Uniaxial Tensile Deformation of Poly(ε-caprolactone) Studied with

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Uniaxial Tensile Deformation of Poly(ε-caprolactone) Studied with SAXS and WAXS Techniques Using Synchrotron Radiation Tahseen Kamal,† Tae Joo Shin,‡ and Soo-Young Park*,† †

Department of Polymer Science, Kyungpook National University, #1370 Sangyuk-dong, Buk-gu, Daegu 702-701, Korea Pohang Accelerator Laboratory, Pohang University of Science and Technology, Pohang 790-784, Korea

‡

S Supporting Information *

ABSTRACT: The structural evolution of poly(ε-caprolactone) (PCL) during uniaxial tensile deformation at 25 °C was examined using small- and wide-angle X-ray scatterings (SAXS and WAXS) techniques with simultaneous stress and strain (S−S) curves. A high-energy X-ray beam at the recently upgraded Pohang synchrotron radiation source revealed the complete lamellar deformation behavior of PCL. Slope-based division of the S−S curves indicated three distinct regions of elastic (region I), yielding (region II) and plastic deformations (region III). In region I, which showed elastic deformation, the WAXS patterns were isotropic, whereas the SAXS patterns became oblate due to elongation of the amorphous chains along the draw direction. In region II, which showed yielding deformation, the WAXS patterns showed a slight orientation, whereas the SAXS patterns exhibited a change from oblate to four-point and to six-point patterns due to the simultaneous fragmentation and melting of the chain-folded lamellae (leading to the four-point pattern) and the subsequent formation of chain-extended lamellae (adding another two maxima along the meridian). In region III, the WAXS patterns revealed the development of the orientation of PCL crystals, whereas SAXS patterns exhibited a two-point pattern. The newly formed chain-extended lamellae in regions II and III might produce network junctions that can transfer an applied force to the PCL crystals for increased orientation. The six-point pattern in region II for PCL was not observed or reported in the past during the uniaxial tensile deformation experiment. This might be due to fast acquisition of the X-ray patterns during mechanical drawing using synchrotron radiation.

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mechanical properties,10 and biodegradability.11 Controlling the amorphous and crystalline phase content in a semicrystalline polymer imparts the final desired properties. The mechanical properties of PCL depend to a large extent on the microstructure and morphology. PCL is readily miscible with a large number of other polymers including poly(styrene-coacrylonitrile),12,13 poly(vinyl methyl ether),14,15 and poly(vinyl chloride).16 The compatibility of PCL with other semicrystalline polymers has attracted the attention of many researchers, and their phase separation upon cooling from the melt or during deformation processes was also investigated. These studies focused mainly on the phase separation of these blends by analyzing the structural details on the small and large length scales using small- and wide-angle X-ray scatterings (WAXS and SAXS) techniques.17 On the other hand, pure PCL has not been analyzed completely during tensile deformation at synchrotron radiation sources, even though structural investigations on the other important semicrystalline polymers, such as polypropylene, polyethylene, poly(ethylene terephthalate), and polyamides, during tensile deformation have been well documented. Moreover, melting and recrystallization of the

INTRODUCTION The simple chain architecture of many polymers forces them to crystallize to a certain extent, normally between 10 and 60%. The reason for less than 100% crystallinity is the kinetic hindrance of entanglements among the polymer chains, which cannot be abolished but can lead to the formation of amorphous layers during solidification from the melt. The process of solidification from the melt leaves the stacked lamellar crystals and entangled amorphous polymeric chains in between lamellae, forming a semicrystalline state.1 The periodicity between the lamellar crystals and amorphous layers in the bulk semicrystalline state occurs typically at a length scale of 10−50 nm. The hard crystalline and soft amorphous parts provide rigidity to the material and toughness to the system, respectively. The presence of these two phases imparts complex deformational behavior upon solid-state drawing. Therefore, an understanding of the deformational behavior is essential because a wide range of day-to-day applications require mechanical stability. Poly(ε-caprolactone) (PCL) is a well-known semicrystalline polymer with good mechanical properties, biocompatibility2 and biodegradability,3 and is recognized as a model semicrystalline polymer, an alternative to polyethylene.4−6 After the discovery of its crystal structure by Bittiger et al. in 1970,7 it was subjected to extensive studies of its crystallization behavior,8,9 © 2012 American Chemical Society

Received: August 14, 2012 Revised: October 19, 2012 Published: November 1, 2012 8752

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Figure 1. (a) S−S curves of PCL during the individual WAXS and SAXS measurements with the division of regions I, II, and III. (b) SAXS and WAXS patterns at different ε in regions I, II, and III; the numbers represent ε, and stretching direction was vertical.

lamellae are arranged in columns, and positions of the lamellae in the neighboring columns are uncorrelated. In another model, tilting of the lamellar surface can occur due to progressive shear between the crystalline stems within the lamellae.18,19 In this model, a two- or four-point pattern can be obtained depending on whether this surface is perpendicular or tilted away from the chain axis, respectively. This tilting of the lamellae can be distinguished by the WAXS patterns, which usually give the off-

lamellae during tensile deformation has been described loosely and requires a step-by-step and close analysis of the SAXS patterns to describe these phenomena completely. The registry of the lamellae gives rise to two- or four-point SAXS patterns. The origin of the two- and four-point patterns is unclear. In one model, a four-point pattern is obtained when the lamellae are arranged in a lattice that resembles a checkerboard, and a two-point pattern is obtained when the 8753

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equatorial hk0 reflections20 if the chain orientation is high enough. The contribution of the lamellae to the SAXS pattern often appeared as a bar that is often regarded as a two-point pattern but might be due to the coexistence of two- and fourpoint patterns. These two- and four-point patterns can be produced or transformed to other patterns during mechanical drawing. To investigate the origin of these two- and four-point patterns, acquiring the SAXS and WAXS patterns with high time resolution during mechanical drawing is necessary using a synchrotron radiation source. In this article, the transition from a four-point to two-point SAXS pattern was observed at the narrow region of the strain in the S−S curve while drawing the PCL film after yielding and before plastic deformation by monitoring the changes in the WAXS and SAXS patterns at the synchrotron radiation source. At this region, the four-point pattern overlapped with the twopoint pattern giving a unique six-point pattern. Such information is important for optimizing the performance of the drawn films (or fibers), predicting their properties with a given structure, and providing possible routes for improving the material.

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π

2

⟨cos α110⟩ =

cos α = cos θ110 cos ϑ

(2)

(3)

where θl10, ϑ, and α denote the Bragg scattering angle of the 110 plane and the azimuthal angles before and after correction, respectively. By solving the eq 1, the f110 values lies between the two extremes of −0.5 and 0. For a perfect orientation of the 110 plane with its normal perpendicular to the drawing direction, the f110 would be −0.5, while for an isotropic sample in the randomly oriented state, the f110 becomes 0. The degree of crystallinity was calculated from the WAXS patterns according to the equation

Xc =

Acr Acr + A am

(4)

where Xc, Aam, and Acr represent the apparent crystallinity and the area under amorphous and crystalline peaks (2θ range was 5°−22°), respectively, after deconvoluting the WAXS pattern (see Figure SI-3 as an example).

EXPERIMENTAL SECTION

3⟨cos2 α110⟩− 1 2

π

∫0 I(α) cos α dα

The above calculations were performed after correcting the azimuthal angle of 110 plane for the flat detector using the following Polanyi equation

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Sample Preparation, X-ray Measurements, and Analysis. PCL with a weight-averaged molecular weight of 80 000 was purchased from Sigma-Aldrich. The PCL chips were melted at 75 °C in a hot press between two steel plates covered with Teflon sheets. After keeping the sample in the molten state for 5 min in a hot press, it was quenched rapidly in ice water. This resulted in a 0.7 mm thick uniform film of PCL, which was subjected to mold cutting for the tensile deformation experiments. A newly upgraded undulated PLS-II 9A USAXS beamline of PAL (Pohang accelerator laboratory) was used for the X-ray diffraction measurements. The SAXS and WAXS measurements were performed with sample-to-detector distances of 2681 and 199 mm, respectively. The synchrotron X-ray radiation had a wavelength of 1.101 Å. The two-dimensional (2D) X-ray patterns were recorded on a CCD detector (Rayonix, USA). A Linkam TST350 tensile tester equipped with a 200 N load cell was used for the uniaxial tensile deformation of the samples during the in-situ experiments. The tensile tester was controlled via a PC through Linksys 32× system control software, which displays and saves the online plots of temperature, force, and distance. The distance and force data were changed into strain (ε = (l − l0)/l0, where l and l0 respectively are the specimen length during elongation and the initial specimen length) and stress, respectively. For tensile testing, dogbone-shaped sample bars with dimensions of 16 mm (length) × 5 mm (neck width) × 0.8 mm (thickness) were prepared. A constant uniaxial stretching velocity of 16.2 μm/s was applied to the specimen during drawing at 25 °C. The 2D patterns of the samples during the continuous stretching process were recorded with a 1 s exposure time and 1.5 s detector readout time. The 2D patterns were scanned with the FIT2D software package21 to obtain the one-dimensional (1D) patterns in the form of intensity vs q (scattering vector, q = (4π/λ) sin(θ/2), where λ is the wavelength of incident X-rays and θ is the scattering angle). The long period (L, distance between the adjacent lamellae along the drawing direction; see Figure 7) was calculated from the Bragg equation of L = 2π/qmax, where qmax corresponds to the peak position calculated using the Gaussian peak fitting method from the 1D SAXS patterns using Origin 7.0 software. Angle ϕ and layer distance calculations were made from the 2D SAXS patterns by ImageJ software. The Herman orientation factor for 110 plane, f110, was calculated by the equation22

f110 =

∫0 I(α) cos2 α dα

RESULTS AND DISCUSSION Division of S−S Curves and Structural Evolution. Figure 1a shows the simultaneous S−S curves during the individual SAXS and WAXS experiments. The two curves were similar so that the deformations at the same ε could be assumed to be similar during the individual SAXS and WAXS experiments, even though those experiments were performed separately. For the sake of simplicity, S−S curve has been divided into three distinct regions (based on slope), and a qualitative description of the SAXS and WAXS data has been given in this section for each region. The shape of the S−S curve is typical of semicrystalline polymers, i.e., the initial region with high slope, necked region with negative slope, and strain hardening regions with moderate positive slope representing the elastic, yielding, and plastic deformations, respectively. These regions were identified as regions I, II, and III in Figure 1a. Figure 1b shows the representative SAXS and WAXS 2D patterns in each region. Before deforming the sample, the SAXS patterns exhibited single isotropic scattering from the stacks of randomly oriented chain-folded lamellae. On the basis of the orthorhombic lattice structure of PCL crystals with a unit cell of a = 0.749 nm, b = 0.498 nm, c = 1.703 nm, and α = β = γ = 90°, the corresponding two strong 110 (innermost) and 200 (outermost) reflections were identified with several other weak reflections on the WAXS pattern. PCL possesses a larger repeating unit as compared to other polyesters. Unlike other semicrystalline polymers, where polymorphism is common and most often observed upon their uniaxial tensile/compression deformation or thermal treatment, no report has been found on the crystal−crystal transition in PCL. This implies the highest stability of the PCL crystal structure. The deformation process in region I before the yield point in the S−S curve did not affect the WAXS patterns. On the other hand, as ε increased in region I, the scattering maximum of the SAXS patterns began shifting toward a low angle along the meridian, whereas the scattering along the equator did not change its position with decreased intensity so that the patterns became oblate. The details will be discussed later in this article. The shifting of the scattering

(1)

where 8754

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maximum toward a low angle along the meridian was attributed to the increase in interlamellar distance, whereas the decrease in intensity along the transverse direction was due to the smaller amounts of lamellae oriented along the transverse direction. In the beginning of region II, the SAXS patterns changed gradually from oblate to a four-point pattern. At the same time, a decrease in stress began in the S−S curves, as shown in Figure 1a. In addition to the four-point pattern, an additional two scatterings (the two-point pattern) appeared at the meridian in the later stage, giving rise to a six-point pattern. The intensity of the two-point pattern increased with a concomitant decrease in the intensity of the four-point pattern as ε increased in region II, indicating that the existing lamellae were gradually melted down and transformed to newly formed lamellae representing the two-point pattern. The decrease in stress in the S−S curves and the appearance of a four-point SAXS pattern in the initial stages of region II suggest that the applied force was consumed by lamellar deformation in such a way that the chain-folded lamellar structure was fragmentized into small pieces, and those fragmented lamellae were tilted against the drawing direction and/or had a checkerboard arrangement to minimize the stress concentration at the lamellae. The scattering along the meridian might be due to crystallization of the extended chains escaping from the chain-folded lamellae during fragmentation after yielding and their uncorrelated arrangement along the stretching direction. To date, PCL has enjoyed a lot of deformation studies17,23 (usually in blends), but such a sixpoint pattern has not been reported or discussed. The observation of this six-point pattern of PCL was made possible through use of a strong incident X-ray beam from the undulated beamline of the recently upgraded PAL. These sixpoint patterns were observed only at the narrow range of ε = 9.5−12% so that the 1 s exposure time of high X-ray flux was needed in order to get good statistical data. The WAXS patterns in region II were changed slightly except for the slight orientation of the 110 and 200 reflections at the later stage. The detailed orientation will be discussed later in this article. In region III, four intensity maxima at the off-meridional direction in the SAXS patterns disappeared completely, and only two scatterings along the meridian remained. The arcs of the 110 and 200 reflections in the WAXS patterns became narrow as ε increased, indicating that the crystalline orientation developed during elongation of the film through strainhardening. Deformation in Region I. Region I represents the elastic deformation of PCL. Figure 2 shows L as a function of ε along the equator and meridian which was calculated from the azimuthally averaged q scans of the SAXS patterns along the mentioned directions (Figure SI-1). The L along the meridian and equator increased and decreased slightly, respectively, upon increase in ε. The L along the meridian increased from 16.2 to 17.1 nm as ε was increased from 0 to 0.05 (5%), so that ΔL/L0 (0.9/16.2) is 0.05 (5%), which is close to the applied strain, where ΔL and L0 are the increased and initial L, respectively. Therefore, the elastic deformation in region I along the meridian occurred in the amorphous region and was almost affine. The reason why the L along the equator decreased slightly might be due to the reorientation of the lamellae oriented along the equator to the meridian. Most of the force might be consumed for elongation of the amorphous chains so that the WAXS patterns changed little in Figure 1a. Deformation in Region II. As mentioned previously, the lamellae in the early stages of region II were fragmentized into

Figure 2. Change in L as a function of ε along the equator and the meridian in region I.

small pieces, and those fragmented lamellae were tilted against the drawing direction and/or had a checkerboard arrangement to minimize stress at the lamellae. During fragmentization of the lamellae, the chains escaped from the initial chain-folded lamellae into the amorphous region, became extended, and easily crystallized into new chain-extended lamellae oriented along the stretching direction with a lack of axial register between the crystals giving a two-point pattern. Figure 3 shows

Figure 3. Intensity at the maximum of the azimuthally averaged q scans of the SAXS patterns as a function of ε along the meridian and the equator with azimuthal angles of 45° and 98°, respectively.

the intensity at the maximum of the azimuthally average q scans of the SAXS patterns as a function of ε along the meridian and equator with azimuthal angles of 45° and 98°, respectively (see Figure SI-2). This azimuthal angle along the meridian and equator covers the scattering from two-point and four-point patterns, respectively. The intensity along the equator decreased continuously with increasing ε, even though the intensity along the meridian initially decreased until ε = 9% but increased at higher ε values. Before ε = 9%, the four-point pattern was not fully oriented due to the low lamellar orientation so that the scattering from the chain-folded lamellae contributed to the intensities along the meridian and equator together. At ε > 9%, the two-point scatterings along the 8755

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by slippage between the lamellae as ε increased. On the other hand, the L of the newly formed lamellae from the two-point pattern was almost constant at 15.3 nm. This L value was much shorter than those of the fragmented chain-folded lamellae less than half of the saturated L from the four-point pattern (35 nm). These results strongly suggest that the newly formed extended-chain lamellae were crystallized in the extended amorphous region between the cleaved lamellae. The WAXS patterns in region II did not change significantly. A slight orientation was observed at the last stage in region II. The detail degree of the orientation will be discussed later in this article. Deformation in Region III. Strain hardening occurred in region III. The stress in the S−S curve increased slowly with increasing ε in this region due to plastic deformation. In the SAXS patterns, only two-point patterns along the meridian were observed except for some weak traces of the four-point pattern in the early stage. L decreased slightly from 15 nm at ε = 10.4% (beginning of the region III) to 14.6 nm at ε = 70% (end of the region III). During plastic deformation, slippage between the lamellae occurred continuously, and the lamellae were divided into many pieces with a relatively constant L. On the other hand, the decrease in L indicates that strain-induced crystallization in the amorphous region might occur simultaneously during slippage between the lamellae. This straininduced crystallization might play a role as an anchoring point during mechanical drawing. This anchoring point could orient the lamellae because of the efficient transfer of the applied force through these junctions. In the WAXS patterns, the orientation developed significantly as ε increased, as shown in Figure 1b. Development of the Orientations of the Crystals, Crystallinity, and Models. Figure 5a shows the azimuthal scans of the 110 reflection as a function of ε which were utilized to determine the f110 and full width at half-maximum (FWHM). At low ε (0−6%) in region I, the scattering intensity was high and uniform against the azimuthal angle due to the isotropic orientation of the 110 reflection. Figure 5b shows the plots of f110 and fwhm against ε. In region II, the orientation of the crystals began to develop at ε = 9%, and its fwhm (and f110) continued to decrease from 60° after ε = 9% (see Figure 5b). The two-point scattering along the meridian in the SAXS pattern was also observed at ε = 9%, as shown in Figure 1. Therefore, the development of the crystal orientation was strongly correlated with the newly formed lamellae, which were produced from the more oriented amorphous chains and might play a role as a cross-linked point (force transferring point) in the amorphous region for further development of the crystal orientation. In region III, the crystals continued to be oriented, and their FWHMs decreased with further elongation. The degree of orientation was developed mainly in region III due to the large amounts of extended-chain lamellae compared to those of the chain-folded lamellae in this region. The chainfolded lamellae in Figure 1b contributed to the azimuthal scanning with an isotropic manner, even though the extended chain lamellae contributed to it with an anisotropic manner. The isotropic scattering (background scattering) would give the same intensity along the azimuthal angle so that the intensities at the meridian (the same as the background intensity) and equator above the background were due mainly to the chainfolded and extended chain lamellae, respectively. Figure 5c shows the intensities of the 110 plane at the meridian and equator as a function of ε. The intensity at the equator was calculated above the background. The 110 plane

meridian were clearly separated from the four-point patterns in Figure 1b. Therefore, the scattering along the meridian after ε = 9% was due mainly to the newly developed two-point patterns. On the other hand, at ε < 9% in region II, the chain-folded lamellae were expected to be transformed continuously to the extended-chain lamellae because the intensity along the equator (mostly due to the chain-folded lamellae) decreased continuously with progressing the ε in region II. To analyze L in the four-point SAXS pattern, the layer distance was calculated from the SAXS pattern. In the fourpoint SAXS pattern, the Y value of the Cartesian coordinates (X, Y) of the maxima of the four-point in Figure 4a represents

Figure 4. (a) Typical SAXS pattern in region II with the definition of angle ϕ and x, y values. (b) L from both the four-point and two-point SAXS patterns and ϕ from the four-point patterns as a function of ε.

the layer distance, which can be calculated by 2π/Y.24 The L is the distance between the adjacent lamellae along the stretching direction so that the layer distance in the reciprocal four-point SAXS pattern represents the L in the checkerboard arrangement in real space. The angle ϕ from the equator represents the angle between the normal of the lamellae and the stretching direction. Figure 4b shows the L from both the four-point and two-point SAXS patterns and ϕ from the four-point patterns as a function of ε. The L from the four-point pattern increased continuously with increasing ε and was saturated at 35 nm after ε = 10%. The L/L0 (35/16.2) after saturation was 2.2, which was almost double the elongation of the film (1 + ε = 1.1). This suggests that the deformation was no longer elastic and affine in the amorphous region (like in region I) and the fragmentation of a lamellar into small ones contributed to the large L as compared to the actual elongation of the film. The ϕ value decreased continuously with increasing ε, indicating that the amounts of stagger between the adjacent lamellae was increased 8756

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Figure 6 shows the change in the apparent crystallinity as a function ε. The constant crystallinity of 0.465 in the region I

Figure 6. Variation of the apparent crystallinity as a function of ε.

started to decrease in region II, and the crystallinity became a constant of 0.405 in region III. In region II, six-point SAXS patterns were observed at the expense of fading-out of intensities of the four-point due to the melting of the chainfolded lamellae and subsequent formation of the chainextended lamellae. The decrease of the crystallinity in region II and lower crystallinity in region III (than that in region I) indicates that the newly formed chain-extended lamellae did not have as much crystallinity as the chain-folded lamellae. Bao et al.25 also observed the reduction of the crystallinity of isotactic polypropylene at 25 °C during tensile deformation. Lamellar deformation mechanisms have been widely discussed in the literature. Galeski,26 Bartczak,27 and Bowden et al.28 proposed the interlamellar and intralamellar slip mechanism while the stress-induced melting and recrystallization was considered to be responsible for the variation of morphology in the deformation process.29,30 Experimental evidence for the above arguments have been extensively reported including microscopic and X-ray diffraction investigations supporting the slip mechanism28,31 and SAXS experiments favoring the melting−recrystallization scheme.32−34 Recently, Men et al.20 explained that the slip mechanism happened at the initial small deformation and was followed by the stress-induced melting and recrystallization during deformation of high-density polyethylene. The SAXS data provided in this article manifest the clearer picture of melting and recrystallization phenomena during lamellar deformation. Figure 7 shows a schematic of the lamellar deformation during in-situ tensile experiment on the basis of SAXS data. The model in region I shows the elongation of the amorphous part while the orientation of the chain-folded lamellae did not change. In region II, the lamellae were fragmentized into a checkerboard-like arrangement. At the end of the region II, coexistence of the chain-extended and chain-folded lamellae was illustrated to represent the six-point SAXS pattern. In the regions III, small chain-extended lamellae were depicted after the complete melting of the chain-folded lamellae with decreased L.

Figure 5. (a) Azimthal scans of the 110 reflection, (b) f110 and FWHM (calculated from the peak at 180° in (a)), and (c) intensity values at the meridian and the equator as a function of ε.

was oriented along the equator so that the intensity at the meridian was affected by the chain-folded lamellae only. The intensity of the 110 reflection at the meridian was constant until ε = 9%, decreased until ε = 14% (the end of region II), and remained low at higher ε values (region III). Therefore, the amounts of the isotropically oriented chain-folded lamellae mainly decreased in region II. The intensity at the equator was almost opposite that at the meridian. The intensity was almost 0 until ε = 9% and increased slowly at ε > 9% with a decreased slope. This strongly indicates that the initial chain-folded lamellae were melted down and recrystallized into extended chain lamellae. From this orientation study, the contribution of the chain-folded and extended chain lamellae could be separated and the transformation from the chain-folded lamellae to the extended chain lamellae was clearly observed. Region II is located in the middle of this transformation. Therefore, the six-point pattern in region II occurred during the melting of the existing chain-folded lamellae and recrystallization of the chain-extended lamellae during mechanical drawing. 8757

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Figure 7. Models in regions I, II, and III for PCL lamellar deformation during in-situ uniaxial tensile deformation where ϕ and L represent the staggering angle in Figure 4 and long period, respectively.

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CONCLUSIONS The lamellar and crystal structural evolutions of PCL during tensile deformation at room temperature were examined using SAXS and WAXS techniques. The three regions in the S−S curves representing elastic, yield, and plastic deformations were observed. In region I, amorphous chains were elongated toward the drawing direction without a change in the PCL crystals of the chain-folded lamellae. In region II, where there was fragmentation of the chain-folded lamellae into checkerboardlike arrangement, melting of the chain-folded lamellae and simultaneous recrystallization of the chain-extended lamellae were observed. In region III, the chain-extended lamellae were further developed simultaneously with a high degree of orientation of PCL crystals. During this deformation process, an oblate SAXS pattern in region I, a transition from a fourpoint to a mixed six-point patterns in region II, and a two-point pattern in region III were observed in the SAXS pattern. The six-point pattern in region II occurred during the melting of the existing chain-folded lamellae and recrystallization of the chainextended lamellae during mechanical drawing. The newly found six-point pattern could be observed with high time resolution of the X-ray patterns during tensile deformation using the recently upgraded synchrotron radiation beam at PAL.

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ASSOCIATED CONTENT

* Supporting Information S

Figures SI-1−SI-3. This material is available free of charge via the Internet at http://pubs.acs.org.

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REFERENCES

AUTHOR INFORMATION

Corresponding Author

*Tel +82-53-950-5630; Fax +82-53-950-6623; e-mail psy@knu. ac.kr. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by MEST and PAL, XFEL project, Korea and the National Research Foundation of NRF-20110020264, Korea, and experiments at PLS-II were supported in part by MEST and POSTECH. 8758

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dx.doi.org/10.1021/ma301714f | Macromolecules 2012, 45, 8752−8759