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Effect of mesophase separation and crystallization on the elastomeric behavior of olefin multi-block copolymers
Polymer 52 (2011) 5221e5230
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Effect of mesophase separation and crystallization on the elastomeric behavior of olefin multi-block copolymers Guoming Liua, Yu Guana, Tao Wena, Xiwei Wanga, Xiuqin Zhanga, Dujin Wanga, *, Xiuhong Lib, Joachim Loosc, Hongyu Chend, Kim Waltone, Gary Marchande a
Beijing National Laboratory for Molecular Sciences, CAS Key Laboratory of Engineering Plastics, Institute of Chemistry, Chinese Academy of Sciences, Beijing 100190, China Shanghai Synchrotron Radiation Facility, Shanghai Institute of Applied Physics, Chinese Academy of Sciences, Shanghai 201204, China School of Physics and Astronomy, University of Glasgow, Glasgow G12 8QQ, Scotland, UK d The Dow Chemical (China) Company Limited, Shanghai 201203, China e The Dow Chemical Company, Freeport, TX 77541, USA b c
a r t i c l e i n f o
a b s t r a c t
Article history: Received 30 July 2011 Received in revised form 6 September 2011 Accepted 8 September 2011 Available online 14 September 2011
The mesophase separation and crystallization as well as the elastomeric properties of olefin multi-block copolymers (OBCs) are studied. The solid state morphologies of the OBCs are determined by the competition between mesophase separation and crystallization of hard blocks. The OBC with lower DC8 (octene content difference between soft and hard blocks) displays a spherulitic superstructure, while the OBC with higher DC8 exhibits mesophase separation. The two OBCs show very different elastomeric properties. Wide-angle X-ray scattering shows that the two OBCs have different deformation mechanisms. It is proposed that the cooperative deformation of crystalline phase makes an important contribution to the elastomeric properties. Ó 2011 Elsevier Ltd. All rights reserved.
Keywords: Elastomer Deformation Morphology
1. Introduction Due to the diversity in morphologies and properties, block copolymers have attracted widespread research interest. Below the orderedisorder transition (ODT) temperature, the melt of a block copolymer can self-assemble into spheres, rods, lamellae, cylinders or more complex structures. For a di-block copolymer, the melt phase behavior depends on two parameters: f, the volume fraction of one component and cN, in which c is the FloryeHuggins interaction parameter and N is the degree of polymerization. Di-block copolymers with both blocks being non-crystalline preserve the microphase separated morphology of the melt during cooling. For amorphous-crystalline di-block copolymers, on the other hand, two competing mechanisms drive solidification, i.e., microphase separation and crystallization. Whether the microphase separated structure is preserved or destroyed during cooling depends on the state of the amorphous blocks upon crystallizing and the segregation strength [1,2]. When the non-crystalline blocks are glassy during crystallization, the overall melt structure is retained.
* Corresponding author. Tel./fax: þ86 10 82618533. E-mail address: [email protected] (D. Wang). 0032-3861/$ e see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.polymer.2011.09.009
Crystallization occurs within the nano-scale domains prescribed by microphase separation [3e5]. However, when the amorphous blocks are rubbery during crystallization, microphase separation competes with crystallization, and the ultimate morphology depends on the segregation strength. Weakly segregated block copolymers undergo enormous structural rearrangement during crystallization, resulting in the formation of lamellar crystallites and even spherulitic superstructures [6e8]; while strong segregated polymers can preserve the morphology formed in the melt state [8e11]. The different crystallization behavior of block copolymers has been classified by Loo et al. [12] as “break out”, “templated” or “confined”. Crystallization from a homogeneous melt for block copolymers containing crystalline blocks and amorphous blocks can also induce microphase separation by the expulsion of amorphous blocks [13,14]. More information about structure, properties and theory of well-defined di- or tri-block copolymers can be found elsewhere [2,15]. Recently, the Dow Chemical Company developed olefin block copolymers (OBCs, INFUSEÔ) based on a chain shuttling catalyst technology [16]. The block copolymers synthesized by such technology consist of ethyleneeoctene segments with low octene concentrations in the hard or crystalline blocks and high octene concentration in the soft or amorphous blocks. This new type of
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2. Experimental
Fig. 1. Sketch showing the molecular architecture of OBC.
olefin block copolymer has a statistical multi-block architecture with a distribution in block lengths and the number of blocks per chain. Compared to random ethyleneeoctene copolymers (EOC), the block architecture imparts a substantially higher melting temperature, a higher crystallization temperature while maintaining a low glass transition temperature. As mentioned above, one of the basic questions in the area of block copolymer research is how the morphology in the melt and potential reorganization during crystallization influence morphology formation and properties of the copolymer. The miscibility (cN) of hard segments and soft segments in OBC depends on molecular weight and the DC8 (the octene content difference between soft and hard segments) [17,18]. Hustad et al. [19] found that polydisperse di-block OBC with DC8 z 35 mol% formed lamellar structures in the solid state with interesting photonic properties. Li et al. [20] reported that nearsymmetric polydisperse di-block OBCs with DC8 exceeding 25 mol% and number averaged molecular weights of 69e93 kg/mol selfassembled into well-defined lamellar structures. After crystallization, the lamellae-type structure was preserved. Crystallites in such lamellae were isotropic oriented and the crystallization kinetics was consistent with a spread growth behavior. The new mode of crystallization within heterogeneous block copolymers was termed as “pass through”. Multi-block OBCs have been assumed to have homogeneous melts in relatively earlier studies based on optical microscopy observations of spherulitic morphology [21,22]. However, phase separation was observed for an OBC with high DC8 of 21 mol% [23], though the melt morphology was still unknown, i.e. whether the solid state phase separated morphology is a replica of the melt structure or is formed during the crystallization process. Recently, based on rheological analysis a first attempt has been made to better understand the melt structure of multi-block OBCs [24]. Large composition fluctuations and mesophase separation transition at temperatures well above the melting point were found in multi-block OBC with DC8 higher than 16 mol%. Jin et al. [25] found the morphology of multi-OBC resulted from the competition between crystallization and mesophase separation. It is well known that structure and mechanical correlation is very important in the application of polymeric materials. In this study, our objective is to understand how the solid state morphology (mesophase separation and crystallization) influences the mechanical properties of OBCs. We explored two OBCs with similar molecular architecture and crystallinity. The solid state morphology was studied by microscopy and small-angle X-ray scattering. Two dimensional (2-D) wide-angle X-ray scattering measurement was carried out on stretched OBCs.
Two multi-block OBCs synthesized via a novel chain shuttling polymerization technology were supplied by the Dow Chemical Company. Fig. 1 shows a sketch of the general molecular structure of OBC with low octene concentration in the hard and high octene concentration in the soft segments. Molecular weight, molecular weight distribution, density and viscosity of the two samples are similar. The two OBCs have comparable low octene content in their hard segments but significantly different octene content in their soft segments. The amount of chain shuttling agent (CSA) illustrated as the Zn content in the samples was determined by X-ray florescence. Because the soft segment octene content in the two OBCs is 22.6 mol% and 35.7 mol%, the OBCs were designated as S23 and S36, respectively. Relatively lower amount of CSA was used in S36, which would principally result in longer segments. The molecular information of the OBCs is shown in Table 1. Note that the crystallinity of S23 is slightly higher than that of S36, because the hard segment octene content in S23 is slightly lower than that in S36. Samples with thickness of 1 mm for mechanical testing and X-ray scattering experiments and ca. 20 mm for optical microscopy were prepared by compression molding at 190 C. Thin films with about 100 nm thickness were made by spin-coating xylene solution at 110 C on carbon covered mica at 80 C. Subsequently, the thin film samples were heated to 190 C for 5 min and then either quenched into liquid nitrogen or slowly cooled to room temperature at a rate of 2 C/min. The thermally treated thin films were cut into pieces for further atomic force microscopy (AFM) and transmission electron microscopy (TEM) investigation. Uniaxial tensile tests were conducted on 1 mm thick ISO-37 Type 2 specimens cut from compression-molded plaques. Specimens were stretched on an Instron 3365 universal mechanical testing machine at room temperature with a crosshead speed of 300 mm/min. Engineering stress was defined conventionally as force per initial cross-sectional area. Polarized light optical micrographs were obtained with a BX51 Olympus optical microscope equipped with a Cannon (50D) camera. The free surface of the thin films was imaged in air with a commercial atomic force microscope (AFM) Nanoscope IIIa from Digital Instruments operated in tapping mode. For TEM observations, specimens were prepared by floating the film samples on water and transferring them onto copper grids. A JEM-2200S TEM was used with an operation voltage of 200 kV. X-ray measurements were carried out at the beamline BL16B1 in the Shanghai Synchrotron Radiation Facility (SSRF). The wavelength of the radiation source was l ¼ 1.24 Å. Scattering patterns were collected by a MAR CCD (MAR-USA) detector with a resolution of 2048 2048 pixels (pixel size: 79 79 mm2). Image acquisition time was 200 s. The sample to detector distance was 3117.4 mm for SAXS and 193.2 mm for WAXS. WAXS patterns for OBC samples with different drawing ratio were collected. The samples were stretched to a specific strain and then fastened on a home-made holder to keep the strain. All the X-ray images were corrected for background scattering, air scattering and beam fluctuations.
Table 1 Molecular characteristic of OBCs. Sample code
Density (g/cm3)
Mn (kg/mol)
Mw (kg/mol)
Octene content in soft segmenta (mol%)
Octene content in hard segmenta (mol%)
Zn content (ppm)
Hard segment (wt%)
Xc, DHb (wt%)
S23 S36
0.867 0.863
76 74
180 249
22.6 35.7
1.13 2.06
177 104
16 18
10 7.8
a b
Determined from 13C NMR. Determined by enthalpy of fusion.
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The melting and crystallization behavior of the OBCs was examined with a PerkineElmer differential scanning calorimeter (DSC 7). The instrument was calibrated with indium before measurement. Temperature scans were performed in the temperature range from 0 to 200 C at a heating/cooling rate of 10 C/min under nitrogen atmosphere. The typical sample weight was about 2e4 mg. 3. Results and discussion 3.1. Mechanical behavior of the OBCs Engineering stressestrain curves of quenched S23 and S36 at room temperature are plotted in Fig. 2. The two OBCs showed typical elastomeric behavior, specifically no yielding and uniform deformation to high strains. They had high elongations, and did not fracture up to an extension ratio of 13. Obvious differences can be found between the two polymers. S23 exhibited strain hardening at high strains while S36 had almost no strain hardening. The initial elastic modulus measured by the slope of the linear part of the engineering stressestrain curves is 6.1 0.4 MPa and 0.95 0.07 MPa for S23 and S36 (Table 2), respectively. In addition, the mechanical data is expressed in terms of modulus G usually used for elastomers [26]:
2 G ¼ s= l 1=l where l is the extension ratio and s is the engineering stress. The modulus gradually decreases at low strains for the two OBCs. In the whole tensile range, the modulus of S23 is much higher than that of S36. The modulus of S23 has an upswing at high strains, while the modulus of S36 has a continuous decrease (Fig. 3). It was also found that for a same material, independent of quenching or slow cooling conditions the general mechanical behavior was similar. Hiltner and coworkers [27,28] found that the slip-link theory well described the stressestrain curves of crystalline thermoplastic elastomers, including ethyleneeoctene random copolymers (EOC) and OBC. The initial modulus and the plateau modulus were primarily determined by the sliplink density, which was shown to scale with crystallinity, while the tightly topological entanglements at high strain served as crosslinks. In our experiments, the crystallinity of S23 is slightly higher than that of S36. Dividing the initial modulus by crystallinity, the normalized initial modulus is 62 4 MPa for S23 and 12 1 MPa for S36 (Table 2), respectively,
Fig. 2. Stressestrain curves of S23 and S36 at room temperature.
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Table 2 Mechanical data of S23 and S36. Sample code
Tangent modulus G (MPa)
G/Xc (MPa)
S23 S36
6.1 0.4 0.95 0.07
62 4 12 1
indicating that the normalized initial modulus of S23 is 5 times larger than that of S36. The plateau modulus of S23 is also much larger than that of S36. Obviously, this result cannot be explained by the previous scaling arguments. Molecular weight seems to play a minor role, as S36 with higher Mw does not impart higher strength compared to S23 with lower Mw.
3.2. Morphology observation To understand the different elastomeric properties of the two multi-block OBCs, we carried out morphological study via optical microscopy, AFM and TEM. For high cooling rates (>10 C/min), no spherulitic structure can be seen in either S23 or S36. When the cooling rate is lower, however, S23 has obvious spherulites (Fig. 4). On the contrary, no spherulitic superstructure can be observed in S36 even at a very low cooling rate of 1 C/min. To further reveal the nanostructure of the OBCs, we have prepared and investigated spin coated thin films quenched from the melt. Fig. 5 shows AFM phase images and bright field TEM images of the OBC samples. In the present case, AFM phase images display higher contrast and resolution than topography images because of the large difference in modulus between the crystalline and amorphous phase. As shown in Fig. 5A, crystalline lamellae and lamellar stacks can be seen in S23, which are brighter due to their relatively higher modulus. The lamellae have a thickness less than 10 nm and a length of 200e300 nm, and they are homogenously distributed in the amorphous matrix. The lamellar stacks consist of a few lamellar crystals, which are confirmed by the corresponding TEM observations shown in Fig. 5C. For sample S36, the AFM image in Fig. 5B shows island-like bright aggregates which consist of several very short lamellar crystals that embed in the “darker” amorphous matrix (Fig. 5D is the corresponding TEM image). Based on the above observations, we notice that the two OBCs have clearly different structural characteristics: relatively long but uncorrelated lamellae for S23, and densely packed lamellae in isolated regions for S36. In other words, S36 exhibits “mesophase separation” (following the term as introduced in literatures [20,24,25]), while crystallization dominated structure with no clear mesophase separation occurs in S23.
Fig. 3. Modulus G vs. inverse extension ratio of S23 and S36 at room temperature.
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Fig. 4. Polarized light optical microscopy images of S23 (A) and S36 (B) cooled at 2 C/min.
The final morphology of a block copolymer containing crystallizable segments is a result of both the melt structure and the crystallization process. An indirect approach to detect the melt morphology is to check whether the morphology after crystallization is influenced by the cooling rate. As shown in Fig. 6A, very long lamellae (up to 5 mm) can be seen in S23 when cooled at 2 C/min, indicating that a lower cooling rate imparts polymer chains more time for orderly packing. The closer look in Fig. 6B shows the coexistence of long lamellae, short lamellae and lamellae stacks. This observation indicates that S23 crystallizes from a homogenous or a weakly segregated melt. In contrast, for S36 only individual or aggregated short lamellae in island-like domains can be seen. The size of these domains is similar to that of the quenched sample. The main difference is that the lamellae of the slow cooled sample are slightly longer and thicker than those of the quenched one. The fact that cooling rate only influences crystal perfection but cannot
distort mesophase segregation indicates that S36 has mesophase separation in the melt and preserves it after crystallization. 3.3. Small-angle X-ray scattering Scattering is a straightforward approach toward analyzing the nano-scale structure. Li et al. [20] have reported that the two polyolefin blocks of polydisperse di-block OBC in the melt had nearzero X-ray contrast, therefore no SAXS intensity above the background was observed in the melt of the polymers they studied. In our system, both S23 and S36 melts showed very weak scattering signals, which is insufficient for illustration of the melt morphology. Scattering intensity increases after crystallization. It can be seen that the scattering intensities for S23, either quenched or slowly cooled samples, decreases with increasing q. No obvious long spacing peak can be identified even in Lorentz corrected plot
Fig. 5. AFM phase images (A, B) and corresponding bright field TEM images (C, D) of S23 (A, C) and S36 (B, D) quenched from the melt.
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Fig. 6. AFM phase images of slowly cooled (2 C/min) OBCs: (A) and (B): S23; (C): S36.
(Fig. 7A). A shoulder peak in the scattering curve is observed for the S36 solid samples, which is more distinct after Lorentz correction (Fig. 7B). The corresponding length (L) can be calculated from the peak position via the following relation:
L ¼
2p q
The quenched and slowly cooled S36 exhibit peak values at a q of about 0.39 nm1 and 0.30 nm1. The corresponding lengths are Lquen z 16 nm and Lslow z 21 nm for quenched and slow cooled samples, respectively. The calculated sizes are within the typical range of long spacing of semi-crystalline polymers and are consistent with our AFM and TEM observations. It might be interesting to discuss the scattering of multi-block OBC solid samples i.e. to discuss the reason for the appearance and disappearance of long spacing peaks. Zuo et al. [29] observed that a multi-block OBC with DC8 ¼ 16.3 mol% had stronger scattering intensity than that of EOC in SAXS. No long spacing peak was found in the Ieq curve of OBC they studied, however a long spacing peak was found after Lorentz correction. Their interpretation was the formation of uncorrelated grain-like amorphous local domains during crystallization, which had lower electron density than the bulk amorphous phase. As to our system, we propose a new interpretation of this phenomenon based on the different phase segregation behavior of the OBCs under investigation. For S36, crystallization is restricted in the segregated domains of crystallizable hard segments. The lamellae are densely packed thus correlate with each other, therefore a long spacing peak can be observed in SAXS. For S23, the crystalline lamellae uniformly disperse in the amorphous matrix. Because the crystallinity of S23
is less than 10%, the inter-lamellar distances have a very broad distribution. In other words, the correlation among the lamellae is lost. Therefore, the intensity profile of S23 decreases monotonically with the scattering vector, displaying particle scattering characteristics. The reason for the different morphologies can be attributed to the different mesophase segregation strength of the two samples which is determined by cN. It is noted that the block length and block length distribution have certain influence on the phase behavior of block copolymers. However, Quantitative discussion of them demands new generation of characterization technique, which unfortunately is currently not available. Nevertheless, both higher DC8 and longer blocks cause higher segregation strength for S36 than for S23. Therefore, mesophase separation is largely suppressed during crystallization in S23, while mesophase separation dominates in S36. 3.4. Crystallization and melting behavior To investigate the crystallization and melting behavior of the OBCs, DSC measurements were performed. Fig. 8 shows relatively sharp crystallization and melting peaks, which is typical for OBC compared with EOC of similar density [27,31]. No additional exothermic or endothermic peaks are identified, indicating that the soft segments are totally amorphous for the high octene content in soft blocks of the two OBCs under investigation. The melting temperature of S23 is about 3 C higher than that of S36, because the octene content in hard segments of S23 is slightly lower than that of S36. The crystallization peak of S36 is asymmetric with a long “tail” at the high temperature side. Table 3 summarizes the
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Fig. 7. SAXS curves of S23 and S36 for different conditions (A); Lorentz corrected SAXS curves (B).
DSC results. We define the temperature at which the endothermic curve diverges from the baseline as Tc, onset, and the peak temperatures of crystallization and melting as Tc, peak, Tm, peak, respectively. Tc, onset of S23 is slightly higher (w3 C), while the Tc, peak is much higher (w13 C) than that of S36. The supercooling calculated by Tm, peak Tc, onset for S23 is similar to that of S36. However, the supercooling for S23 is w9 C lower than that of S36 if the supercooling is defined as Tm, peak Tc, peak. The long crystallization tail at higher temperatures manifested by the large difference between crystallization onset and peak temperature might be caused by the conserved mesophase separation of the hard blocks during crystallization, which results in confinement and low mobility of polymer chains. At such
conditions long and linear segments between branches crystallize first at higher temperatures. Currently a systematic study on the relation between DC8 and crystallization kinetics is in progress. 3.5. Structure evolution during deformation Fig. 9 shows the 2-D WAXS patterns of S23 and S36 at different stretching ratios at room temperature. At non-stretched state, a clear amorphous halo at low diffraction angle and two distinct rings can be observed for both S23 and S36. The two rings can be indexed to be the (110) and (200) reflections from the orthorhombic PE crystals. All reflections are isotropic, indicating that the crystalline phase is unoriented in the initial samples. When strain is applied to S23, the WAXS patterns become anisotropic. At l ¼ 4, the (110) and (200) reflections turn to arcs indicating the preferential orientation of the crystals. The (110) and (200) arcs become shorter and gradually reduce to spots at the meridian with increasing the stretching ratio, indicating that the c-axis of orthorhombic crystals gradually orients parallel to the stretching direction. At l ¼ 7, an additional meridian diffraction at a Bragg angle between amorphous halo and (110) reflection is seen, the intensity of which becomes stronger with increasing the stretching ratio. This can be the (010) reflection from the monoclinic crystals. At high stretching
Table 3 DSC data of S23 and S36.
Fig. 8. DSC thermographs of S23 and S36 obtained with cooling and heating rates of 10 C/min.
Sample code
Tc,
S23 S36
99.6 C 96.6 C
onset
Tc,
peak
88.4 C 75.2 C
Tm,
peak
116.5 C 112.0 C
Tm,
peak
28.1 C 36.8 C
Tc,
peak
Tm,
peak
16.9 C 15.4 C
Tc,
onset
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Fig. 9. 2-D WAXS patterns of S23 and S36 at different stretching ratios. The stretching direction is horizontal.
ratios, the WAXS pattern looks like fiber X-ray scattering pattern. This can be interpreted by the formation of micro fibrils [29]. The deformation feature of S23 that strain can induce orientation and transformation of crystals fits the general understanding and agrees well with the previous reports [29,30]. However, we are surprised to find that the WAXS patterns of S36 almost remain isotropic, suggesting that the crystalline phase has very low orientation upon stretching. The data are further analyzed in the following part. Integrated 1-D WAXS profiles were obtained by 360 scan (Fig. 10). For S23, the heights of both (110) and (200) reflections
Fig. 10. Integrated WAXS profiles of S23 and S36 at different stretching ratio.
relative to that of the amorphous peak decreased with increasing the stretching ratio. When l ¼ 12, it is difficult to separate the (110) and (200) reflections from the broad overlap of the amorphous peak and the potential monoclinic (010) reflection. Therefore, the fraction of the orthorhombic phase cannot be estimated using the 1-D peak fitting method at this stretching ratio. For S36, however, the heights of both the (110) and (200) reflections relative to that of the amorphous peak almost keep unchanged. To quantitatively analyze the structural change, the fractions of the orthorhombic crystal in S23 and S36 were calculated using the 1-D peak fitting method. We assume that the stretched sample has a cylindrical symmetry around the stretching axis, where the 2-D WAXS patterns can be used to calculate the fraction of each phase. The 1-D profile with a 2 theta range of 8 e22 was fitted by three Gaussian þ Lorentz type peaks: the amorphous phase, the orthorhombic (110) and (200) peaks, respectively. Since the amorphous scattering of S23 is superposed with the monoclinic (010) reflection and cannot be deconvoluted by 1-D peak fitting, we only estimate the area of orthorhombic (110) and (200) reflections, which are intense and can be easily separated from the broad reflection from 8 to 22 . Furthermore, the results are satisfactory with 3 peak fitting (except l ¼ 12 sample of S23). The fraction of orthorhombic crystal are calculated by the area ratio of orthorhombic (110) and (200) reflections and total area of diffraction, as illustrated in Fig. 11. Obvious differences between S23 and S36 can be observed. The fraction of orthorhombic crystal in S23 gradually decreases with increasing stretching ratio, while that in S36 varies very little with increasing stretching ratio.
Fig. 11. Fraction of the orthorhombic crystal in S23 and S36 at different stretching ratio.
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The full width at half maximum (FWHM) of the orthorhombic (110) reflection at different stretching ratios is shown in Fig. 12. The FWHM of S23 increases with increasing stretching ratio, while the FWHM of S36 keeps almost unchanged. According to the Scherrer equation, the crystal size is inversely proportional to FWHM. Therefore, the orthorhombic crystal size in S23 decreases and that in S36 almost keeps unchanged with increasing strain. The crystal size decrease in S23 can be explained by the crystal fragmentation during deformation. The orientation degree of orthorhombic crystal in S23 and S36 was calculated by Hermans orientation factor [32]:
f ¼
3 cos2 4 1 2
where 4 is the angle between the chain axis and the reference axis (stretching axis). hcos2 4i is defined as:
Zp=2 D E cos2 4 ¼
Ið4Þcos2 4 sin4d4
0
Zp=2
Ið4Þsin4d4
Fig. 13. Hermans orientation factor of crystals for S23 and S36 as a function of stretching ratio.
preferentially orients parallel to the tensile direction. Finally, a highly oriented structure with micro fibrils forms. The deformation process of S36 distinctly differs from that of S23. With increasing strain, the amorphous matrix deforms accordingly, however, the inner structure of the segregated crystalline region largely keeps intact.
0
where I(4) is the scattering intensity along the angle 4. The range of Hermans orientation factor is in the range of 0.5 to 1. When f is 0.5, it means that all the chains align perpendicularly to the reference axis. When f ¼ 1, all the chains are perfectly parallel to the reference axis. In our case, the c-axis orientation factor (fc) could not be measured directly due to the absence of (00l) plane, therefore the fc is obtained indirectly from the (110) and (200) reflections using the following expression [29]
D E D E D E cos2 4 ¼ 1 1:435 cos2 4110 0:565 cos2 4200 The calculated fc values in S23 and S36 are illustrated in Fig. 13. The fc of S23 increases asymptotically toward the value of unity with increasing stretching ratio, indicating that the chains gradually align along the stretching direction. For S36, however, the fc increases slowly to a small value of about 0.14 at l ¼ 12. This indicates that the orientation of crystals in S36 is very weak even at a very high stretching ratio. Fig. 14 is a schematic illustration showing the different deformation process of S23 and S36. The initial structure of S23 is featured as nearly homogeneous distributed lamellae, while the crystalline lamellae in S36 segregate into domains. Upon stretching, the lamellae in S23 slip with each other and fragment into blocks. With increasing the strain, the c-axis of crystals of S23 gradually
Fig. 12. FWHM of orthorhombic (110) reflection at different stretching ratio.
3.6. Factors influencing the elastomeric properties In the slip-link treatment of Hiltner and coworkers [27,28], crystalline phase acted as junctions, which contributed to the elastomeric response by increasing the sliplink density. Notwithstanding the limited amount, the crystalline phase, with much higher modulus than amorphous phase, also contributes to the mechanical properties itself as filler. The crystalline lamellae in S23 are dispersed in the amorphous matrix, while the lamellae in S36 are confined in isolated meso-domains. It is well known that fibrils/ lamellar fillers are more effective than particle fillers in reinforcing modulus. This is probably an important reason for the different modulus of S23 and S36. During tensile deformation, the crystal lamellae of S23 undergo rotating, tilting, slipping and finally breaking off and transforming to other phases. This indicates that crystalline phase has contribution to the tensile force in the whole strain range. However, the “hard” crystalline domains in S36 do not deform cooperatively, therefore the crystalline phase in S36 has less important contribution to the mechanical response. It has long been realized that microphase separation is important for the mechanical response of poly(styrene-b-diene-b-styrene) (SBS) thermoplastic elastomers [32]. Meanwhile, the success of ethylene random copolymer elastomers indicates that crystallization is also an important factor for elastomeric properties. It is natural to ask which factor governs the elastomeric properties in a system that has both mesophase separation and crystallization. Seguela and Prudhomme [33] reported that melt crystallized hydrogenated poly(butadiene-b-isoprene-b-butadiene) tri-block polymer (HBIB) exhibited much higher tensile strength than solution crystallized ones. The meltecrystallized samples undergo incompatibilityinduced microphase separation, whereas the solution-crystallized samples segregated due to crystallization. Koo et al. [13,14] got similar results. At a first glance, our results are contradictory to theirs. However, several important differences must be noted. Firstly, in their cases, the microphase separated morphology is 1-D [33] (cylindrical) or 2-D [13,14] (lamellar), while in our case it is 0-D (spherical). Cylinders or lamellae with high length in one or two direction are intrinsically easier to carry load than isolated spheres. To date, the relationship between micro/mesophase separation, crystallization and mechanical properties of block copolymers is still
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Fig. 14. Schematic illustration of the deformation mechanism of S23 and S36.
not fully understood. Moreover, there are other factors that also influence the mechanical properties. For instance, Park et al. [24] reported that the molecular weight between entanglements (Me) in OBC increased with increasing the commoner content, which could be another factor for the lower modulus of S36. Sarva et al. [34] reported that microstructure of elastomer had influence on the rate dependent stressestrain behavior. It has also been reported that the number of arms in star shaped block copolymers has strong influence on the mechanical properties [35,36]. Despite the complexity in factors governing the mechanical properties, we stress here based on our results that the cooperative deformation of crystalline phase makes an important contribution to the mechanical properties of crystalline elastomers. 4. Conclusions We investigated the mesophase phase separation and crystallization of multi-block OBCs with different chain composition. Crystallization within mesophase separated domains led to a densely packed lamellar structure with relatively stronger correlation between lamellae, showing a long spacing peak in SAXS. However, crystallization in OBC with lower segregation strength suppressed the mesophase separation. The uniform dispersed crystalline lamellae had lower inter-lamellar correlation, exhibiting a monotonically decreased intensity profile in SAXS. The two OBCs had distinct difference in mechanical behavior. One had a higher modulus and showed strain hardening while the other had a much lower modulus and almost no strain hardening. Tensile WAXS measurements showed different micro-structure transformation process. The orthorhombic crystal exhibited orientation, fragmentation and transformation to monoclinic crystal under stretching in one OBC; while the orthorhombic crystal in the other almost kept intact and had very low orientation under stretching. The filler effect of crystalline phase differed between the two OBCs due to the different morphology. The cooperative deformation of crystalline phase was proposed to make important contribution to the mechanical response of crystalline elastomers. Acknowledgments Financial support from National Natural Science Foundation of China is gratefully acknowledged (50903089, 21074141, 50925313).
Joachim Loos thanks CAS for his Visiting Professorship for Senior International Scientists (2009J2-28). The Dow Chemical Company is acknowledged for partial financial support. The SSRF is acknowledged for providing the beam time. Prof. Chenyang Liu (ICCAS) is thanked for constructive discussions. References [1] Hamley IW. Adv Polym Sci 1999;148:113e37. [2] Hamley IW. Developments in block copolymer science and technology. Chichester: John Wiley & Sons Ltd.; 2004. [3] Zhu L, Cheng SZD, Calhoun BH, Ge Q, Quirk RP, Thomas EL, et al. J Am Chem Soc 2000;122(25):5957e67. [4] Quiram DJ, Register RA, Marchand GR, Adamson DH. Macromolecules 1998; 31(15):4891e8. [5] Loo YL, Register RA, Adamson DH. Macromolecules 2000;33(22):8361e6. [6] Rangarajan P, Register RA, Adamson DH, Fetters LJ, Bras W, Naylor S, et al. Macromolecules 1995;28(5):1422e8. [7] Ryan AJ, Hamley IW, Bras W, Bates FS. Macromolecules 1995;28(11): 3860e8. [8] Quiram DJ, Register RA, Marchand GR. Macromolecules 1997;30(16): 4551e8. [9] Quiram DJ, Register RA, Marchand GR, Ryan AJ. Macromolecules 1997;30(26): 8338e43. [10] Loo YL, Register RA, Ryan AJ. Phys Rev Lett 2000;84(18):4120e3. [11] Loo YL, Register RA, Ryan AJ, Dee GT. Macromolecules 2001;34(26):8968e77. [12] Loo YL, Register RA, Ryan AJ. Macromolecules 2002;35(6):2365e74. [13] Koo C, Wu L, Lim L, Mahanthappa M, Hillmyer M, Bates F. Macromolecules 2005;38(24):6090e8. [14] Koo C, Hillmyer M, Bates F. Macromolecules 2006;39(2):667e77. [15] Hamley IW. The physics of block copolymers. London: Oxford University Press; 1998. [16] Arriola DJ, Carnahan EM, Hustad PD, Kuhlman RL, Wenzel TT. Science 2006; 312(5774):714e9. [17] Graessley WW, Krishnamoorti R, Balsara NP, Butera RJ, Fetters LJ, Lohse DJ, et al. Macromolecules 1994;27(14):3896e901. [18] Reichart GC, Graessley WM, Register RA, Lohse DJ. Macromolecules 1998; 31(22):7886e94. [19] Hustad PD, Marchand GR, Garcia-Meitin EI, Roberts PL, Weinhold JD. Macromolecules 2009;42(11):3788e94. [20] Li S, Register RA, Landes BG, Hustad PD, Weinhold JD. Macromolecules 2010; 43(10):4761e70. [21] Wang HP, Khariwala DU, Cheung W, Chum SP, Hiltner A, Baer E. Macromolecules 2007;40(8):2852e62. [22] Khariwala DU, Taha A, Chum SP, Hiltner A, Baer E. Polymer 2008;49(5): 1365e75. [23] Kamdar AR, Wang HP, Khariwala DU, Taha A, Hiltner A, Baer E. J Polym Sci Part B Polym Phys 2009;47(16):1554e72. [24] Park HE, Dealy JM, Marchand GR, Wang J, Li S, Register RA. Macromolecules 2010;43(16):6789e99. [25] Jin J, Du J, Xia Q, Liang Y, Han CC. Macromolecules 2010;43(24):10554e9. [26] Treloar LRG. The physics of rubber elasticity. London: Oxford University Press; 1967.
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G. Liu et al. / Polymer 52 (2011) 5221e5230
[27] Bensason S, Stepanov EV, Chum S, Hiltner A, Baer E. Macromolecules 1997; 30(8):2436e44. [28] Wang HP, Chum SP, Hiltner A, Baer E. J Appl Polym Sci 2009;113(5):3236e44. [29] Zuo F, Burger C, Chen X, Mao Y, Hsiao BS, Chen H, et al. Macromolecules 2010; 43(4):1922e9. [30] Wang HP, Chum SP, Hiltner A, Baer E. J Polym Sci Part B Polym Phys 2009; 47(13):1313e30. [31] Wilchinsky ZW. J Appl Phys 1960;31(11):1969e72.
[32] Holden G, Legge NR. Thermoplastic elastomers. 2nd ed. New York: Hanser Publishers; 1996. [33] Seguela R, Prudhomme J. Polymer 1989;30(8):1446e55. [34] Sarva SS, Hsieh AJ. Polymer 2009;50(13):3007e15. [35] Juhari A, Mosnácek J, Yoon JA, Nese A, Koynov K, Kowalewski T, et al. Polymer 2010;51(21):4806e13. [36] Mosnácek J, Yoon JA, Juhari A, Koynov K, Matyjaszewski K. Polymer 2009; 50(9):2087e94.
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Effect of mesophase separation and crystallization on the elastomeric behavior of olefin multi-block copolymers Guoming Liua, Yu Guana, Tao Wena, Xiwei Wanga, Xiuqin Zhanga, Dujin Wanga, *, Xiuhong Lib, Joachim Loosc, Hongyu Chend, Kim Waltone, Gary Marchande a
Beijing National Laboratory for Molecular Sciences, CAS Key Laboratory of Engineering Plastics, Institute of Chemistry, Chinese Academy of Sciences, Beijing 100190, China Shanghai Synchrotron Radiation Facility, Shanghai Institute of Applied Physics, Chinese Academy of Sciences, Shanghai 201204, China School of Physics and Astronomy, University of Glasgow, Glasgow G12 8QQ, Scotland, UK d The Dow Chemical (China) Company Limited, Shanghai 201203, China e The Dow Chemical Company, Freeport, TX 77541, USA b c
a r t i c l e i n f o
a b s t r a c t
Article history: Received 30 July 2011 Received in revised form 6 September 2011 Accepted 8 September 2011 Available online 14 September 2011
The mesophase separation and crystallization as well as the elastomeric properties of olefin multi-block copolymers (OBCs) are studied. The solid state morphologies of the OBCs are determined by the competition between mesophase separation and crystallization of hard blocks. The OBC with lower DC8 (octene content difference between soft and hard blocks) displays a spherulitic superstructure, while the OBC with higher DC8 exhibits mesophase separation. The two OBCs show very different elastomeric properties. Wide-angle X-ray scattering shows that the two OBCs have different deformation mechanisms. It is proposed that the cooperative deformation of crystalline phase makes an important contribution to the elastomeric properties. Ó 2011 Elsevier Ltd. All rights reserved.
Keywords: Elastomer Deformation Morphology
1. Introduction Due to the diversity in morphologies and properties, block copolymers have attracted widespread research interest. Below the orderedisorder transition (ODT) temperature, the melt of a block copolymer can self-assemble into spheres, rods, lamellae, cylinders or more complex structures. For a di-block copolymer, the melt phase behavior depends on two parameters: f, the volume fraction of one component and cN, in which c is the FloryeHuggins interaction parameter and N is the degree of polymerization. Di-block copolymers with both blocks being non-crystalline preserve the microphase separated morphology of the melt during cooling. For amorphous-crystalline di-block copolymers, on the other hand, two competing mechanisms drive solidification, i.e., microphase separation and crystallization. Whether the microphase separated structure is preserved or destroyed during cooling depends on the state of the amorphous blocks upon crystallizing and the segregation strength [1,2]. When the non-crystalline blocks are glassy during crystallization, the overall melt structure is retained.
* Corresponding author. Tel./fax: þ86 10 82618533. E-mail address: [email protected] (D. Wang). 0032-3861/$ e see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.polymer.2011.09.009
Crystallization occurs within the nano-scale domains prescribed by microphase separation [3e5]. However, when the amorphous blocks are rubbery during crystallization, microphase separation competes with crystallization, and the ultimate morphology depends on the segregation strength. Weakly segregated block copolymers undergo enormous structural rearrangement during crystallization, resulting in the formation of lamellar crystallites and even spherulitic superstructures [6e8]; while strong segregated polymers can preserve the morphology formed in the melt state [8e11]. The different crystallization behavior of block copolymers has been classified by Loo et al. [12] as “break out”, “templated” or “confined”. Crystallization from a homogeneous melt for block copolymers containing crystalline blocks and amorphous blocks can also induce microphase separation by the expulsion of amorphous blocks [13,14]. More information about structure, properties and theory of well-defined di- or tri-block copolymers can be found elsewhere [2,15]. Recently, the Dow Chemical Company developed olefin block copolymers (OBCs, INFUSEÔ) based on a chain shuttling catalyst technology [16]. The block copolymers synthesized by such technology consist of ethyleneeoctene segments with low octene concentrations in the hard or crystalline blocks and high octene concentration in the soft or amorphous blocks. This new type of
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2. Experimental
Fig. 1. Sketch showing the molecular architecture of OBC.
olefin block copolymer has a statistical multi-block architecture with a distribution in block lengths and the number of blocks per chain. Compared to random ethyleneeoctene copolymers (EOC), the block architecture imparts a substantially higher melting temperature, a higher crystallization temperature while maintaining a low glass transition temperature. As mentioned above, one of the basic questions in the area of block copolymer research is how the morphology in the melt and potential reorganization during crystallization influence morphology formation and properties of the copolymer. The miscibility (cN) of hard segments and soft segments in OBC depends on molecular weight and the DC8 (the octene content difference between soft and hard segments) [17,18]. Hustad et al. [19] found that polydisperse di-block OBC with DC8 z 35 mol% formed lamellar structures in the solid state with interesting photonic properties. Li et al. [20] reported that nearsymmetric polydisperse di-block OBCs with DC8 exceeding 25 mol% and number averaged molecular weights of 69e93 kg/mol selfassembled into well-defined lamellar structures. After crystallization, the lamellae-type structure was preserved. Crystallites in such lamellae were isotropic oriented and the crystallization kinetics was consistent with a spread growth behavior. The new mode of crystallization within heterogeneous block copolymers was termed as “pass through”. Multi-block OBCs have been assumed to have homogeneous melts in relatively earlier studies based on optical microscopy observations of spherulitic morphology [21,22]. However, phase separation was observed for an OBC with high DC8 of 21 mol% [23], though the melt morphology was still unknown, i.e. whether the solid state phase separated morphology is a replica of the melt structure or is formed during the crystallization process. Recently, based on rheological analysis a first attempt has been made to better understand the melt structure of multi-block OBCs [24]. Large composition fluctuations and mesophase separation transition at temperatures well above the melting point were found in multi-block OBC with DC8 higher than 16 mol%. Jin et al. [25] found the morphology of multi-OBC resulted from the competition between crystallization and mesophase separation. It is well known that structure and mechanical correlation is very important in the application of polymeric materials. In this study, our objective is to understand how the solid state morphology (mesophase separation and crystallization) influences the mechanical properties of OBCs. We explored two OBCs with similar molecular architecture and crystallinity. The solid state morphology was studied by microscopy and small-angle X-ray scattering. Two dimensional (2-D) wide-angle X-ray scattering measurement was carried out on stretched OBCs.
Two multi-block OBCs synthesized via a novel chain shuttling polymerization technology were supplied by the Dow Chemical Company. Fig. 1 shows a sketch of the general molecular structure of OBC with low octene concentration in the hard and high octene concentration in the soft segments. Molecular weight, molecular weight distribution, density and viscosity of the two samples are similar. The two OBCs have comparable low octene content in their hard segments but significantly different octene content in their soft segments. The amount of chain shuttling agent (CSA) illustrated as the Zn content in the samples was determined by X-ray florescence. Because the soft segment octene content in the two OBCs is 22.6 mol% and 35.7 mol%, the OBCs were designated as S23 and S36, respectively. Relatively lower amount of CSA was used in S36, which would principally result in longer segments. The molecular information of the OBCs is shown in Table 1. Note that the crystallinity of S23 is slightly higher than that of S36, because the hard segment octene content in S23 is slightly lower than that in S36. Samples with thickness of 1 mm for mechanical testing and X-ray scattering experiments and ca. 20 mm for optical microscopy were prepared by compression molding at 190 C. Thin films with about 100 nm thickness were made by spin-coating xylene solution at 110 C on carbon covered mica at 80 C. Subsequently, the thin film samples were heated to 190 C for 5 min and then either quenched into liquid nitrogen or slowly cooled to room temperature at a rate of 2 C/min. The thermally treated thin films were cut into pieces for further atomic force microscopy (AFM) and transmission electron microscopy (TEM) investigation. Uniaxial tensile tests were conducted on 1 mm thick ISO-37 Type 2 specimens cut from compression-molded plaques. Specimens were stretched on an Instron 3365 universal mechanical testing machine at room temperature with a crosshead speed of 300 mm/min. Engineering stress was defined conventionally as force per initial cross-sectional area. Polarized light optical micrographs were obtained with a BX51 Olympus optical microscope equipped with a Cannon (50D) camera. The free surface of the thin films was imaged in air with a commercial atomic force microscope (AFM) Nanoscope IIIa from Digital Instruments operated in tapping mode. For TEM observations, specimens were prepared by floating the film samples on water and transferring them onto copper grids. A JEM-2200S TEM was used with an operation voltage of 200 kV. X-ray measurements were carried out at the beamline BL16B1 in the Shanghai Synchrotron Radiation Facility (SSRF). The wavelength of the radiation source was l ¼ 1.24 Å. Scattering patterns were collected by a MAR CCD (MAR-USA) detector with a resolution of 2048 2048 pixels (pixel size: 79 79 mm2). Image acquisition time was 200 s. The sample to detector distance was 3117.4 mm for SAXS and 193.2 mm for WAXS. WAXS patterns for OBC samples with different drawing ratio were collected. The samples were stretched to a specific strain and then fastened on a home-made holder to keep the strain. All the X-ray images were corrected for background scattering, air scattering and beam fluctuations.
Table 1 Molecular characteristic of OBCs. Sample code
Density (g/cm3)
Mn (kg/mol)
Mw (kg/mol)
Octene content in soft segmenta (mol%)
Octene content in hard segmenta (mol%)
Zn content (ppm)
Hard segment (wt%)
Xc, DHb (wt%)
S23 S36
0.867 0.863
76 74
180 249
22.6 35.7
1.13 2.06
177 104
16 18
10 7.8
a b
Determined from 13C NMR. Determined by enthalpy of fusion.
G. Liu et al. / Polymer 52 (2011) 5221e5230
The melting and crystallization behavior of the OBCs was examined with a PerkineElmer differential scanning calorimeter (DSC 7). The instrument was calibrated with indium before measurement. Temperature scans were performed in the temperature range from 0 to 200 C at a heating/cooling rate of 10 C/min under nitrogen atmosphere. The typical sample weight was about 2e4 mg. 3. Results and discussion 3.1. Mechanical behavior of the OBCs Engineering stressestrain curves of quenched S23 and S36 at room temperature are plotted in Fig. 2. The two OBCs showed typical elastomeric behavior, specifically no yielding and uniform deformation to high strains. They had high elongations, and did not fracture up to an extension ratio of 13. Obvious differences can be found between the two polymers. S23 exhibited strain hardening at high strains while S36 had almost no strain hardening. The initial elastic modulus measured by the slope of the linear part of the engineering stressestrain curves is 6.1 0.4 MPa and 0.95 0.07 MPa for S23 and S36 (Table 2), respectively. In addition, the mechanical data is expressed in terms of modulus G usually used for elastomers [26]:
2 G ¼ s= l 1=l where l is the extension ratio and s is the engineering stress. The modulus gradually decreases at low strains for the two OBCs. In the whole tensile range, the modulus of S23 is much higher than that of S36. The modulus of S23 has an upswing at high strains, while the modulus of S36 has a continuous decrease (Fig. 3). It was also found that for a same material, independent of quenching or slow cooling conditions the general mechanical behavior was similar. Hiltner and coworkers [27,28] found that the slip-link theory well described the stressestrain curves of crystalline thermoplastic elastomers, including ethyleneeoctene random copolymers (EOC) and OBC. The initial modulus and the plateau modulus were primarily determined by the sliplink density, which was shown to scale with crystallinity, while the tightly topological entanglements at high strain served as crosslinks. In our experiments, the crystallinity of S23 is slightly higher than that of S36. Dividing the initial modulus by crystallinity, the normalized initial modulus is 62 4 MPa for S23 and 12 1 MPa for S36 (Table 2), respectively,
Fig. 2. Stressestrain curves of S23 and S36 at room temperature.
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Table 2 Mechanical data of S23 and S36. Sample code
Tangent modulus G (MPa)
G/Xc (MPa)
S23 S36
6.1 0.4 0.95 0.07
62 4 12 1
indicating that the normalized initial modulus of S23 is 5 times larger than that of S36. The plateau modulus of S23 is also much larger than that of S36. Obviously, this result cannot be explained by the previous scaling arguments. Molecular weight seems to play a minor role, as S36 with higher Mw does not impart higher strength compared to S23 with lower Mw.
3.2. Morphology observation To understand the different elastomeric properties of the two multi-block OBCs, we carried out morphological study via optical microscopy, AFM and TEM. For high cooling rates (>10 C/min), no spherulitic structure can be seen in either S23 or S36. When the cooling rate is lower, however, S23 has obvious spherulites (Fig. 4). On the contrary, no spherulitic superstructure can be observed in S36 even at a very low cooling rate of 1 C/min. To further reveal the nanostructure of the OBCs, we have prepared and investigated spin coated thin films quenched from the melt. Fig. 5 shows AFM phase images and bright field TEM images of the OBC samples. In the present case, AFM phase images display higher contrast and resolution than topography images because of the large difference in modulus between the crystalline and amorphous phase. As shown in Fig. 5A, crystalline lamellae and lamellar stacks can be seen in S23, which are brighter due to their relatively higher modulus. The lamellae have a thickness less than 10 nm and a length of 200e300 nm, and they are homogenously distributed in the amorphous matrix. The lamellar stacks consist of a few lamellar crystals, which are confirmed by the corresponding TEM observations shown in Fig. 5C. For sample S36, the AFM image in Fig. 5B shows island-like bright aggregates which consist of several very short lamellar crystals that embed in the “darker” amorphous matrix (Fig. 5D is the corresponding TEM image). Based on the above observations, we notice that the two OBCs have clearly different structural characteristics: relatively long but uncorrelated lamellae for S23, and densely packed lamellae in isolated regions for S36. In other words, S36 exhibits “mesophase separation” (following the term as introduced in literatures [20,24,25]), while crystallization dominated structure with no clear mesophase separation occurs in S23.
Fig. 3. Modulus G vs. inverse extension ratio of S23 and S36 at room temperature.
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Fig. 4. Polarized light optical microscopy images of S23 (A) and S36 (B) cooled at 2 C/min.
The final morphology of a block copolymer containing crystallizable segments is a result of both the melt structure and the crystallization process. An indirect approach to detect the melt morphology is to check whether the morphology after crystallization is influenced by the cooling rate. As shown in Fig. 6A, very long lamellae (up to 5 mm) can be seen in S23 when cooled at 2 C/min, indicating that a lower cooling rate imparts polymer chains more time for orderly packing. The closer look in Fig. 6B shows the coexistence of long lamellae, short lamellae and lamellae stacks. This observation indicates that S23 crystallizes from a homogenous or a weakly segregated melt. In contrast, for S36 only individual or aggregated short lamellae in island-like domains can be seen. The size of these domains is similar to that of the quenched sample. The main difference is that the lamellae of the slow cooled sample are slightly longer and thicker than those of the quenched one. The fact that cooling rate only influences crystal perfection but cannot
distort mesophase segregation indicates that S36 has mesophase separation in the melt and preserves it after crystallization. 3.3. Small-angle X-ray scattering Scattering is a straightforward approach toward analyzing the nano-scale structure. Li et al. [20] have reported that the two polyolefin blocks of polydisperse di-block OBC in the melt had nearzero X-ray contrast, therefore no SAXS intensity above the background was observed in the melt of the polymers they studied. In our system, both S23 and S36 melts showed very weak scattering signals, which is insufficient for illustration of the melt morphology. Scattering intensity increases after crystallization. It can be seen that the scattering intensities for S23, either quenched or slowly cooled samples, decreases with increasing q. No obvious long spacing peak can be identified even in Lorentz corrected plot
Fig. 5. AFM phase images (A, B) and corresponding bright field TEM images (C, D) of S23 (A, C) and S36 (B, D) quenched from the melt.
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Fig. 6. AFM phase images of slowly cooled (2 C/min) OBCs: (A) and (B): S23; (C): S36.
(Fig. 7A). A shoulder peak in the scattering curve is observed for the S36 solid samples, which is more distinct after Lorentz correction (Fig. 7B). The corresponding length (L) can be calculated from the peak position via the following relation:
L ¼
2p q
The quenched and slowly cooled S36 exhibit peak values at a q of about 0.39 nm1 and 0.30 nm1. The corresponding lengths are Lquen z 16 nm and Lslow z 21 nm for quenched and slow cooled samples, respectively. The calculated sizes are within the typical range of long spacing of semi-crystalline polymers and are consistent with our AFM and TEM observations. It might be interesting to discuss the scattering of multi-block OBC solid samples i.e. to discuss the reason for the appearance and disappearance of long spacing peaks. Zuo et al. [29] observed that a multi-block OBC with DC8 ¼ 16.3 mol% had stronger scattering intensity than that of EOC in SAXS. No long spacing peak was found in the Ieq curve of OBC they studied, however a long spacing peak was found after Lorentz correction. Their interpretation was the formation of uncorrelated grain-like amorphous local domains during crystallization, which had lower electron density than the bulk amorphous phase. As to our system, we propose a new interpretation of this phenomenon based on the different phase segregation behavior of the OBCs under investigation. For S36, crystallization is restricted in the segregated domains of crystallizable hard segments. The lamellae are densely packed thus correlate with each other, therefore a long spacing peak can be observed in SAXS. For S23, the crystalline lamellae uniformly disperse in the amorphous matrix. Because the crystallinity of S23
is less than 10%, the inter-lamellar distances have a very broad distribution. In other words, the correlation among the lamellae is lost. Therefore, the intensity profile of S23 decreases monotonically with the scattering vector, displaying particle scattering characteristics. The reason for the different morphologies can be attributed to the different mesophase segregation strength of the two samples which is determined by cN. It is noted that the block length and block length distribution have certain influence on the phase behavior of block copolymers. However, Quantitative discussion of them demands new generation of characterization technique, which unfortunately is currently not available. Nevertheless, both higher DC8 and longer blocks cause higher segregation strength for S36 than for S23. Therefore, mesophase separation is largely suppressed during crystallization in S23, while mesophase separation dominates in S36. 3.4. Crystallization and melting behavior To investigate the crystallization and melting behavior of the OBCs, DSC measurements were performed. Fig. 8 shows relatively sharp crystallization and melting peaks, which is typical for OBC compared with EOC of similar density [27,31]. No additional exothermic or endothermic peaks are identified, indicating that the soft segments are totally amorphous for the high octene content in soft blocks of the two OBCs under investigation. The melting temperature of S23 is about 3 C higher than that of S36, because the octene content in hard segments of S23 is slightly lower than that of S36. The crystallization peak of S36 is asymmetric with a long “tail” at the high temperature side. Table 3 summarizes the
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Fig. 7. SAXS curves of S23 and S36 for different conditions (A); Lorentz corrected SAXS curves (B).
DSC results. We define the temperature at which the endothermic curve diverges from the baseline as Tc, onset, and the peak temperatures of crystallization and melting as Tc, peak, Tm, peak, respectively. Tc, onset of S23 is slightly higher (w3 C), while the Tc, peak is much higher (w13 C) than that of S36. The supercooling calculated by Tm, peak Tc, onset for S23 is similar to that of S36. However, the supercooling for S23 is w9 C lower than that of S36 if the supercooling is defined as Tm, peak Tc, peak. The long crystallization tail at higher temperatures manifested by the large difference between crystallization onset and peak temperature might be caused by the conserved mesophase separation of the hard blocks during crystallization, which results in confinement and low mobility of polymer chains. At such
conditions long and linear segments between branches crystallize first at higher temperatures. Currently a systematic study on the relation between DC8 and crystallization kinetics is in progress. 3.5. Structure evolution during deformation Fig. 9 shows the 2-D WAXS patterns of S23 and S36 at different stretching ratios at room temperature. At non-stretched state, a clear amorphous halo at low diffraction angle and two distinct rings can be observed for both S23 and S36. The two rings can be indexed to be the (110) and (200) reflections from the orthorhombic PE crystals. All reflections are isotropic, indicating that the crystalline phase is unoriented in the initial samples. When strain is applied to S23, the WAXS patterns become anisotropic. At l ¼ 4, the (110) and (200) reflections turn to arcs indicating the preferential orientation of the crystals. The (110) and (200) arcs become shorter and gradually reduce to spots at the meridian with increasing the stretching ratio, indicating that the c-axis of orthorhombic crystals gradually orients parallel to the stretching direction. At l ¼ 7, an additional meridian diffraction at a Bragg angle between amorphous halo and (110) reflection is seen, the intensity of which becomes stronger with increasing the stretching ratio. This can be the (010) reflection from the monoclinic crystals. At high stretching
Table 3 DSC data of S23 and S36.
Fig. 8. DSC thermographs of S23 and S36 obtained with cooling and heating rates of 10 C/min.
Sample code
Tc,
S23 S36
99.6 C 96.6 C
onset
Tc,
peak
88.4 C 75.2 C
Tm,
peak
116.5 C 112.0 C
Tm,
peak
28.1 C 36.8 C
Tc,
peak
Tm,
peak
16.9 C 15.4 C
Tc,
onset
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Fig. 9. 2-D WAXS patterns of S23 and S36 at different stretching ratios. The stretching direction is horizontal.
ratios, the WAXS pattern looks like fiber X-ray scattering pattern. This can be interpreted by the formation of micro fibrils [29]. The deformation feature of S23 that strain can induce orientation and transformation of crystals fits the general understanding and agrees well with the previous reports [29,30]. However, we are surprised to find that the WAXS patterns of S36 almost remain isotropic, suggesting that the crystalline phase has very low orientation upon stretching. The data are further analyzed in the following part. Integrated 1-D WAXS profiles were obtained by 360 scan (Fig. 10). For S23, the heights of both (110) and (200) reflections
Fig. 10. Integrated WAXS profiles of S23 and S36 at different stretching ratio.
relative to that of the amorphous peak decreased with increasing the stretching ratio. When l ¼ 12, it is difficult to separate the (110) and (200) reflections from the broad overlap of the amorphous peak and the potential monoclinic (010) reflection. Therefore, the fraction of the orthorhombic phase cannot be estimated using the 1-D peak fitting method at this stretching ratio. For S36, however, the heights of both the (110) and (200) reflections relative to that of the amorphous peak almost keep unchanged. To quantitatively analyze the structural change, the fractions of the orthorhombic crystal in S23 and S36 were calculated using the 1-D peak fitting method. We assume that the stretched sample has a cylindrical symmetry around the stretching axis, where the 2-D WAXS patterns can be used to calculate the fraction of each phase. The 1-D profile with a 2 theta range of 8 e22 was fitted by three Gaussian þ Lorentz type peaks: the amorphous phase, the orthorhombic (110) and (200) peaks, respectively. Since the amorphous scattering of S23 is superposed with the monoclinic (010) reflection and cannot be deconvoluted by 1-D peak fitting, we only estimate the area of orthorhombic (110) and (200) reflections, which are intense and can be easily separated from the broad reflection from 8 to 22 . Furthermore, the results are satisfactory with 3 peak fitting (except l ¼ 12 sample of S23). The fraction of orthorhombic crystal are calculated by the area ratio of orthorhombic (110) and (200) reflections and total area of diffraction, as illustrated in Fig. 11. Obvious differences between S23 and S36 can be observed. The fraction of orthorhombic crystal in S23 gradually decreases with increasing stretching ratio, while that in S36 varies very little with increasing stretching ratio.
Fig. 11. Fraction of the orthorhombic crystal in S23 and S36 at different stretching ratio.
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The full width at half maximum (FWHM) of the orthorhombic (110) reflection at different stretching ratios is shown in Fig. 12. The FWHM of S23 increases with increasing stretching ratio, while the FWHM of S36 keeps almost unchanged. According to the Scherrer equation, the crystal size is inversely proportional to FWHM. Therefore, the orthorhombic crystal size in S23 decreases and that in S36 almost keeps unchanged with increasing strain. The crystal size decrease in S23 can be explained by the crystal fragmentation during deformation. The orientation degree of orthorhombic crystal in S23 and S36 was calculated by Hermans orientation factor [32]:
f ¼
3 cos2 4 1 2
where 4 is the angle between the chain axis and the reference axis (stretching axis). hcos2 4i is defined as:
Zp=2 D E cos2 4 ¼
Ið4Þcos2 4 sin4d4
0
Zp=2
Ið4Þsin4d4
Fig. 13. Hermans orientation factor of crystals for S23 and S36 as a function of stretching ratio.
preferentially orients parallel to the tensile direction. Finally, a highly oriented structure with micro fibrils forms. The deformation process of S36 distinctly differs from that of S23. With increasing strain, the amorphous matrix deforms accordingly, however, the inner structure of the segregated crystalline region largely keeps intact.
0
where I(4) is the scattering intensity along the angle 4. The range of Hermans orientation factor is in the range of 0.5 to 1. When f is 0.5, it means that all the chains align perpendicularly to the reference axis. When f ¼ 1, all the chains are perfectly parallel to the reference axis. In our case, the c-axis orientation factor (fc) could not be measured directly due to the absence of (00l) plane, therefore the fc is obtained indirectly from the (110) and (200) reflections using the following expression [29]
D E D E D E cos2 4 ¼ 1 1:435 cos2 4110 0:565 cos2 4200 The calculated fc values in S23 and S36 are illustrated in Fig. 13. The fc of S23 increases asymptotically toward the value of unity with increasing stretching ratio, indicating that the chains gradually align along the stretching direction. For S36, however, the fc increases slowly to a small value of about 0.14 at l ¼ 12. This indicates that the orientation of crystals in S36 is very weak even at a very high stretching ratio. Fig. 14 is a schematic illustration showing the different deformation process of S23 and S36. The initial structure of S23 is featured as nearly homogeneous distributed lamellae, while the crystalline lamellae in S36 segregate into domains. Upon stretching, the lamellae in S23 slip with each other and fragment into blocks. With increasing the strain, the c-axis of crystals of S23 gradually
Fig. 12. FWHM of orthorhombic (110) reflection at different stretching ratio.
3.6. Factors influencing the elastomeric properties In the slip-link treatment of Hiltner and coworkers [27,28], crystalline phase acted as junctions, which contributed to the elastomeric response by increasing the sliplink density. Notwithstanding the limited amount, the crystalline phase, with much higher modulus than amorphous phase, also contributes to the mechanical properties itself as filler. The crystalline lamellae in S23 are dispersed in the amorphous matrix, while the lamellae in S36 are confined in isolated meso-domains. It is well known that fibrils/ lamellar fillers are more effective than particle fillers in reinforcing modulus. This is probably an important reason for the different modulus of S23 and S36. During tensile deformation, the crystal lamellae of S23 undergo rotating, tilting, slipping and finally breaking off and transforming to other phases. This indicates that crystalline phase has contribution to the tensile force in the whole strain range. However, the “hard” crystalline domains in S36 do not deform cooperatively, therefore the crystalline phase in S36 has less important contribution to the mechanical response. It has long been realized that microphase separation is important for the mechanical response of poly(styrene-b-diene-b-styrene) (SBS) thermoplastic elastomers [32]. Meanwhile, the success of ethylene random copolymer elastomers indicates that crystallization is also an important factor for elastomeric properties. It is natural to ask which factor governs the elastomeric properties in a system that has both mesophase separation and crystallization. Seguela and Prudhomme [33] reported that melt crystallized hydrogenated poly(butadiene-b-isoprene-b-butadiene) tri-block polymer (HBIB) exhibited much higher tensile strength than solution crystallized ones. The meltecrystallized samples undergo incompatibilityinduced microphase separation, whereas the solution-crystallized samples segregated due to crystallization. Koo et al. [13,14] got similar results. At a first glance, our results are contradictory to theirs. However, several important differences must be noted. Firstly, in their cases, the microphase separated morphology is 1-D [33] (cylindrical) or 2-D [13,14] (lamellar), while in our case it is 0-D (spherical). Cylinders or lamellae with high length in one or two direction are intrinsically easier to carry load than isolated spheres. To date, the relationship between micro/mesophase separation, crystallization and mechanical properties of block copolymers is still
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Fig. 14. Schematic illustration of the deformation mechanism of S23 and S36.
not fully understood. Moreover, there are other factors that also influence the mechanical properties. For instance, Park et al. [24] reported that the molecular weight between entanglements (Me) in OBC increased with increasing the commoner content, which could be another factor for the lower modulus of S36. Sarva et al. [34] reported that microstructure of elastomer had influence on the rate dependent stressestrain behavior. It has also been reported that the number of arms in star shaped block copolymers has strong influence on the mechanical properties [35,36]. Despite the complexity in factors governing the mechanical properties, we stress here based on our results that the cooperative deformation of crystalline phase makes an important contribution to the mechanical properties of crystalline elastomers. 4. Conclusions We investigated the mesophase phase separation and crystallization of multi-block OBCs with different chain composition. Crystallization within mesophase separated domains led to a densely packed lamellar structure with relatively stronger correlation between lamellae, showing a long spacing peak in SAXS. However, crystallization in OBC with lower segregation strength suppressed the mesophase separation. The uniform dispersed crystalline lamellae had lower inter-lamellar correlation, exhibiting a monotonically decreased intensity profile in SAXS. The two OBCs had distinct difference in mechanical behavior. One had a higher modulus and showed strain hardening while the other had a much lower modulus and almost no strain hardening. Tensile WAXS measurements showed different micro-structure transformation process. The orthorhombic crystal exhibited orientation, fragmentation and transformation to monoclinic crystal under stretching in one OBC; while the orthorhombic crystal in the other almost kept intact and had very low orientation under stretching. The filler effect of crystalline phase differed between the two OBCs due to the different morphology. The cooperative deformation of crystalline phase was proposed to make important contribution to the mechanical response of crystalline elastomers. Acknowledgments Financial support from National Natural Science Foundation of China is gratefully acknowledged (50903089, 21074141, 50925313).
Joachim Loos thanks CAS for his Visiting Professorship for Senior International Scientists (2009J2-28). The Dow Chemical Company is acknowledged for partial financial support. The SSRF is acknowledged for providing the beam time. Prof. Chenyang Liu (ICCAS) is thanked for constructive discussions. References [1] Hamley IW. Adv Polym Sci 1999;148:113e37. [2] Hamley IW. Developments in block copolymer science and technology. Chichester: John Wiley & Sons Ltd.; 2004. [3] Zhu L, Cheng SZD, Calhoun BH, Ge Q, Quirk RP, Thomas EL, et al. J Am Chem Soc 2000;122(25):5957e67. [4] Quiram DJ, Register RA, Marchand GR, Adamson DH. Macromolecules 1998; 31(15):4891e8. [5] Loo YL, Register RA, Adamson DH. Macromolecules 2000;33(22):8361e6. [6] Rangarajan P, Register RA, Adamson DH, Fetters LJ, Bras W, Naylor S, et al. Macromolecules 1995;28(5):1422e8. [7] Ryan AJ, Hamley IW, Bras W, Bates FS. Macromolecules 1995;28(11): 3860e8. [8] Quiram DJ, Register RA, Marchand GR. Macromolecules 1997;30(16): 4551e8. [9] Quiram DJ, Register RA, Marchand GR, Ryan AJ. Macromolecules 1997;30(26): 8338e43. [10] Loo YL, Register RA, Ryan AJ. Phys Rev Lett 2000;84(18):4120e3. [11] Loo YL, Register RA, Ryan AJ, Dee GT. Macromolecules 2001;34(26):8968e77. [12] Loo YL, Register RA, Ryan AJ. Macromolecules 2002;35(6):2365e74. [13] Koo C, Wu L, Lim L, Mahanthappa M, Hillmyer M, Bates F. Macromolecules 2005;38(24):6090e8. [14] Koo C, Hillmyer M, Bates F. Macromolecules 2006;39(2):667e77. [15] Hamley IW. The physics of block copolymers. London: Oxford University Press; 1998. [16] Arriola DJ, Carnahan EM, Hustad PD, Kuhlman RL, Wenzel TT. Science 2006; 312(5774):714e9. [17] Graessley WW, Krishnamoorti R, Balsara NP, Butera RJ, Fetters LJ, Lohse DJ, et al. Macromolecules 1994;27(14):3896e901. [18] Reichart GC, Graessley WM, Register RA, Lohse DJ. Macromolecules 1998; 31(22):7886e94. [19] Hustad PD, Marchand GR, Garcia-Meitin EI, Roberts PL, Weinhold JD. Macromolecules 2009;42(11):3788e94. [20] Li S, Register RA, Landes BG, Hustad PD, Weinhold JD. Macromolecules 2010; 43(10):4761e70. [21] Wang HP, Khariwala DU, Cheung W, Chum SP, Hiltner A, Baer E. Macromolecules 2007;40(8):2852e62. [22] Khariwala DU, Taha A, Chum SP, Hiltner A, Baer E. Polymer 2008;49(5): 1365e75. [23] Kamdar AR, Wang HP, Khariwala DU, Taha A, Hiltner A, Baer E. J Polym Sci Part B Polym Phys 2009;47(16):1554e72. [24] Park HE, Dealy JM, Marchand GR, Wang J, Li S, Register RA. Macromolecules 2010;43(16):6789e99. [25] Jin J, Du J, Xia Q, Liang Y, Han CC. Macromolecules 2010;43(24):10554e9. [26] Treloar LRG. The physics of rubber elasticity. London: Oxford University Press; 1967.
5230
G. Liu et al. / Polymer 52 (2011) 5221e5230
[27] Bensason S, Stepanov EV, Chum S, Hiltner A, Baer E. Macromolecules 1997; 30(8):2436e44. [28] Wang HP, Chum SP, Hiltner A, Baer E. J Appl Polym Sci 2009;113(5):3236e44. [29] Zuo F, Burger C, Chen X, Mao Y, Hsiao BS, Chen H, et al. Macromolecules 2010; 43(4):1922e9. [30] Wang HP, Chum SP, Hiltner A, Baer E. J Polym Sci Part B Polym Phys 2009; 47(13):1313e30. [31] Wilchinsky ZW. J Appl Phys 1960;31(11):1969e72.
[32] Holden G, Legge NR. Thermoplastic elastomers. 2nd ed. New York: Hanser Publishers; 1996. [33] Seguela R, Prudhomme J. Polymer 1989;30(8):1446e55. [34] Sarva SS, Hsieh AJ. Polymer 2009;50(13):3007e15. [35] Juhari A, Mosnácek J, Yoon JA, Nese A, Koynov K, Kowalewski T, et al. Polymer 2010;51(21):4806e13. [36] Mosnácek J, Yoon JA, Juhari A, Koynov K, Matyjaszewski K. Polymer 2009; 50(9):2087e94.