

∫0 I(q) dq

(1)

where z denotes the location measured along a trajectory normal to the lamellar surfaces, and the multiplication of I(q) with q2 (Lorentz correction) is performed. For systems with a structure of stacks of lamellae, the correlation function shows characteristic features that allow the long spacing defined as the average thickness of a lamella together with one interlamellar amorphous layer measured along the lamellar normal to be determined. To compare the melt temperature in the experiment, the equilibrium melting point of random butene-1/ethylene copolymer should be computed according to the Flory equation,66 which implies the exclusion of counits from the crystal lattice. 1 1 R − = ln XB ° Tm° Tm,c ΔHm°

(2)

where T°m and T°m,c are the equilibrium melting temperatures of ° homopolymer and copolymer, respectively, R is the gas constant, ΔHm is the melting enthalpy of ideal crystals, and XB is the total molar fraction of counits. However, the interchain composition distribution of Ziegler−Natta catalyzed copolymers is very broad, thus making an estimation of the equilibrium melting temperature with eq 2 unsound. Therefore, the equilibrium melting temperature T°m of form II and I of PB-1 homopolymer being 401 and 414 K42 are considered as upper limits for the copolymer under investigation. In this investigation, all the melt temperatures were equal to or higher than 130 °C, meaning that the random butene-1/ethylene copolymer was melted at temperatures over the equilibrium melting point of the homopolymer in form II and mostly also that of in form I.

■

RESULTS AND DISCUSSION Figure 1 shows the DSC results of butene-1/ethylene copolymer PBcoE10 with different thermal histories. First, the original PBcoE10 sample in stable form I was heated up to 180 °C and kept at this temperature for 10 min (black curve). The melting temperature of form I was 98 °C, and the small peak at 47 °C was associated with the melting of crystallites of low stability.67 Then the melt was cooled down from 180 to 0 °C at 20 K/min (upper red curve). Clearly, no exothermic peak can be observed during the cooling process, indicating that no C


Macromolecules

Article

Figure 3. Selected WAXD curves of PBcoE10 after crystallized at 50 °C with the melts at different temperatures.

In an effort to elucidate the effect of annealing time at molten state for the observed melt memory results, additional measurements were performed in such a way that the sample was held at 130 °C for 1 and 600 min and at 150 °C for 600 min before cooling down for crystallization at 50 °C. The resultant structure has been also measured by WAXD as is given in Figure 3. At 130 °C, the annealing time does not influence the final crystalline structure. One observes always diffraction data showing pure form I crystalline structure. At 150 °C, however, one observes a gradual increase in the content of form II crystallites after long time annealing. Clearly, the memory effect observed in this case is similar to the one observed in polyethylene copolymer that a heterogeneous melt structure can be removed only at very long time at high temperature.29 In general, form I refers to the stable crystal modification generated via a solid-state transformation of form II by aging at room temperature, and form I′ refers to the same hexagonal crystalline structure obtained by direct crystallization from melt or solution. Because of the identical crystallographic structures, form I and I′ crystals cannot be distinguished directly by WAXD. However, it is possible to identify the two forms using DSC due to the lower melting temperature of form I′. Figure 4 presents the DSC heating curves of crystals obtained after isothermal crystallization at 50 °C from the melts at different temperatures and the evolution of WAXD patterns of samples in form I′ and form II during heating. The small DSC melting peaks at about 58 °C are due to crystallites of lower stability.67 The melting temperature of form I′ crystals is 77 °C, lower than the melting temperature of form II at 81 °C. After melting at 150 and 160 °C, form I′ and form II coexist and the two peaks both appear from the corresponding crystals. The peak at 89 °C is due to the melting−recrystallization process of initial crystals. For comparison, the heating curve of stable form I transformed from form II is shown, whose melting temperature is 108 °C. It is obvious that the melting temperature of from I is much higher than the ones of form I′ and II, and form I′ shows even lower melting temperature than form II. Therefore, the crystalline form I obtained directly from the melt in the butene1/ethylene copolymer should be defined as form I′. WAXD measurements indicated that form I′ and form II in the heating process both melted without appearance of new crystalline phase, even though there exists melting and recrystallization process. Figure 5 presents the DSC heating curves of PBcoE10 in form I′ and form II after isothermal crystallization at different

Figure 2. Half-time of crystallization of PBcoE10 as a function of melt temperature Tm crystallized at 50 °C (top) and the half-time as a function of crystallization temperature Tc after melting at 130 and 180 °C (bottom).

homogeneous state which had no effect on the kinetics of crystallization. The results of crystallization kinetics in this butene-1/ ethylene copolymer imply that the melt state has a strong memory effect on the followed crystallization process. As a polymorphic polymer, it is necessary to investigate the effect of melt memory on the crystalline form. Thus, the thermal treatment shown in Scheme 1 was applied, while in situ WAXD measurements were carried out to follow the evolution of crystalline phase. PBcoE10 sample in stable form I was first heated up to a certain temperature ranging from 130 to 180 °C and held for 10 min. Then the melt was quickly cooled to 50 °C and kept for 60 min to complete the crystallization. Figure 3 presents the selected WAXD patterns of PBcoE10 after crystallized at 50 °C from the melt at different temperatures. The sample crystallized from the 180 °C melt exhibits three major peaks at 11.8°, 16.9°, and 18.4°, corresponding to the crystallographic planes of (200), (220), and (213) of form II crystals, respectively. However, when crystallized from the melt at 130 °C, three different diffraction peaks appear at 9.9°, 17.3°, and 20.5°, corresponding to the crystallographic planes of (110), (300), and (220 + 211) of form I crystals. For the samples crystallized from the melts of temperatures in between, diffraction peaks of both forms I and II can be observed, and the fraction of form I increases gradually with decreasing the melt temperature. As for all melt temperatures investigated only 130 °C is lower than the equilibrium melting point for homopolybutene-1, the observed occurrence of form I cannot be due to self-seeding induced by residue crystallites of form I upon heating. D


Macromolecules

Article

Figure 5. DSC heating curves of PBcoE10 samples obtained from 130 °C melt in form I′ (top) and from 180 °C melt in form II (bottom) after crystallized at different temperatures.

effect on the followed crystallization. Combined with WAXD results in Figure 3, the crystals in the clearly visible spherulites from the 180 °C melt were in form II, while the crystalline phase from the 130 °C melt was in form I′ without distinct spherulites in the current scale. The 150 °C melt resulted in the mixture of form I′ and II, where the diameter of spherulites was much smaller than the one from the 180 °C melt. The POM results state clearly that forms I′ and II own very different morphology. The WAXD, DSC, and POM results provide a direct evidence that form I′ is generated from the melt directly at ambient pressure. When the stable form I of butene-1/ethylene copolymer was melted at different temperature, the mixture of forms I′ and II can be generated from the melt. The relative content of two crystalline forms was simply controlled by the melt temperature. Clearly, the observed occurrence of form I′ in this PBcoE10 after its the initial form I crystals were melted completely at temperatures higher than the corresponding equilibrium melting point of homopolymer is different from the other results. In homopolymer, Li et al.68 found a partial recovery of form I crystals after they were melted partly at 128 °C. Cavallo et al.33 observed the generation of from I′ with low self-nucleation temperature where original form I crystals were also partly melted. They both considered that the survived form I crystals served as the nuclei of new crystals. In our case, as is shown in Figure 6, crystallization of form I′ in this PBcoE10 sample is dominated by extensive nucleation events in the heterogeneous melt, resulting in a rather diffuse morphology

Figure 4. DSC heating curves (top) of PBcoE10 after crystallized at 50 °C out of the melts being at different temperatures and the heating curve of stable form I obtained by phase transition from form II is shown for comparison. The evolution of WAXD patterns for samples obtained from 130 °C melt in form I′ (middle) and 180 °C melt in form II (bottom) during heating.

temperatures, and similar thermal behavior is detected. The melting peaks from crystallites of limited stability67 in all curves are about 8 °C higher than crystallization temperature Tc. When crystallization temperature is below 65 °C, the melting peaks from initial crystals and the ones obtained by melting and recrystallization both appear. However, when the sample was isothermally crystallized at temperatures higher than 65 °C, only the melting peak from initial crystals is observed, meaning that the crystals melt directly without recrystallization. Figure 6 shows the selected POM images of PBcoE10 after isothermal crystallization at 50 °C from the melt at 130, 150, and 180 °C. It is clear that the melt temperature had strong E


Macromolecules

Article

Figure 6. Selected POM images of PBcoE10 after isothermal crystallization at 50 °C from the melt at 130, 150, and 180 °C.

without clearly identifiable spherulitical structure at the current scale. In order to explore the source of different crystalline forms, another experiment was carried out following the procedure given in Scheme 2. PBcoE10 samples in stable form I was first heated up to 180 °C and held for 10 min to eliminate all possible traces of form I crystallites. The melt was quickly cooled to 50 °C and kept for different times from 0 to 60 min to produce form II crystallites to different extent. For clarity, this first crystallization time at 50 °C was defined as t1,50cry. Then the samples with different t1,50cry (differing in content of form II crystals) were heated up to 130 °C and held for 10 min. The 130 °C melt was again cooled down to 50 °C and kept for 60 min to complete the second crystallization. In the top of Figure 7, selected WAXD curves in PBcoE10 first crystallized at 50 °C for different t1,50cry from the 180 °C melt are present. When t1,50cry was 0 min, there was no diffraction peak of crystalline phase, meaning that the system was amorphous. Three diffraction peaks of form II showed up at t1,50cry = 20 min. The diffraction intensities of pure form II increased with the increase of t1,50cry. The bottom part in Figure 7 presents the selected WAXD curves in PBcoE10 completely crystallized at 50 °C after the samples with different fractions of form II crystals were melted at 130 °C. When the first crystallization time t1,50cry was 0 min, the corresponding WAXD curve in the second crystallization process shows only the diffraction peaks of form II. With increasing t1,50cry, the diffraction peaks of form I′ appeared and enhanced while the diffraction intensity of peaks of form II decreased gradually. At 60 min of t1,50cry, only the diffraction peaks from form I′ present. Clearly, form I′ crystals were generated from the melt even though no form I crystals exist in the previous system. The same thermal process was followed by POM as present in Figures 8 and 9. In Figure 8a, the system was crystallized partly at t1,50cry = 12 min, showing randomly distributed spherulites with crystals in form II were generated from the 180 °C melt. The partially crystallized sample was then heated up to 130 °C and kept for 10 min before being quickly cooled down to 50 °C where image shown in Figure 8b was taken. Clearly, the film showed an amorphous morphology without any visible crystalline structure in accord with DSC results in Figure 1. After the sample was kept at 50 °C for 5 min, as is shown in Figure 8c, crystallization of the sample into form I′ can be observed occurring exactly at positions previous occupied by spherulites with form II crystals. Notably, the newly formed crystals possess very different morphology compared to the previous form II spherulites. No clear spherulitical morphology can be observed for these newly formed form I′ crystallites at the current scale. Further prolonging the crystallization time at

Figure 7. Selected WAXD curves in PBcoE10 during isothermal crystallization at 50 °C cooled down from the melt at 180 °C (top) and WAXD curves after completion of isothermal crystallization at 50 °C for samples being melted at 130 °C after partially crystallization in form II shown in the top plot (bottom).

50 °C led to a development of form II spherulites at positions where no previous crystalline structures existed before the sample was melted at 130 °C. Such mixed morphology of form I′ and form II crystals can be seen in Figure 8d. Clearly, the 180 °C melt still kept the random state even when it was cooled to low temperature, which led to the formation of form II crystals during isothermal crystallization at 50 °C. The appearance of much smaller spherulites in form I′ within precious bigger spherulites of form II indicates the occurrence of massive nucleation of form I′ crystallites in the heterogeneous melt which was faster than the nucleation and growth of form II crystallites in the surrounding homogeneous melt, which was also consistent with DSC results. POM pictures in Figure 9 reinforced this view that with increasing t1,50cry from 12 to 60 min the crystalline phase in form I′ enhanced obviously, and F


Macromolecules

Article

Figure 8. Selected POM images of PBcoE10 at different stage of Scheme 2. (a) Form II crystallized at 50 °C for 12 min from the melt at 180 °C. (b−d) Form II crystallites melted at 130 °C and crystallized at 50 °C for 0, 5, and 60 min.

Figure 9. Selected POM images of PBcoE10 after completion of isothermal crystallization at 50 °C for samples having been melted at 130 °C after crystallization in form II for 12, 20, and 60 min.

the number of imperfect spherulites in form II reduced gradually, which confirmed the WAXD results in Figure 7. The WAXD and POM results both verified that the only prerequisite for forming form I′ crystals is that the sample should be in a crystalline state regardless of its crystalline form prior to melting treatment below certain temperature. Phenomenologically, the whole processes can be summarized as follows. First, the form I crystals changed into a homogeneous melt upon heating to 180 °C. In the first crystallization, form II crystals were generated from the homogeneous melt, whose content increased with the crystallization time. Here, the melt, which did not crystallize into form II, kept the homogeneous state even at 50 °C. Then, the mixture of form II crystals and remaining homogeneous melt were heated up to 130 °C, producing a mixture of heterogeneous and homogeneous melts. At last, the mixture of form I′ and form II crystals were generated in the second crystallization process. It was obvious that form I′ was from the heterogeneous melt while form II was from the homogeneous melt. The results provided an evidence for the memory effect, which is due to the heterogeneous melt. Clearly, the heterogeneous melt must lose crystallographic packing in close proximity due to a relatively slow thermal mobility or be

some type of segmental ordered structures independently of previous packing of crystalline sequences.29 No matter whether it was from form I or form II, the heterogeneous melt lost the helical conformation of previous crystals but kept some sequence-length segregation and induced the formation of nuclei of form I′. The optical observations suggest that the growth of form I′ cannot cross the boundary between heterogeneous and homogeneous melts. Such behavior can be better evidenced in in situ WAXD experiments. As is shown in Figure 10, WAXD data were collected during isothermal crystallization of a PBcoE10 sample at 50 °C after being melted at 130 °C from a partially crystallized state in form II. Clearly, one observes first the only appearance of diffraction peaks from form I′ at short crystallization time followed by gradual development of form II diffraction peaks. In the bottom of Figure 10, time-dependent fractions of form I′, form II, and amorphous phase were given. Clearly, the developments of form I′ and form II crystallinity are two independent events with different time constants. For form I′, crystallization practically stopped at around 15 min whereas for form II crystallization continued until around 30 min. Therefore, both POM and WAXD data indicated that the G


Macromolecules

Article

Figure 10. Evolution of WAXD curves (top) and fraction of form I′ and II and amorphous phase (bottom) during isothermal crystallization at 50 °C for samples being melted at 130 °C after partial crystallization in form II for 30 min.

occurrence of form I′ crystallites was restricted within the heterogeneous melt only. To further explore the mechanism of the direct formation of form I′ in the system, structural parameters at the lamellar level are necessary. Thus, SAXS measurements were conducted, and the long spacing, lamellar thickness, and amorphous layer thickness were obtained for samples of different crystalline modification produced isothermally at different crystallization temperatures. Because of the similar electron density difference between crystalline and amorphous phase in form II, the SAXS intensity is very weak. Therefore, structural parameters of the samples with form II were determined by measuring their corresponding phase transformed form I samples with higher crystalline density, which provided a much-enhanced contrast in SAXS measurements. The SAXS results are shown in Figure 11. In the Lorentz-corrected one-dimensional scattering intensity distribution profiles, the scattering peaks from form I′ and I both moved to smaller q with increasing crystallization temperature, indicating that the long spacing increased gradually. In addition, the scattering peaks of samples with form I′ locate at much larger q values than the corresponding samples with form I, meaning that the long spacing of form I′ was much smaller. By means of the correlation function approach shown in the inset of the top part in Figure 11, the long spacing dac, lamellar thickness dc, and linear crystallinity Xl were calculated. The long spacing of form I crystalline lamellar stacks ranged from 18.6 to 27.2 nm, much higher than the one of form I′ crystalline lamellar stacks being from 11.0 to 16.5 nm in the current range of crystallization temperature. The linear crystallinity of form I is a little higher than the one of form I′, and they both keep constant with increasing the crystallization temperature, which is similar to the behavior in our previous

Figure 11. 1D SAXS intensity distribution profiles with Lorentz correction of PBcoE10 in form I obtained after phase transition from isothermal crystallized form II crystallites (top), directly isothermal crystallized form I′ (middle), and the lamellar thickness dc of form I′ and I as a function of crystallization temperature Tc (bottom). The inset in top part shows how the average thicknesses of the long spacing dac and lamellar layers dc are derived from the correlation function curve. The linear crystallinity Xl is calculated by the equation Xl = dc/ dac.

studies.69,70 Phase transformation from form II to I involves an extension of the 113 helical conformation (form II) into the 31 helix (form I). Therefore, it is necessary to acquire the thickness in form II by the relationship dc(II) = dc(I)/1.12, which is due to the ratio between the axial repeating units of two conformations.69 The lamellar layer thickness in form II was therefore from 5.6 to 8.4 nm, which is much thicker than the one in form I′ from 3.7 to 5.0 nm after isothermal crystallization at different temperatures. The result is consistent with the discovery that the minimum thickness required for the lamellar crystals to grow was lower for form I′ than for form II and increased with crystallization temperature.54 H


Macromolecules

■

Article

CONCLUSION In the present study, the effect of melt memory on the crystallization of butene-1/ethylene copolymer was investigated by means of DSC, WAXD, POM, and SAXS techniques. The kinetics of crystallization depended on the melt temperatures strongly, which were even above the equilibrium melting point. The crystallization rate increased with decreasing melt temperature. The evolution of crystalline phase with different thermal histories was explored. When the crystals were melted at high temperature, only form II crystallites were generated, while form I′ crystals were generated from the melt at low temperature. The formation of form I′ and II was influenced by the melt state, independent of the previous crystalline phase. The results can be understood that at high melt temperature the system formed a random homogeneous melt, where the nucleation of form I′ crystallites was suppressed so that form II crystallites were developed. However, when sample with previous crystalline structures was melted at low temperature yet above the equilibrium melting temperature, a heterogeneous melt state was formed which significantly favored the nucleation of form I′ crystallites. The fraction of forms I′ and II can be controlled by changing the content of heterogeneous and homogeneous melts in the system. In addition, the lamellae in form I′ was thinner than the ones in form II. Therefore, the melt memory was due to the preordered sequence-length segregation in the heterogeneous melt, which was from the previous crystals but lost their crystallographic conformation. This result provides a direct route to acquire the stable form I′ of PB-1 copolymers from the melt at atmospheric pressure without passing through metastable form II in industrial process.

The results obtained in this work enabled us to propose a scheme to explain the crystallization habits in PB-1 and its random copolymers. The major findings include (i) the butene1/ethylene random copolymer of 9.88 mol % ethylene counits can crystallize into either form II or form I′ depending only on the melt temperature regardless of previous crystalline form before melting, (ii) the direct crystallization in form I′ is dominated by the massive nucleation process occurring in the heterogeneous melt, and (iii) the lamellar thickness of both form I′ and form II crystallites depends on crystallization temperature regardless of original crystalline structure before melting. First of all, the above findings suggest that heterogeneities in sequential length segregation were preserved during melting at temperatures above the equilibrium melting point, which promoted the nucleation of form I′ crystals after cooling down. The thickness of newly formed form I′ lamellae had no memory on previous crystalline structure but only depended on crystallization temperature. Clearly, the localized segmental near parallel packing in the heterogeneous melt, although without any preferred helical conformation, effectively initiated the massive nucleation and growth of form I′ crystallites before the development of nucleation and growth of normally favored form II crystallites. Although, at this moment, it is still not certain about the exact driving force for the formation of form I′ nuclei in such heterogeneous melt, it is most likely that the localized segmental segregation significantly lowered down the nucleation barrier for form I′ nuclei. In PB-1 homopolymers, the crystallization of form II is always easier than that of form I′ as was shown in various experiments with partially melted form I samples that form II crystallization occurred extensively despite the existence of a large number of form I crystallites working as nuclei for the growth of form I′.68 It is also very important to notice that the growth of form II crystallites is significantly slowed down by the inclusion of counits on the polymer chain.71 Such influence of counits on the form II crystallization rate can be understood as a consequence of much thicker crystalline lamellae in the system in form II, requiring therefore much longer crystallizable chain segments diffusing onto the growing front where an additional segmental length selection is necessary. In the extreme case, when the counits content reaches such an amount that statistically the crystallizable chain segments are not long enough to build up form II crystallites, only form I′ crystallites can be realized as was reported in the literature. When both crystalline modifications are possible, as the thickness of form I′ lamellar crystallites is much smaller than the corresponding form II crystallites obtained at same isothermal conditions, it is possible that at certain counits content the nucleation and growth rate of form I′ becomes similar to or even higher than that of form II. Indeed, such a case was found in this work. Following this viewpoint, we could also offer a possible explanation for the direct formation of form I′ for PB-1 homopolymer in thin film as reported by Yan and co-workers.47 Although no exact film thickness was given in ref 47, the film could be considered very thin as it was obtained by spin-coating of a dilute solution of PB-1. In such a thin film at the given high isothermal crystallization temperature, nucleation of form II crystallites would be strongly suppressed due to the fact that the film thickness might be much thinner than that would be required for the formation of form II crystals. Clearly, when the formation of form II is limited, a much thinner but normally kinetically slower form I′ would develop.

■

AUTHOR INFORMATION

Corresponding Author

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

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS We thank Prof. Zhonghua Wu and Dr. Guang Mo of BSRF for assistance during SAXS experiments and Ms Ran Chen for providing computer code to calculate crystallinity from WAXD data. This work is supported by the National Natural Science Foundation of China (21134006).

■

REFERENCES

(1) Strobl, G. The Physics of Polymers, 2nd ed.; Springer: Berlin, Germany, 1997. (2) Wunderlich, B. Macromolecular Physics; Academic Press: New York, 1976; Vol. 2. (3) Alfonso, G. C.; Ziabicki, A. Colloid Polym. Sci. 1995, 273, 317− 323. (4) Supaphol, P.; Lin, J.-S. Polymer 2001, 42, 9617−9626. (5) Mamun, A.; Umemoto, S.; Okui, N.; Ishihara, N. Macromolecules 2007, 40, 6296−6303. (6) Xu, J. J.; Ma, Y.; Hu, W. B.; Rehahn, M.; Reiter, G. Nat. Mater. 2009, 8, 348−353. (7) Martins, J. A.; Zhang, W.; Brito, A. M. Polymer 2010, 51, 4185− 4194. (8) Lorenzo, A. T.; Arnal, M. L.; Sánchez, J. J.; Müller, A. J. J. Polym. Sci., Part B: Polym. Phys. 2006, 44, 1738−1750. (9) Zhang, Y.-S.; Zhong, L.-W.; Yang, S.; Liang, D.-H.; Chen, E.-Q. Polymer 2012, 53, 3621−3628. I


Macromolecules

Article

(10) Fillon, B.; Wittmann, J. C.; Lotz, B.; Thierry, A. J. Polym. Sci., Part B: Polym. Phys. Ed. 1993, 31, 1383−1393. (11) Androsch, R.; Wunderlich, B. Macromolecules 2000, 33, 9076− 9089. (12) Zhu, X.; Li, Y.; Yan, D.; Zhu, P.; Lu, Q. Colloid Polym. Sci. 2001, 279, 292−296. (13) Cho, K.; Saheb, D. N.; Choi, J.; Yang, H. Polymer 2002, 43, 1407−1416. (14) Cho, K.; Saheb, D. N.; Yang, H. C.; Kang, B. I.; Kim, J.; Lee, S. S. Polymer 2003, 44, 4053−4059. (15) Häfele, A.; Heck, B.; Hippler, T.; Kawai, T.; Kohn, P.; Strobl, G. Eur. Phys. J. E 2005, 16, 217−224. (16) Maus, A.; Hempel, E.; Thurn-Albrecht, T.; Saalwächter, K. Eur. Phys. J. E 2007, 23, 91−101. (17) Khanna, Y. P.; Kumar, R.; Reimschuessel, A. C. Polym. Eng. Sci. 1988, 28, 1607−1611. (18) Ziabicki, A.; Alfonso, G. C. Colloid Polym. Sci. 1994, 272, 1027− 1042. (19) Reinshagen, J. H.; Dunlap, R. W. J. Appl. Polym. Sci. 1973, 17, 3619−3625. (20) Supaphol, P.; Spruiell, J. E. J. Appl. Polym. Sci. 2000, 75, 337− 346. (21) Rastogi, S.; Lippits, D. R.; Peters, G. W. M.; Graf, R.; Yao, Y.; Spiess, H. W. Nat. Mater. 2005, 4, 635−641. (22) Lippits, D. R.; Rastogi, S.; Höhne, G. W. H. Phys. Rev. Lett. 2006, 96, 218303/1−4. (23) Lippits, D. R.; Rastogi, S.; Höhne, G. W. H.; Mezari, B.; Magusin, P. C. M. M. Macromolecules 2007, 40, 1004−1010. (24) Talebi, S.; Duchateau, R.; Rastogi, S.; Kaschta, J.; Peters, G. W. M.; Lemstra, P. J. Macromolecules 2010, 43, 2780−2788. (25) Pandey, A.; Toda, A.; Rastogi, S. Macromolecules 2011, 44, 8042−8055. (26) Hikosaka, M.; Watanabe, K.; Okada, K.; Yamazaki, S. Adv. Polym. Sci. 2005, 191, 137−186. (27) Yamazaki, S.; Gu, F.; Watanabe, K.; Okada, K.; Toda, A.; Hikosaka, M. Polymer 2006, 47, 6422−6428. (28) Luo, C. F.; Sommer, J.-U. Phys. Rev. Lett. 2014, 112, 195702/1− 5. (29) Reid, B. O.; Vadlamudi, M.; Mamun, A.; Janani, H.; Gao, H. H.; Hu, W. B.; Alamo, R. G. Macromolecules 2013, 46, 6485−6497. (30) Gao, H. H.; Vadlamudi, M.; Alamo, R. G.; Hu, W. B. Macromolecules 2013, 46, 6498−6506. (31) De Rosa, C.; Ruiz de Ballesteros, O.; Di Gennaro, M.; Auriemma, F. Polymer 2003, 44, 1861−1870. (32) Sorrentino, A.; Pantani, R.; Titomanlio, G. J. Polym. Sci., Part B: Polym. Phys. 2007, 45, 196−207. (33) Cavallo, D.; Gardella, L.; Portale, G.; Müller, A. J.; Alfonso, G. C. Polymer 2014, 55, 137−142. (34) Meille, S. V. Polymer 1994, 35, 2607−2612. (35) Wu, M. C.; Woo, E. M. Polym. Int. 2005, 54, 1681−1688. (36) Pan, H.; Na, B.; Lv, R.; Li, C.; Zhu, J.; Yu, Z. J. Polym. Sci., Part B: Polym. Phys. 2012, 50, 1433−1437. (37) Natta, G.; Corradini, P.; Bassi, I. W. Nuovo Cimento Suppl. 1960, 15, 52−67. (38) Miller, R. L.; Holland, V. F. J. Polym. Sci., Part B: Polym. Lett. 1964, 2, 519−521. (39) Turner-Jones, A. T. J. Polym. Sci., Part B: Polym. Lett. 1963, 1, 455−456. (40) Kopp, S.; Wittmann, J. C.; Lotz, B. Polymer 1994, 35, 908−915. (41) Boor, J.; Mitchell, J. C. J. Polym. Sci., Part A: Gen. Pap. 1963, 1, 59−84. (42) Powers, J.; Hoffman, J. D.; Weeks, J. J.; Quinn, F. A., Jr. J. Res. Natl. Bur. Stand. 1965, 69A, 335−345. (43) Azzurri, F.; Flores, A.; Alfonso, G. C.; Sics, I.; Hsiao, B. S.; Baltá Calleja, F. J. Polymer 2003, 44, 1641−1645. (44) Azzurri, F.; Flores, A.; Alfonso, G. C.; Baltá Calleja, F. J. Macromolecules 2002, 35, 9069−9073. (45) Liu, Y. P.; Cui, K. P.; Tian, N.; Zhou, W. Q.; Meng, L. P.; Li, L. B.; Ma, Z.; Wang, X. L. Macromolecules 2012, 45, 2764−2772.

(46) Wang, Y. T.; Jiang, Z. Y.; Wu, Z. H.; Men, Y. F. Macromolecules 2013, 46, 518−522. (47) Zhang, B.; Yang, D. C.; Yan, S. K. J. Polym. Sci., Part B: Polym. Phys. 2002, 40, 2641−2645. (48) Yamashita, M.; Ueno, S. Cryst. Res. Technol. 2007, 42, 1222− 1227. (49) Yamashita, M.; Hoshino, A.; Kato, M. J. Polym. Sci., Part B: Polym. Phys. 2007, 45, 684−697. (50) Armeniades, C. D.; Baer, E. J. Macromol. Sci., Part B: Phys. 1967, 1, 309−334. (51) Nakafuku, C.; Miyaki, T. Polymer 1983, 24, 141−148. (52) Kalay, G.; Kalay, C. R. J. Appl. Polym. Sci. 2003, 88, 814−824. (53) Li, L.; Liu, T.; Zhao, L. Polymer 2011, 52, 5659−5668. (54) De Rosa, C.; Auriemma, F.; Ruiz de Ballesteros, O.; Esposito, F.; Laguzza, D.; Di Girolamo, R.; Resconi, L. Macromolecules 2009, 42, 8286−8297. (55) Androsch, R.; Hohlfeld, R.; Frank, W.; Nase, M.; Cavallo, D. Polymer 2013, 54, 2528−2534. (56) Marega, C.; Spataro, S.; Fassone, E.; Camurati, I.; Marigo, A. J. Appl. Polym. Sci. 2014, 131, 40119/1−12. (57) Stolte, I.; Cavallo, D.; Alfonso, G. C.; Portale, G.; Drongelen, M. V.; Androsch, R. Eur. Polym. J. 2014, 60, 22−32. (58) De Rosa, C.; Ruiz de Ballesteros, O.; Auriemma, F.; Di Girolamo, R.; Scarica, C.; Giusto, G.; Esposito, S.; Guidotti, S.; Camurati, I. Macromolecules 2014, 47, 4317−4329. (59) Miyoshi, T.; Hayashi, S.; Imashiro, F.; Kaito, A. Macromolecules 2002, 35, 2624−2632. (60) Lotz, B.; Thierry, A. Macromolecules 2003, 36, 286−290. (61) Jiang, Z. Y.; Sun, Y. Y.; Tang, Y. J.; Lai, Y. Q.; Funari, S. S.; Gehrke, R.; Men, Y. F. J. Phys. Chem. B 2010, 114, 6001−6005. (62) Nakamura, K.; Aoike, T.; Usaka, K.; Kanamoto, T. Macromolecules 1999, 32, 4975−4982. (63) Glatter, O.; Kratky, O. Small-Angle X-ray Scattering; Academic Press: London, 1982. (64) Strobl, G.; Schneider, M. J. J. Polym. Sci., Polym. Phys. 1980, 18, 1343−1359. (65) Strobl, G.; Schneider, M. J.; Voigt-Martin, I. G. J. Polym. Sci., Polym. Phys. 1980, 18, 1361−1381. (66) Flory, P. J. Principles of Polymer Chemistry; Cornell University Press: Ithaca, NY, 1953; p 570. (67) Al-Hussein, M.; Strobl, G. Macromolecules 2002, 35, 1672− 1676. (68) Su, F. M.; Li, X. Y.; Zhou, W. M.; Zhu, S. S.; Ji, Y. X.; Wang, Z.; Qi, Z. M.; Li, L. B. Macromolecules 2013, 46, 7399−7405. (69) Wang, Y. T.; Jiang, Z. Y.; Fu, L. L.; Lu, Y.; Men, Y. F. Macromolecules 2013, 46, 7874−7879. (70) Wang, Y. T.; Lu, Y.; Jiang, Z. Y.; Men, Y. F. Macromolecules 2014, 47, 6401−6407. (71) Stolte, I.; Androsch, R. Polymer 2013, 54, 7033−7040.

J


Ethylene

Recommend Documents