Article pubs.acs.org/Macromolecules
Stretching Temperature Dependency of Lamellar Thickness in StressInduced Localized Melting and Recrystallized Polybutene‑1 Yaotao Wang, Zhiyong Jiang, Lianlian Fu, Ying Lu, 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 structural parameters such as long spacing and thicknesses of crystalline lamellae and amorphous layers of polybutene-1 crystallized from quiescent melt at different isothermal temperatures and established by tensile stretching the isothermally crystallized samples at different temperatures via a stress induced localized melting and recrystallization process were determined by small-angle X-ray scattering experiments. In the two cases, almost parallel crystallization lines where reciprocal lamellar thickness and temperature at which the structure formed are in linear relationship were observed. As it turns out, the samples crystallized from the quiescent melt possess much thicker crystalline lamellae than those obtained via stretching at same temperature. The results yield two limiting temperatures for the formation of lamellae with infinite thickness in the two cases which can be understood by considering the influence of the stress on the thermodynamic parameters in the melt and thus the crystallization process.
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INTRODUCTION The understanding of crystallization in bulk polymers and morphological changes on a microscopic level during tensile deformation are essential to controlling the mechanical properties of semicrystalline polymeric materials and thus to possible improvement. When the fundamental of the structure of semicrystalline polymersstacks of layer-like crystalline lamellae with thickness in the nanometer-range embedded in an amorphous matrix1was revealed in the 1950s, considerations about the mechanism of formation of this structure started immediately. Over the years, various conceptual models about the crystallization in bulk polymers have been proposed. One approach developed by Hoffman and Lauritzen gained the ascendancy.2−6 The model interprets the characteristic exponential law for the variation of the crystallization time with the crystallization temperature as originating from an increase in an activation barrier. This “secondary nucleation” then determines the crystal growth rate. The longitudinal extension of the nucleus fixes also the thickness of the crystalline lamellae. And the maximum for a crystal thickness is just above the stability limit of the crystallites as determined by the Gibbs−Thomson equation. In recent years, Strobl7−14 introduced a multistage model which is different from the conventional nucleation and growth theory. By carrying out the time- and temperature- dependent small-angle X-ray scattering (SAXS) experiments, they found that the formation of crystallites is governed by the same general laws. Both the crystallization temperature Tc and the melting temperature Tm are linearly dependent on the reciprocal crystalline layer thickness dc−1, but with different slopes and limiting temperatures for dc−1 → 0. These two lines are termed as crystallization line and melting line of a specified sample, © XXXX American Chemical Society
respectively. The occurrence of two well-defined independent lines may be understood as an indication that the transformation from the melt into the partially crystalline state is generally a two-step process. The process starts with an attachment of chain sequences from the melt onto a growth face of a mesomorphic layer with minimum thickness, which then spontaneously thickens. When a critical thickness is reached, the layer solidifies immediately by the formation of block-like crystallites. The last, but equally important step in the crystal development is the stabilization of the crystallites in time, leading to a further decrease in the Gibbs free energy. The microstructure and crystallization conditions could influence the crystallization process. Heck et al.11 observed the effect of diluents on the crystallization line in poly(ethyleneco-octene). The addition of diluents of C16H32 and C15H30 introduced a shift of the crystallization line to lower temperature. This was explained by the inclusion of diluent molecules in the mesomorphic phase, then affected the crystal formation and changed the crystal thickness. Miyoshi et al.15 found that the polybutene-1 (PB-1) sample with low isotacticity, which crystallized directly into form I, and the one with high isotacticity, which experienced the crystallization into form II and a subsequent solid−solid transition into form I, showed different supercooling dependence of lamellar thickness dc. The difference was attributed to a huge difference in mobility of chain segments in crystalline phase in forms I and II.16 Extremely fast dynamics of chain segments in form II led Received: June 26, 2013 Revised: August 30, 2013
A
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sample-to-detector distance was 1760 mm, and the effective range of the scattering vector q (q = 4π(sinθ)/λ, where 2θ is the scattering angle and λ is the wavelength) was 0.071−0.796 nm−1. Each SAXS pattern obtained in the center of the sample was collected within 30 min which was then background corrected and normalized using the standard procedure. One dimensional scattering intensity distributions have been obtained by integrating the two-dimensional scattering patterns. For the isotropic melt crystallized samples a 360° integration of intensity at each scattering vector q has been performed. Whereas for the deformed samples showing oriented structural features integrations were confined within ±20° along the stretching direction. Beside one-dimensional scattering intensity distributions integrated from 2D SAXS patterns, the technique of one-dimensional electron density correlation function analysis has been also used to give detailed structural information on the systems. The electron density correlation function K(z) can be derived from the inverse Fourier transformation of the experimentally intensity distribution I(q) as follows:35−37
to large variations in dc. In contrast, the direct crystallization of form I into a fixed crystal resulted in very thin lamellae. Under stretching, the semicrystalline polymer transforms from an original isotropic spherulitic morphology into a highly oriented fibrillar one.17 As it turns out, based on the results of true stress−strain curves and recovery property studies,18−21 block slippage22−26 within the crystalline lamellae takes place first, and followed by the stress-induced melting and recrystallization27−32 starting at certain strain determined by the stability of crystalline blocks and the state of entangled amorphous network.33 Recently, we presented the phase transition from form I to II upon stretching in PB-1 at the elevated temperature.32 On the basis of the fact that a solid state I to II phase transition cannot take place due to the restriction in chain conformations and lattice dimensions in both phases, the observed occurrence of transition from form I to form II must proceed via a two-step process. First, those form I crystallites with their polymer chain direction tilted with respect to the stretching direction underwent a stress induced melting process. Second, the freed polymer chain segments recrystallized into metastable form II crystallites with their chain direction preferentially aligned along the stretching direction. This result is considered to provide a direct evidence for the stress-induced melting and recrystallization mechanism. The process is similar to the crystallization in the quiescent melt with both experiencing the process from the molten state to the crystalline phase. In this study, we explored the relationship between melt crystallization and stress-induced localized melting and recrystallization behavior in PB-1. By discussing the effect of stress on the crystallization behavior, we made an attempt to understand the nature of the phenomenon.
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∞
K (z) =
EXPERIMENTAL SECTION
d0 d
∞
∫0 I(q)q2 dq
(2)
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 because of isotropically distributed stacks of parallel lamellar crystallites in the sample. 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. For comparison, all the one-dimensional scattering intensity distribution curves were performed with Lorentz correction in this article, no matter the lamellae crystallized in the bulk or generated from the stress-induced melting and recrystallization. It must be mentioned that for perfectly oriented lamellar stack structure no Lorentz correction should be performed. However, when the structures are not perfectly oriented like in the current case, the scattering intensity at certain scattering vector q would still be distributed over certain area on the reciprocal sphere with a radius of q which certainly needs a corresponding correction. Since there is no unique recipe to correct the intensity distribution for such case, Lorentz correction was used also for samples after deformation. DSC measurements were conducted to acquire the melting temperature (Tm) and weight fraction crystallinity (Φw). A DSC1 Stare System (Mettler Toledo Instruments, Swiss), which had been calibrated for temperature and melting enthalpy by using indium as a standard, was used during the experiments at a heating rate of 10 K/ min. For the calculation of crystallinity, a value for the heat of fusion at 100% crystallinity of ΔHid = 125 J/g38 were used. And the DSC crystallinity is summarized in below table.
The PB-1 is produced by BASELL Polyolefins with a trade name of PB0110 M with a melt flow rate (MFR) of 0.4 g/10 min (190 °C/2.16 kg). It has a weight-average molecular weight of 4.39 × 105 g/mol.34 Pellets of PB-1 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 (0, 30, 40, 50, 60, 70, 80, 90, and 100 °C) and held isothermally for more than 5 h to complete the crystallization in the samples. The isothermally crystallized PB-1 samples were stored at room temperature for 1 month to allow a complete phase transition from the metastable tetragonal phase II to the stable hexagonal phase I. “Dog bone” tensile bars with dimensions of 10 × 5 × 0.5 mm3 were obtained with the aid of a punch. The samples crystallized at 0, 30, and 40 °C were stretched at different temperatures (30, 40, 50, 60, 70, 80, 90, and 100 °C) using a portable tensile testing machine (TST350, Linkam, U.K.) until an engineer strain of 300%, then relaxed. The cross-head speed during stretching was kept at 50 μm/s. In order to measure the strain of the deformed area located at certain spots on the samples accurately, optical photo images of the samples were employed. The Hencky measure of strain εH is used as a basic quantity of the extension, which is defined as
εH = 2 ln
∫0 I(q)q2 cos(qz) dq
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RESULTS AND DISCUSSION Figure 1 shows the SAXS results in PB-1 with the stable crystalline modification of form I crystallized at different temperatures. In the one-dimensional scattering intensity distribution profiles (top), the scattering peaks move to smaller q with the increase of crystallization temperature, meaning that the long spacing increased gradually. In the middle, the resultant correlation functions are presented as a function of crystallization temperature and the inset shows how the average thickness of the amorphous layers da and long spacing dac are derived by the correlation function. 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 given in Table 1. As shown in the bottom, the long spacing, the average thickness of the lamellar and amorphous layers all increase with increasing
(1)
where d0 and d are the widths of undeformed and deformed area, respectively. SAXS experiments were carried out on a modified Xeuss system of Xenocs France equipped with a semiconductor detector (Pilatus 100 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. The wavelength of the X-ray radiation was 0.154 nm. The B
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Table 1. DSC Crystallinity of PB-1 at Different Conditionsa Tc (°C)
Φw (%)
Td (°C)
Φw (%)
Tc (°C)
Φw (%)
quench to 0 30 40 50 60 70 80 90
50.3 54.1 56.1 56.9 58.9 63.7 65.9 67.8
30 40 50 60 70 80 90 100
44.0 43.6 41.7 41.3 46.4 51.6 50.4 52.2
30 40 50 60
43.9 44.7 42.8 45.3
a
Left: PB-1 crystallized at different temperatures (Tc). Middle: quenched PB-1 after stretched at different temperatures (Td). Right: PB-1 crystallized at different temperatures (Tc) after being stretched at 75°C.
With the increasing of strain, the SAXS pattern transforms from an isotropic ring to four-point scattering diagram at moderate strains and eventually to highly anisotropic scattering intensity distribution with two broad scattering peaks aligned in the stretching direction. The middle plot shows the one-dimensional scattering intensity distributions along the stretching direction. The scattering intensity distribution changes from a one peak curve to a two peaks one after the strain-hardening point indicating the formation of new lamellar stacks of different long period. Evidently, the scattering peak denoting the original lamellar structure weakens accompanied by a gradual increase of the new peak showing the gradual transition at larger strains. The two scattering peaks can be separated by a fitting procedure present in the inset of the middle plot. The bottom plot shows the evolution of the two scattering peaks during the strainhardening stage. The position of the scattering peak from newly established lamellar stacks locates at larger q than the one from the original lamellar stacks. With the increase of strain, the scattering peak position of the new lamellar stacks keeps almost constant. This result is in agreement with previous findings that the long spacing of new lamellar stacks obtained after tensile stretching is determined by the drawing temperature, regardless of the initial morphology.21,28 After being relaxed, the sample showed a slightly reduced long spacing due to the relaxation of highly oriented amorphous phase and the orientation of the crystallites. Obviously, the structure of the new lamellar stacks generated through stress-induced melting and recrystallization depends on the drawing temperature. It is reasonable to obtain lamellar stacks generated at different drawing temperatures and to compare them with the lamellae crystallized in the bulk from the molten state. Figure 3 shows the SAXS results of PB-1 samples stretched at elevated temperatures and relaxed. The top shows the one-dimensional scattering intensity distributions along the stretching direction. In the middle, the scattering peaks from new lamellar stacks are acquired by the fit procedure present in Figure 2 (the inset in the middle plot). Clearly, the scattering peak moves to smaller q with the increase of drawing temperature indicating an increase in long spacing. With the help of the correlation function analysis, the long spacing, the thickness of lamellae and amorphous layers were obtained and present in the bottom plot of Figure 3. Here, the DSC crystallinity of less than 50% of the deformed samples was used to assist the assignment of the smaller value in correlation function profiles to crystalline lamellar thickness. Clearly, structural parameters of long spacing, lamellar and amorphous
Figure 1. Top: the one-dimensional scattering intensity distribution profiles in PB-1 crystallized at elevated temperatures; Middle: the resultant correlation function curves, where the long spacing dac and the amorphous thickness da can be obtained as shown in the inset and the lamellar thickness dc was obtained as dac − da. Bottom: the evolution of the dac, dc, da, and αc,l as a function of crystallization temperature.
crystallization temperature. The linear crystallinity, however, shows no crystallization temperature dependency and keeps essentially constant. In the current range of crystallization temperature, the lamellar thickness varies from about 13 to 35 nm. Figure 2 shows the SAXS results of quenched PB-1 (crystallized at 0 °C) stretched at 75 °C at different strains. C
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Figure 2. Top: True stress σH - strain εH curve and the selected 2D SAXS patterns of quenched PB-1 stretched at 75 °C at different strains as indicated on the graphs. Stretching direction horizontal. Middle: the corresponding one-dimensional scattering intensity distribution profiles integrated along the stretching direction, where the two scattering peaks shown in the sample after the strain-hardening can be separated by a fitting procedure present in the inset. Bottom: the evolution of the scattering peaks from the two populations of lamellar stacks. Figure 3. Top: the one-dimensional scattering intensity distribution profiles integrated along the stretching direction in PB-1 stretched at elevated temperatures and relaxed. Middle: the scattering curves from newly established lamellae obtained by fitting the curves in the top plot. Bottom: the long spacing (dac), the thickness of lamellae (dc) and amorphous phase (da) as a function of stretching temperature.
thicknesses of all three samples crystallized at different conditions after stretching at different temperatures follow the same master curve irrespective of their original structure. This result again presents the process of stress induced localized melting and recrystallization during stretching the PB1 samples. Notably, the linear crystallinity after stretching keeps also essentially constant at 40%. In previous works, Strobl et al. carried out investigations with the crystallization line dc−1 vs T on syndiotactic polypropylene,9,39 polyethylene,11 poly(ε-caprolactone),7 isotactic polypropylene,40 and isotactic poly(butene-1)10 for an understanding of the molecular processes involved in the formation of polymer crystallites. Because of the feature that the thickness
of the newly established lamellae was determined by the drawing temperature upon stretching, it is possible to construct a similar crystallization line to explore the stress-induced melting and recrystallization process further. As is known, PB-1 crystallizes into form II crystalline structure from the melt followed by a solid−solid phase transition into form I upon D
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storage.41−44 Similarly, during the stress-induced melting and recrystallization process, the system also formed at the beginning form II crystallites which can be transformed into form I quickly due to the existence of stress.32,45 Phase transformation from form II to I involves an extension of the 11/3 helical conformation (form II) into the 3/1 helix (form I). The ratio between the axial repeating units of these conformations is 1.12.46 Therefore, it is necessary to acquire the thickness in form II (the original crystalline phase after crystallization) by the relationship dc (II) = dc (I)/1.12. From these experiments we derived the relationships between lamellar thickness dc (II), crystallization temperature Tc and drawing temperature Td. The results obtained for isothermally crystallized samples from the random melt state, and the samples with different lamellar thicknesses stretched at different temperatures are collected in Figure 4. Clearly, two
structural unit, and σacn and σam denote the surface free energy of the native crystal layer and the mesomorphic layer, respectively. With the help of this equation, we might speculate the underlying mechanism in the current case that a significant reduction in dc was observed when the sample was crystallized from a localized oriented melt state. A close look at the above equation reveals that dc hardly changes if the crystallization of the system proceeded in the same manner of via a mesomorphic phase during the localized melting and recrystallization. This is due to the fact that all thermodynamic parameters included in the equation that governs the crystalline lamellar thickness are related mostly to the mesomorphic and crystalline phase only independent of the state of amorphous phase. This result leads to a consideration of validity of the above equation in current case of stress induced melting and recrystallization. Indeed, if the localized crystallization process from a stretched molten state proceeds via a direct rout to crystalline phase without passing though the mesomorphic phase, a reduction in lamellar thickness is expected. In such case, the crystallization line can be described by the following equation: Tc∞ − T ≈
crystallization lines with similar slope were observed. The data for the quenched sample deviated slightly to the isothermal crystallization line because it crystallized during the quenching process above 0 °C. The lamellae established by stress induced localized melting and recrystallization were much thinner than the one crystallized in the bulk at the same temperature. The two crystallinization lines yield two different limiting temperatures Tc∞ when extrapolated to infinite crystalline thickness (dc → ∞) at about 146 °C for lamellae crystallized in the bulk and 294 °C for lamellae generated by stress-induced melting and recrystallization, respectively. What is the difference between the two crystallization processes, which influenced the crystallization behavior? Theoretical description of the crystallization line of polymers crystallized from a quiescent molten state has been given by Strobl as12−14
dc
ga − gc ≈ ΔS(Tc∞ − T )
Here ΔS denote the entropy per monomer, and Tc∞ represents the equilibrium crystallization temperature associated with a sample of macroscopic size. The relationship holds as the system crystallizes without passing though the mesomorphic phase in the case of localized melting and recrystallization during stretching. As was discussed above, the melt with oriented chain segments upon stretching influenced the crystallization behavior through affecting the equilibrium condition. Stretching effectively increase the free energy of the melt state. Therefore, the difference ga − gc becomes larger and then the value of ΔS(Tc∞ − T) increases. Since the entropy S in the melt with oriented chain segments decreases, the change in entropy (ΔS) in the phase transition thus decreases. Clearly, Tc∞ must increase to keep the product increasing.
(2σacn − 2σam)Tc∞ Δz Δhcm
Δhca
Here Δhca denotes the heat of transition from amorphous phase to the crystalline one. As was proven experimentally,14 the increase in the surface free energy term at the numerator is much smaller than the increase in Δhca at the denominator of the equation leading to a strong decrease in dc. In an effort to understand the very strong increase in the limiting temperature of crystallization at infinite lamellar thickness for the stretching induced crystallization, we perform a simple thermodynamic analysis. During tensile deformation, those crystallites with their polymer chain direction tilted with respect to the stretching direction underwent a stress induced localized melting with the freed polymeric chains then preferentially aligned along stretching direction. Here the melt due to stretching was different from the random melt in the bulk because the stress changes the thermodynamic parameters of the system. Compared with the random melt, the enthalpy H in the melt induced by stress slightly increased accompanied by a decrease in the entropy S.47 Therefore, the Gibbs free energy G increased because G = H − TS. For Gibbs free energies ga and gc of a monomer in the melt and in an infinite perfect crystal, respectively, a linear approximation for the difference ga − gc reads:
Figure 4. Relations between the inverse crystalline lamellar thickness dc−1 in form II and the crystallization temperature Tc for isothermally crystallized samples from the random molten state, and the drawing temperature Td for the samples from the stress-induced melting and recrystallization.
Tc∞ − T ≈
2σacnTc∞ Δz
dc
where Δhcm is the heat of transition from mesomorphic phase to a crystalline phase, Δz is the stem length increment per E
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Indeed, this is just what has been observed that the value of Tc∞ changed from 146 °C for lamellae crystallized from quiescent melt to 294 °C for lamellae generated by stress-induced localized melting and recrystallization.
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CONCLUSION For isothermally crystallized samples from the quiescent melt state and the samples recrystallized from localized stress induced oriented melt state, the crystallization lines in PB-1 representing relationship between reciprocal lamellar thickness 1/dc(II) and crystallization temperature Tc and drawing temperature Td were obtained by means of SAXS experiments. The two crystallization lines with similar slope resulted in two different equilibrium crystallization temperatures. The lamellae generated from stress-induced melting and recrystallization process were much thinner than the one crystallized in the bulk from quiescent melt at the same temperature. The difference was interpreted as the change of thermodynamic parameters in the melt induced by stretching.
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AUTHOR INFORMATION
Corresponding Author
*E-mail: (Y.M)
[email protected]. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This work is supported by the National Natural Science Foundation of China (21134006). We thank Prof. G. Strobl for his insightful discussion during the revision of this manuscript.
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REFERENCES
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