Molecular Weight Dependency of Crystallization Line

Sep 12, 2014 - Isothermal crystallization and subsequent melting processes in semicrystalline polymers can be described by several linear dependencies...
0 downloads 0 Views 406KB Size
Download PDF
Article pubs.acs.org/Macromolecules

Molecular Weight Dependency of Crystallization Line, Recrystallization Line, and Melting Line of Polybutene‑1 Yaotao Wang, Ying Lu, 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: Isothermal crystallization and subsequent melting processes in semicrystalline polymers can be described by several linear dependencies termed as crystallization line, recrystallization line, and melting line in a phase diagram of inversed crystalline lamellar thickness as a function of temperature. The crystallization line denotes the lamellar thickness dependency on crystallization temperature. Upon subsequent heating, the system either undergoes a process involving a continuous change in lamellar thickness following a recrystallization line or a direct melting at a constant thickness defining a melting line. The limiting temperatures of infinite lamellar thickness for crystallization line and recrystallization line coincide with each other and are higher than that of melting line. The above-mentioned experimental results are interpreted as a consequence of involvement of a mesomorphic phase during crystallization. In this work, the effect of molecular weight on the crystallization line, recrystallization line, and melting line was investigated for three polybutene-1 (PB-1) samples in metastable form II by means of temperature-dependent small-angle X-ray scattering experiments. Three crystallization lines with different slopes but same limiting temperature were observed for the three PB-1 samples. However, a common recrystallization line and melting line were observed for all three samples. The results can be understood as follows. The surface free energy of the native crystals formed from the mesomorphic blocks increases with molecular weight in PB-1, leading to an increase of the slop of the crystallization line for samples with higher molecular weight. However, during the stabilization process after the formation of native crystals forming stable crystals the surface free energy decreases to such an extent that the differences among samples of different molecular weight vanish so that a common recrystallization line and melting line were observed.

■

INTRODUCTION Semicrystalline polymers are composed of stacked crystalline lamellae and entangled amorphous polymeric chains in between.1 The crystallization in semicrystalline polymers which transfers the entangled melt into a semicrystalline state is a process of primary importance that crucially affects the final properties of polymeric materials and has been studied for a long time. In recent years, by means of systematic temperaturedependent small-angle X-ray scattering (SAXS) measurements, Strobl et al. studied the melting behavior, subsequent to an isothermal crystallization, of several polymers, including polyethylene (PE),2−5 isotactic polypropylene (iPP),6 syndiotactic polypropylene (sPP),2,7−11 polybutene-1 (PB-1),12 poly(ε-caprolactone) (PCL),2 polystyrene (PS),9,13 and poly(Llactide) (PLLA).14 A phase diagram between inversed lamellar thickness (dc−1) and temperature was constructed, in which there exist several linear relationships between the inversed lamellar thickness and the corresponding temperatures such as a crystallization line denoting the linear relationship between dc−1 and crystallization temperature (Tc), a recrystallization line showing a continues thickening process during heating before final melting and a melting line depicting the linear dependency of dc−1 on the melting temperature (Tm). Notably, the limiting © 2014 American Chemical Society

temperatures of infinite lamellar thickness for crystallization line and recrystallization line coincide with each other and are higher than that of the melting line. The above-mentioned experimental results are interpreted as a consequence of involvement of a mesomorphic phase during crystallization. Therefore, four different phases are included in the diagram, such as the amorphous melt (a), mesomorphic layers (m), native crystals (cn), and stabilized crystals (cs).The crystallization line describes therefore a transition from mesomorphic phase to native crystalline phase. The recrystallization line represents a transition between stabilized crystals and mesomorphic phase. The melting line stands for a transition from stabilized crystals to amorphous melt state directly. The limiting temperature for infinite thick lamellae on crystallization line and recrystallization is therefore denoted as T∞ mc, and the 15,16 17−22 one for melting line is T∞ . Stribeck et al. investigated ac the oriented crystallization of PE and PP from the quiescent melt by SAXS measurements and utilized the multidimensional chord distribution function (CDF) to acquire the structural Received: June 20, 2014 Revised: August 18, 2014 Published: September 12, 2014 6401

dx.doi.org/10.1021/ma501272a | Macromolecules 2014, 47, 6401−6407

Macromolecules

Article

parameters. Their studies corroborated independently Strobl’s models on the pathways of crystallization via secondary crystallites creating a block mesostructure. The crystallization of polymer depends on different parameters, and molecular weight (MW) is one of the most important factors governing the properties of polymer. The MW dependency of crystallization kinetics and the morphology has been explored in many polymers, such as PE,23−28 iPP,28−30 PB-1,12,31−34 which is necessary to understand the crystallization mechanism of polymers. However, systematic quantitative studies of the effect of MW on crystallization of polymer are not common, and the dependence of the overall crystallization rate and lamellar thickness on MW is not yet clear. Mandelkern24 and John Manley25 both found that the surface free energy of a mature crystallite in PE increased significantly with MW. Fatou et al.27 studied the overall crystallization rate of PE over a wide MW range and found a peculiar behavior: as MW increased, the half crystallization time decreased, then passed through a minimum, and increased again to a MW-independent plateau. Cheng29 observed that at certain crystallization temperature the melting temperature of iPP increased with MW, which meant that the lamellae in iPP with larger MW were thicker, and Yamada31 also found the crystalline lamellar thickness of iPP increased with MW. Fu et al.12 found a transition from folded-chain to chain-extended crystallization with increasing the crystallization temperature in PB-1. Acierno et al.34 investigated the MW dependence of critical temperature at which two crystalline morphologies coexisted. It is well-known that polymers can possess different crystalline forms. PB-1, for example, may exist at least in three crystal modifications: the hexagonal form I, tetragonal form II, and orthorhombic form III. Lamellar crystallites of metastable form II are generally formed when PB-1 crystallizes from the melt, which transforms into stable form I crystals when keeping the polymer at room temperature for several days.36−40 With the help of tensile deformation, the transformation process will be accelerated,41−43 while the reversible process can occur when deformed at high temperature.44 And PB-1 with low tacticity can be crystallized into form I directly.45−47 In addition, the effect of cross-nucleation and self-nucleation on the crystallization behavior of PB-1was investigated recently.48−51 Form III has been observed in the case of crystallization from solution.52 In this study, we explored the effect of MW on the crystallization and melting behaviors of three PB-1 samples in form II. DSC was used to study the kinetics of isothermal crystallization from the melt. By means of temperaturedependent SAXS experiments, we observed three crystallization lines, a common recrystallization line, and melting line in PB-1 samples with different MW. The results can be understood by considering changes in thermal dynamic parameters during the formation of native crystals and stabilized crystals.

■

Table 1. Characterization of Samples trade name

melt flow rate (MFR) (190 °C/2.16)

Mw (kg/mol)

Tm (form II)a (°C)

crystallinity, Φwa (%)

PB0110M PB0400M PB0800M

0.4 16.4 200

711 188 77

115.9 115.4 115.1

61 64 71

a The melting temperature (Tm) in form II and the corresponding crystallinity (Φw) were derived from the second DSC heating scan at a heating rate of 10 K/min after a cooling from the molten state at a rate of 10 K/min.

DSC measurements were conducted in two ways by means of a DSC1 Stare system (Mettler Toledo Instruments, Swiss), which had been calibrated for temperature and melting enthalpy using indium as a standard. The first way was to investigate the crystallization kinetics of PB-1 by cooling samples quickly from melt to a desired crystallization temperature followed by an observation of crystallization process isothermally. The second one was to measure the heat flow of PB-1 in form II at a heating rate of 10 K/min from 0 to 180 °C, from which the melting temperature Tm and weight fraction crystallinity Φw were obtained. For the calculation of crystallinity, a value of ΔHid = 62 J/g were chosen for the melting enthalpy of ideal crystals in form II.54 In-situ synchrotron SAXS measurements were performed at beamline 1W2A, BSRF, Beijing, China. A piece of PB-1 sample in form II directly after isothermal crystallization, tightly wrapped with a thin aluminum foil in order to promote thermal conductivity, was heated from crystallization temperature (Tc) to melting by a portable heating device (TST350, Linkam, UK) installed at the beamline to follow the microstructure evolution during heating. The distance from sample to detector was 4959 mm, the wavelength of X-ray was 0.154 nm, and then 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 0.065−0.600 nm−1. In-situ SAXS measurements were carried out during the heating scan from Tc to 140 °C at a rate of 1 K/min, and each pattern was collected within 60 s. All the two-dimensional (2D) 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 (1D) 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.55 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 experimentally intensity distribution I(q) as follows:55−57 ∞

K (z) =

∫0 I(q)q2 cos(qz) dq ∞

∫0 I(q)q2 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) was 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.

EXPERIMENTAL SECTION

■

Three commercial grade isotactic PB-1 samples with different MW produced by LyondellBasell were used in this study. A summary of the samples’ physical properties is listed in Table 1.53 PB-1 samples were first compression-molded into films of about 0.5 mm in thickness at 180 °C and held in the molten state for 5 min to erase the processing history. The molten films were then transferred rapidly into isothermal water bath at different preset temperatures (Tc = 30, 40, 50, 60, 70, 80, and 90 °C) and held isothermally for 5 h to complete the crystallization in the samples.

RESULTS AND DISCUSSION It is well-known that crystallization of polymers is kinetically controlled. The isothermal crystallization process strongly depends on the crystallization condition. The lamellar thickness increases with the crystallization temperature, accompanied by an increase of crystallization time. The top part in Figure 1 shows the conversion−crystallization time curves of PB0110M 6402

dx.doi.org/10.1021/ma501272a | Macromolecules 2014, 47, 6401−6407

Macromolecules

Article

ment of the chains, resulting in “chain-folded” crystallites. However, if the crystal thickness surpassed the diameter, the polymer chains should begin to disentangle, resulting in the formation of “chain-extended” crystallites. It is reasonable that the disentanglement process may reduce the rate of crystallization and explain the interesting behavior of the ratio of crystallization half-times in PB0110M and PB0800M. To check the change of crystallization mechanism, one should compare the radius of gyration Rg with the lamellar thickness dc. The radius of gyration Rg is an estimate for the diameter of the volume occupied by the chains in the melt, and Rg and the mean-squared end-to-end distance R0 are for Gaussian chains related by eq 2

R 02 (2) 6 and R0 can be calculated using the characteristic ratio C∞ as Rg2 =

R 0 2 = C∞ab 2N

(3)

ab2

where stands for the sum of the squares of the lengths of the backbone bonds of one monomer unit and N is the degree of polymerization. For PB-1, C∞ is 18 and ab2 is 4.74 × 10−2 nm2.58 The values calculated for Rg and dc at different crystallization temperatures are shown in Table 2. At 50 °C, Table 2. Values of the Radius of Gyration Rg and Lamellar Thickness dc of Three PB-1 Samples

Figure 1. Top: selected conversion−crystallization time curves of PB0110M from melt to form II crystallites. Bottom: the half-time of form II formation in three PB-1 samples with different molecular weight Mw and the ratio between them as a function of crystallization temperature Tc.

trade name

Mw (kg/mol)

C∞

PB0800M PB0400M PB0110M

77 188 711

18 18 18

ab2

2

(nm )

4.74 × 10−2 4.74 × 10−2 4.74 × 10−2

Rg (nm)

dc at 50 °C (nm)

dc at 90 °C (nm)

14.0 21.9 42.4

12.2 13.2 14.5

20.8 22.1 24.5

Rg of PB0110M and PB0400M are much larger than dc, indicating that the “chain-folded” crystals were formed, and the crystallization rate increased with MW. However, Rg of PB0800M is much smaller than dc at 90 °C, resulting in the formation of “chain-extended” crystals. Based on MW dependency of transition temperature in PB-1 measured by Acierno,34 the transition temperature of PB0800M and PB0400M with MW below 200 kg/mol was about 90 °C, and the one of PB0110M with MW of 711 kg/mol was so high that only the “chain-folded” crystallites could be formed. Figure 2 depicts the DSC results in three PB-1 samples. In the top plot, the DSC thermograms measured during heating immediately after the completion of isothermal crystallization at the indicated temperature are present. One observes that the endothermic peak for the sample with Tc below 70 °C keeps almost constant at about 115 °C, and then the peak moves to higher temperatures for the one with Tc above 70 °C. The results indicate the presence of melting and recrystallization processes. This means that the lamellae below a critical thickness experienced a melting and recrystallization process during heating, while the lamellae above the critical value melted directly.11 The middle plot collects the crystallinities (Φw) measured after completion of the crystallization process and the lamellar thickness (dc). It is noticed that despite the large change of lamellar thickness, the change of crystallinity remains small, which is also observed in other systems.2,6 A change of the crystallization temperature modifies the length scales of the partially crystalline structure but has no effect on

from melt to form II crystallites derived from DSC measurements. The values present in this conversion plot were normalized to the final crystallinity at each temperature. The crystallization rate of form II decreases with increasing crystallization temperature Tc. The crystallization time at conversion fraction of 0.5 is denoted as the half-time t1/2 of isothermal crystallization process. The bottom part presents the evolution of the half-time of crystallization of three PB-1 samples and the ratio between them as a function of crystallization temperature Tc. The half-time t1/2 varies from tens to thousands of seconds when Tc increases from 80 to 100 °C. The half-time t1/2 increased with MW at low Tc, while at higher Tc the half-time of PB0110M with highest MW became smaller than the other two samples. For clarity, the ratio of the half-times of crystallization in samples with different MW is also present. The ratio of half-time between PB0400M and PB0800M is always above 1, indicating that the crystallization rate of sample with higher MW is smaller. However, with increasing Tc, the ratio between PB0110M and PB0800M is above 1 first and then becomes smaller than 1, indicating a possible change in the mechanism of crystallization for PB-1. Fu et al.12 observed the transition from chain-folded to chainextended crystallization in similar systems crystallized at different temperatures. When the crystalline thickness was smaller than the diameter of the volume occupied by the chains in the melt, the crystallization does not require a disentangle6403

dx.doi.org/10.1021/ma501272a | Macromolecules 2014, 47, 6401−6407

Macromolecules

Article

Figure 3 presents the evolution of temperature-dependent SAXS results of PB0110M isothermally crystallized at 30 °C in form II during heating. In the Lorentz corrected onedimensional scattering intensity distribution profiles (top), the scattering peak moved to smaller q with increasing temperature, indicating that long spacing increased gradually.

Figure 2. Top: DSC melting curvesof PB0110M in form II measured after isothermal crystallization at the indicated temperatures at a heating rate of 10 K/min. Middle: the evolution of weight fraction crystallinity (Φw) from DSC and lamellar thickness (dc) from SAXS in PB0110M as a function of crystallization temperature (Tc). Bottom: the evolution of melting temperature (Tm) of three PB-1 samples as a function of Tc.

the global fractions of the crystallized and noncrystallized chain parts.2 The bottom plot presents the melting temperature of three PB-1 samples as a function of crystallization temperature. Below the critical crystallization temperature, three samples all show constant melting peaks, indicating that the three samples with different MW all experienced recrystallization processes. However, above the critical crystallization temperature, the melting peaks increase with crystallization temperature. Overall, there is a positive MW dependency on melting temperature, indicating that when crystallized at same temperature the lamellae of PB-1 with higher MW were a little thicker.

Figure 3. PB0110M isothermally crystallized at 30 °C in form II and heated up to melt: temperature dependencies of selected onedimensional scattering intensity distribution profiles (top), the resultant correlation function curves (middle), and long spacing dac, lamellar thickness dc, amorphous thickness da, and linear crystallinity Φl = dc/dac (bottom). The dac and da can be obtained as shown in the inset, and dc was obtained by the equation dc = dac − da. 6404

dx.doi.org/10.1021/ma501272a | Macromolecules 2014, 47, 6401−6407

Macromolecules

Article

from a quiescent molten state has been respectively introduced by Strobl as follows:15,16

The resultant correlation functions in middle part of Figure 3 are presented, and the inset shows how the average thickness of the long spacing dac and amorphous layers da are derived from the correlation function curve. The assignment of the smaller value present in the correlation function curves of this PB-1 sample to the thickness of amorphous layers is due to its high crystallinity being larger than 50% as shown in DSC results. In the bottom of Figure 3, one finds that the long spacing and the average thicknesses of the lamellar and amorphous layers all increased during heating, while the linear crystallinity kept almost constant. The gradual increase in the crystal thickness during heating scan indicates the occurrence of a melting and recrystallization process forming thicker crystallites. The variation of the reciprocal lamellar thickness for the three PB-1 samples with different MW during heating up to the molten state is shown in Figure 4. For the three PB-1 samples

∞ Tmc

∞ (2σac n − 2σam)Tmc Δz −T≈ dc Δhcm

(4)

∞ (2σacs − 2σam)Tmc Δz dc Δhcm

(5)

∞ Tmc −T≈

Tac∞ − T ≈

2σacsTac∞ Δz Δhca

dc

(6)

∞ where T∞ ac and Tmc are the equilibrium melting and crystallization temperature; Δhcm and Δhca are the heat of transition from mesomorphic phase to crystalline phase and the one from crystalline phase to melt, respectively; Δz is the stem length increment per structural unit; σam, σacn, and σacs denote the surface free energy of the mesomorphic layers, native crystals, and stabilized crystals, respectively. Based the three equations, it is possible to explain the molecular weight dependency of crystallization and melting behaviors in PB-1. In the current condition, with varying MW, the three crystallization lines present different slopes. Because of the constant transition heat Δhcm and equilibrium crystallization temperature T∞ mc in eq 4 describing crystallization line, the parameter that affects the slope should be the surface free energy of native crystals (σacn). Some researchers have found that the surface free energy of crystallites in PE24,25 and PB-132 increases with molecular weight. Here, the similar behavior was observed that the surface free energy of three PB-1 samples increased with molecular weight, which resulted in the increase of slope in crystallization line. The dependency of surface free energy on molecule weight may contribute to the interlamellar strain. Most polymers begin to undergo intermolecular tangling in the melt over a certain molecule weight.60 As polymer crystallizes from such an entangled melt, interlamellar linkages should be generated which increase in number with molecule weight.61 An effect of interlamellar linkages was to strain the interlamellar layer by increasing its density,62 and the interlamellar strain should result in a high apparent surface free energy.63 Therefore, the surface free energy of “native crystallites” increases with molecular weight. As has been pointed out by Devoy64 and Lovering,65 the molecular weight dependency of the surface free energy could be related to the relative crystallite size in the chain direction. At low molecular weight, the crystallite size was comparable to the extended chain length, while at higher molecular weight, regular chainfolded crystallites formed and the size of the crystallites was very much smaller than the extended chain length. This opinion is consistent with the results in Table 2, where the chainextended crystallites in the sample with lower molecular weight formed at higher crystallization temperature and the chain-fold crystallites present in the one with higher molecular weight at any crystallization temperature. Therefore, the molecular weight changed the surface free energy of native crystals (σacn) and then influenced the crystallization line. It must be mentioned that evidence of a transition from chain-folded crystallization to extended-chain crystallization in the low molecular weight sample is not observed. Ideally, one could consider an observation of two crystallization lines for the PB0800M sample with lowest molecular weight because such a transition would involve a change in surface free energy and

Figure 4. Variation of dc−1of three PB-1 samples during heating scans subsequent to isothermal crystallization at the indicated temperature Tc as derived from the temperature-dependent SAXS experiments. Three crystallization lines (solid), a common recrystallization line (dash), and a melting line (dot) are extrapolated to dc−1 = 0.

with different MW, there exist three crystallization lines with different slopes, a common recrystallization line, and a common melting line. And the crystallization and recrystallization lines all extrapolate to a same limiting temperature (T∞ mc = 146 ± 2 °C), which is similar to the one in our previous work,59 while the melting line does to a lower equilibrium melting temperature (T∞ ac = 133 ± 2 °C). In addition, the melting line and recrystallization line intersect each other at a certain temperature and lamellar thickness, and the point, named Xs, marks the end of recrystallization.15,16 For an initial thickness below the critical lamellar thickness, recrystallization always happens before melting. If the initial crystal thickness is above the critical value, no recrystallization occurs, and the crystal directly melts. It is necessary to explorer the particular behavior that the crystallization line is dependent on MW, while the recrystallization line and melting line are both independent of MW. A theoretical description of the crystallization, recrystallization, and melting lines of semicrystalline polymers crystallized 6405

dx.doi.org/10.1021/ma501272a | Macromolecules 2014, 47, 6401−6407

Macromolecules

Article

crystallization and melting behavior was analyzed. The three crystallization lines were in different slopes but same equilibrium temperature due to the variation of surface free energy of the native crystals with molecular weight. Subsequent stabilization of the native crystallites into stable crystallites after crystallization leads to a similar surface free energy for all samples, resulting in a common recrystallization line and melting line.

thus the slope of the crystallization line. The observed effect of only one single crystallization line for each sample indicates that only one dominant crystallization type was effectively observed. Indeed, as is shown in Table 2, even at low crystallization temperature of 50 °C the lamellar thickness of the PB0800M sample has been very close to the Rg value. Therefore, within the accessible range of crystallization, only single crystallization line was obtained. After crystallization, the native crystals stabilized gradually and transformed into a state of lower surface free energy,16 where the difference of surface free energy of native crystals nearly disappears. Therefore, the stabilized crystals of three PB1 samples are in a state with similar surface free energy (σacs). Combined with recrystallization and melting lines of eqs 5 and 6, the parameters, which determine their slopes, are not influenced by the molecular weight. Thus, a common recrystallization line and a common melting line are observed. Here, the crystallization and melting behaviors of PB-1 in form II provided an experimental evidence for the multistage model of polymer crystallization.15,16 It was clear that the formation and melting of polymer crystals were not reverse processes. In different systems, a main feature for the multistage model was a crystallization line and a melting line as two independent features, with different slopes and limiting temperatures. And in some systems, a rerystallization process also took place. The crystallization line represents the transition from the mesomorphic phase to the native crystalline state whereas the recrystallization line denotes the transition from the mesomorphic phase to the stabilized crystallites, and the melting line described the melting of the stabilized crystallites. In this work, the molecular weight dependency of crystallization and melting behaviors in PB-1 was discussed. Three crystallization lines, a common recrystallization line, and a common melting line were observed. The three crystallization lines were due to the variation of surface free energy of native crystallites with molecular weight, which did not change the multistage characters of polymer crystallization. And the surface free energy of stabilized crystallites was almost not influenced by molecular weight, resulting in a common recrystallization line and a common melting line.

■

AUTHOR INFORMATION

Corresponding Author

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

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS We thank Prof. Gert Strobl of University of Freiburg for his insightful discussions and Prof. Zhonghua Wu and Dr. Guang Mo of BSRF for assistance during synchrotron SAXS measurements. 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, 1997; p 349. (2) Heck, B.; Hugel, T.; Iijima, M.; Sadiku, E.; Strobl, G. New J. Phys. 1999, 1, 17.1−29. (3) Heck, B.; Strobl, G.; Grasruck, M. Eur. Phys. J. 2003, E11, 117− 130. (4) Cho, T. Y.; Heck, B.; Strobl, G. Colloid Polym. Sci. 2004, 282, 825−832. (5) Strobl, G. Eur. Phys. J. 2005, E18, 295−309. (6) Iijima, M.; Strobl, G. Macromolecules 2000, 33, 5204−5214. (7) Hauser, G.; Schmidtke, J.; Strobl, G. Macromolecules 1998, 31, 6250−6258. (8) Hugel, T.; Strobl, G.; Thomann, R. Acta Polym. 1999, 50, 214− 217. (9) Al-Hussein, M.; Strobl, G. Eur. Phys. J. 2001, E6, 305−314. (10) Grasruck, M.; Strobl, G. Macromolecules 2003, 36, 86−91. (11) Heck, B.; Siegenführ, S.; Strobl, G.; Thomann, R. Polymer 2007, 48, 1352−1359. (12) Fu, Q.; Heck, B.; Strobl, G.; Thomann, Y. Macromolecules 2001, 34, 2502−2511. (13) Al-Hussein, M.; Strobl, G. Macromolecules 2002, 35, 1672− 1676. (14) Cho, T. Y.; Strobl, G. Polymer 2006, 47, 1036−1043. (15) Strobl, G. Eur. Phys. J. 2000, E3, 165−183. (16) Strobl, G. Prog. Polym. Sci. 2006, 31, 398−442. (17) Stribeck, N.; Almendarez Camarillo, A.; Cunis, S.; Bayer, R. K.; Gehrke, R. Macromol. Chem. Phys. 2004, 205, 1445−1454. (18) Stribeck, N.; Bösecke, P.; Bayer, R.; Almendarez Camarillo, A. Prog. Colloid Polym. Sci. 2005, 130, 127−139. (19) Stribeck, N.; Bayer, R.; Bösecke, P.; Almendarez Camarillo, A. Polymer 2005, 46, 2579−2583. (20) Stribeck, N.; Nöchel, U.; Almendarez Camarillo, A.; Roth, S. V.; Dommach, M.; Bösecke, P. Macromolecules 2007, 40, 4535−4545. (21) Almendarez Camarillo, A.; Roth, S. V.; Bösecke, P.; Buchner, S.; Krenn, K.; Gehrke, R.; Stribeck, N. J. Mater. Sci. 2007, 42, 6212−6221. (22) Stribeck, N.; Nöchel, U.; Funari, S. S. Macromolecules 2009, 42, 2093−2101. (23) Fatou, J. G.; Mandelkern, L. J. Phys. Chem. 1965, 69, 417−428. (24) Mandelkern, L. J. Polym. Sci., Part C 1966, 15, 129−162. (25) Leung, W. M.; John Manley, R. St.; Panaras, A. R. Macromolecules 1985, 18, 746−752. (26) Hoffman, J. D.; Miller, R. L. Macromolecules 1988, 21, 3038− 3051.

■

CONCLUSION In the present study, we have investigated the molecular weight dependency of crystallization and melting behaviors of PB-1 in form II by DSC and temperature-dependent SAXS techniques. The kinetics of crystallization of PB-1 in form II showed the different molecular weight dependencies with crystallization temperature. At low crystallization temperatures, the chainfolded crystallites were generated in all three PB-1 samples, and the crystallization rate decreased with increasing molecular weight. However, the crystallization rate of the sample with low MW was slower than the one with high MW over certain temperature. This was because that the chain-extended crystallites formed in the low-MW sample, accompanied by the disentanglement process resulting in a much slower crystallization than the one that chain-folded crystallites formed for the high-MW sample. The microstructure evolution during heating was followed for PB-1 in form II isothermally crystallized at different temperatures. Three crystallization lines, a common recrystallization line, and a common melting line were observed. Based on the multistage model of polymer crystallization, the effect of molecular weight on the 6406

dx.doi.org/10.1021/ma501272a | Macromolecules 2014, 47, 6401−6407

Macromolecules

Article

(63) Hoffman, J. D.; Lauritzen, J. I. J. Res. Natl. Bur. Stand., Sect. A 1961, 65A, 297. (64) Devoy, C.; Mandelkern, L. J. Polym. Sci., Part A-2 1969, 7, 1883−1894. (65) Lovering, E. G. J. Polym. Sci., Part C 1970, 30, 329−338.

(27) Fatou, J. G.; Macro, C.; Mandelkern, L. Polymer 1990, 31, 1685−1693. (28) Hocquet, S.; Dosière, M.; Tanzawa, Y.; Koch, M. H. J. Macromolecules 2002, 35, 5025−5033. (29) Cheng, S. Z. D.; Janimak, J. J.; Zhang, A. Q.; Cheng, H. N. Macromolecules 1990, 23, 298−303. (30) Bicerano, J. J. Macromol. Sci., Rev. Macromol. Chem. Phys. 1998, C38, 391−479. (31) Yamada, K.; Hikosaka, M.; Toda, A.; Yamazaki, S.; Tagashira, K. J. Macromol. Sci., Phys. 2003, B42, 733−752. (32) Cortazar, M.; Guzman, G. M. Makromol. Chem. 1982, 183, 721−729. (33) Monasse, B.; Haudin, J. M. Makromol. Chem., Macromol. Symp. 1988, 20/21, 295−302. (34) Acierno, S.; Grizzuti, N.; Winter, H. H. Macromolecules 2002, 35, 5043−5048. (35) Kajioka, H.; Yamada, K.; Taguchi, K.; Toda, A. Polymer 2011, 52, 2051−2058. (36) Natta, G.; Corradini, P.; Bassi, I. W. Makromol. Chem. 1956, 21, 240−244. (37) Danusso, F.; Gianotti, G. Makromol. Chem. 1965, 88, 149−158. (38) Azzurri, F.; Flores, A.; Alfonso, G. C.; Baltá Calleja, F. J. Macromolecules 2002, 35, 9069−9073. (39) Azzurri, F.; Flores, A.; Alfonso, G. C.; Sics, I.; Hsiao, B. S.; Baltá Calleja, F. J. Polymer 2003, 44, 1641−1645. (40) Su, F. M.; Li, X. Y.; Zhou, W. M.; Chen, W.; Li, H. L.; Cong, Y. H.; Hong, Z. H.; Qi, Z. M.; Li, L. B. Polymer 2013, 54, 3408−3416. (41) Gohil, R. M.; Miles, M. J.; Petermann, J. J. Macromol. Sci. Phys., Part B: Phys. 1982, 21, 189−201. (42) 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. (43) Chen, W.; Li, X. Y.; Li, H. L.; Su, F. M.; Zhou, W. M.; Li, L. B. J. Mater. Sci. 2013, 48, 4925−4933. (44) Wang, Y. T.; Jiang, Z. Y.; Wu, Z. H.; Men, Y. F. Macromolecules 2013, 46, 518−522. (45) De Rosa, C.; Auriemma, F.; de Ballesteros, O. R.; Esposito, F.; Laguzza, D.; Di Girolamo, R.; Resconi, L. Macromolecules 2009, 42, 8286−8297. (46) Miyoshi, T.; Mamun, A.; Reichert, D. Macromolecules 2010, 43, 3986−3989. (47) Miyoshi, T.; Mamun, A. Polym. J. 2012, 44, 65−71. (48) 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. (49) Cavallo, D.; Gardella, L.; Portale, G.; Müller, A. J.; Alfonso, G. C. Polymer 2013, 54, 4637−4644. (50) Cavallo, D.; Gardella, L.; Portale, G.; Müller, A. J.; Alfonso, G. C. Polymer 2014, 55, 137−142. (51) Cavallo, D.; Gardella, L.; Portale, G.; Müller, A. J.; Alfonso, G. C. Macromolecules 2014, 47, 870−873. (52) Danusso, F.; Gianotti, G.; Polizzotti, G. Makromol. Chem. 1964, 80, 13−21. (53) Product information, Lyondell Basell. Available from https:// polymers.lyondellbasell.com/portal/site/basell/polybutene-1, 2013. (54) Rubin, I. D. J. Polym. Sci. 1964, B2, 747−749. (55) Glatter, O.; Kratky, O. Small-Angle X-ray Scattering; Academic Press: London, 1982. (56) Strobl, G.; Schneider, M. J. J. Polym. Sci., Polym. Phys. 1980, 18, 1343−1359. (57) Strobl, G.; Schneider, M. J.; Voigt-Martin, I. G. J. Polym. Sci., Polym. Phys. 1980, 18, 1361−1381. (58) Brandrup, J.; Immergut, E. H.; Grulke, E. A. Polymer Handbook, 4th ed.; John Wiley & Sons: New York, 1999. (59) Wang, Y. T.; Jiang, Z. Y.; Fu, L. L.; Lu, Y.; Men, Y. F. Macromolecules 2013, 46, 7874−7879. (60) Fox, T. G.; Loshaek, S. J. Appl. Phys. 1955, 26, 1080−1082. (61) Keith, H. D.; Padden, F. J.; Vadimsky, R. G. Science 1965, 150, 1026−1027. (62) Schultz, J. M.; Robinson, W. H.; Pound, G. M. J. Polym. Sci., Part A-2 1967, 5, 511−533. 6407

dx.doi.org/10.1021/ma501272a | Macromolecules 2014, 47, 6401−6407