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Chain Trajectory of Semicrystalline Polymers As Revealed by SolidState NMR Spectroscopy You-lee Hong,*,† Wei Chen, Shichen Yuan, Jia Kang, and Toshikazu Miyoshi* Department of Polymer Science, The University of Akron, Akron, Ohio 44325-3909, United States ABSTRACT: Over the last half century, a chain-folding structure of semicrystalline polymers has been debated in polymer science. Recently, 13 C−13C double quantum (DQ) NMR spectroscopy combined with 13C selective isotope labeling has been developed to investigate re-entrance sites of the folded chains, mean values of adjacent re-entry number ⟨n⟩ and fraction ⟨F⟩ of semicrystalline polymers. This viewpoint highlights the versatile approaches of using solid-state (ss) NMR and isotope labeling for revealing (i) chain trajectory in melt- and solution-grown crystals, (ii) conformation of the folded chains in single crystals, (iii) self-folding in the early stage of crystallization, and (iv) unfolding of the folded chains under stretching.
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structure (trajectory) of the melt-grown crystals.34 The DQ method provides a measure of the apparent dipole−dipole couplings for a given 13C site, which represents the RMS sum of many 13C−13C pair couplings. It provides the number and topology of interacted spins at internuclear distances less than ∼7 Å. In the crystalline system, atomic coordinates of the 13Clabeled sites were determined by WAXD35 and further refined by DQ-NMR.34 Thus, the experimental DQ buildup curve was utilized to trace the chain trajectory of 13C-labeled polymers in the crystalline regions by utilizing the influence of 13C spins. Figure 1A−E shows several chain-folding models, including CF0 (isolated model) and CFI-III of iPB1 in form I crystals, with the corresponding simulated DQ buildup curves of 35%
rystallization of long polymer chains induces drastic structural change from random coils in melt and solution state into folded chains in thin crystalline lamellae with thickness of 5−20 nm.1−3 Much attention has been given to understand the chain trajectory of polymers over the last half century, because chain-level structures include specific information regarding when, where, and how long polymer chains fold during crystallization.4−20 In the early days, neutron scattering (NS) combined with 2H isotope labeling proved that radius of gyration (Rg) of polyethylene (PE) in the melt-grown crystals is comparable to that of the melt state and that Rg in the solution-grown single crystals is much smaller than that of the melt state.12 Furthermore, local chain-folding patterns of PE and other polymers have been investigated by NS,13−17 infrared (IR),18−20 surface decoration,21 and atomic force microscopy (AFM) for direct visualization22−24 and force detection;25 however, there were inconsistent views on the detailed chainfolding structures. Moreover, most characterization techniques require either specific morphology, 21,22 monodisperse oligomer,16 or large supercooling.13−15,17,18 Thereby, it is still challenging to systematically study the kinetic effect on the chain-folding structure of semicrystalline polymers in both melt- and solution-grown crystals. ssNMR spectroscopy is a very powerful tool to characterize the structure and dynamics of organic molecules and inorganic materials. In polymers, during the past two decades, molecular dynamics (time scale and geometry of molecular motion) in the glassy,26 crystalline,27−29 and melt state30,31 and local chain packing32 and conformation33 in the glassy state have been successfully characterized by modern ssNMR techniques. Several works that combined ssNMR with isotope labeling, such as 13C, 2H, and 15N, highlighted structure and dynamics for specific sites of the polymer chains.27,29,33 Recently, Hong and Miyoshi focused on selective 13C isotope labeling of CH3 in isotactic-poly(1-butene) (iPB1) and dipolarbased 13C−13C DQ NMR spectroscopy to elucidate chain-level © XXXX American Chemical Society
Figure 1. Four chain-folding models of (A) CF0, (B) CFI, (C) CFII, and (D) CFIII and (E) the corresponding simulated DQ buildup curves for 35% 13C-labeled iPB1 form I crystal blends with nonlabeled iPB1 at a mixing ratio of 1:9 under infinite folding.34,38 Received: January 18, 2016 Accepted: February 9, 2016
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ACS Macro Letters C CH3-labeled iPB1 as a function of excitation time.34 CF0 is the major structure in solidification, switchboard, and random re-entry models and shows the slowest buildup curve with the lowest height (pink curve in Figure 1E) among the four CF models. On the contrary, different chain-folding patterns, as shown in the CFI-III models, result in largely different DQ curves,38 for example, the CFII model, having zigzag reentrance pattern, leads to the highest spin density around a detected 13C spin resulting in the fastest and highest DQ curve (red curve in Figure 1E) among the CFI-III models. Simulation results emphasize that the DQ approach is sensitive to identify different chain-folding patterns. When isotope labeling is used for structural analysis in crystalline systems, special care must be taken to avoid segregations of the labeled and nonlabeled components during crystallization. In the case of 2H and 1H PE blend, slow crystallization induced the segregation of the two components.35 Thus, most experiments were performed under rapid crystallization.13−15,17 In the 13C labeling case, since the mass difference between 13C and 12C is quite small, and since singlesite labeling in the monomer unit was used at a low labeling ratio of 15−35%,34,36−42 cocrystallization of 13C-labeled and nonlabeled polymers at the stem level was observed even at very high Tc (e.g., isotactic-poly(propylene), Tc = 150 °C).34,36−42 Moreover, ssNMR can selectively observe the crystalline signal by filtering out the amorphous signal using pulse techniques. Thereby, comparing samples crystallized under wide Tc range, the kinetic effect on the local chain-level structure of 13C-labeled polymers in the crystalline regions could be analyzed by 13C DQ NMR with a relaxation filter.34,36−42 According to the secondary nucleation theory,3,4 it is expected that the adjacent re-entry number and fraction change depending on the crystallization conditions. In refs34 and38, successive adjacent re-entry number ⟨n⟩ and adjacent re-entry fraction ⟨F⟩ were defined as the ensemble average over all chains of stem number in adjacent re-entry clusters and of the ratio of stem numbers in the adjacent re-entry clusters to the total stem number, respectively. Hong et al. investigated the chain-folding structure of 13Clabeled iPB1 with ⟨Mw⟩ = 37 K g/mol in solution-grown crystals as a function of Tc.38,39 AFM and transmission electron microscopy (TEM) demonstrated that iPB1 crystallized at Tc = 60 °C adopts form I hexagonal crystals, which have well-defined growth planes, while iPB1 at Tc = ∼0 °C results in circular crystals (kinetic roughness). Considering the thickness of the hexagonal single crystals and the ⟨Mw⟩, the maximum folding number ⟨nmax⟩ was estimated to be 21. By comparing DQ experimental and simulation curves based on several chainfolding models in a single row (Figure 1B−D), it was found that only the CFII model is feasible to reproduce the experimental data at Tc = 60 °C (black open circle in Figure 2A). After adjusting either ⟨n⟩ or ⟨F⟩ values, two limit models of ⟨nmax⟩ = 21 at ⟨F⟩ = 90%, and ⟨nmin⟩ = 8 at ⟨F⟩ = 100% could reproduce the experimental data (black open circle), as depicted in Figure 2A.38,39 The DQ curve at Tc = ∼0 °C (green open circle in Figure 2A) was consistent with the DQ curve at Tc = 60 °C. This result indicated that, at low concentration (0.03 wt %), the majority of the individual chains adopt nearly perfect adjacent re-entry structure and create clusters of single molecules at both Tc = 60 and ∼0 °C.38,39 The change in crystal habits from hexagonal to circular shape was reasonably explained by kinetically driven depositions of single 13
Figure 2. (A) 13C−13C DQ buildup curves for 35% 13C CH3-labeled iPB1 in solution-grown single crystal blends with nonlabeled iPB1 at a mixing ratio of 1:9, crystallized at Tc = 60 (black open circle) and ∼0 °C (green open circle) and melt-grown crystal blend at Tc = 95 (black triangle) and ∼0 °C (green triangle) with the corresponding chainfolding models in solution- (B) and melt-grown crystals (C).38
molecule clusters on the growth front (Figure 2B). This explanation contradicts the secondary nucleation theory, which hypothesizes that microscopic chain-level structure, dominated by kinetics, determines the macroscopic morphology.3,4,12 According to the nucleation theory, secondary nucleation induces two-dimensional (2D) conformation (linear) of the folded chains under sufficiently low supercooling, while primary nucleation induces 3D conformation (see Figure 3C,D).1
Figure 3. (A) Relaxation filtered 13C CPMAS (black) and DQ (green) NMR spectra for 35% 13C CH3-labeled iPB1 in form III single crystal blends with nonlabeled iPB1 at a mixing ratio of 1:9, formed at Tc = 50 °C and asymmetric orthorhombic packing structure of form III. (B) The DQ buildup curves for the 13C-labeled CH3 signals at 14.2 (A site) and 15.1 ppm (B site) and the corresponding simulated curves based on the 3D conformation of 12 stems in four rows at ⟨F⟩ = 65%.40 Possible chain-folding mechanisms and folded chain conformations (C) on the growth front and (D) in dilute solution.40 356
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coarse grained Poly(vinyl alcohol) (PVA) predicted ⟨n⟩ to be 2.4 at ⟨F⟩ = 100% under wide supercoolings.44 The simulation result suggested that the chain-folding structure is determined by the chain length between entanglement points. The ⟨n⟩ value obtained in the simulation is very close to the ⟨n⟩ value in iPB1 determined by DQ NMR. Understanding the structural formation during early stages of crystallization is one of the challenging issues in polymer science. Small angle X-ray scattering (SAXS), wide-angle X-ray diffraction (WAXD), and IR enabled to observe structural evolution of polymers prior to crystallization.43 Olmsted et al.9 and Kaji et al.10 proposed conformational assisted spinodal decomposition (SD) prior to nucleation and growth below the SD line. Yuan et al. reported chain-folding structure of the mesomorphic form of iPP which was made by rapid and deep quenching to glass transition temperature at ∼0 °C.37,41 This condition limits lateral growth and thus retains the structure formed during the early stage of crystallization. The mesomorphic domain includes only ∼60 stems and its size is comparable to critical nuclei in the early stage of crystallization predicted by the MD simulation.45 Figure 4 shows the 13C−13C
Determination of the cluster shape (2D vs 3D conformation) and the size of the folded chains under low supercooling is a particularly important subject to elucidate crystallization mechanisms. Unfortunately, the DQ curve at single resonance of iPB1 form I could not distinguish between the 2D and 3D conformation.38,39 Note that the simulated curve based on 2D conformation leads to approximately 10% higher ⟨F⟩ value than using 3D conformation.38 To confirm the conformation of the folded chains and folding mechanism, Hong et al. chose form III of iPB1 with ⟨Mw⟩ = 37 K g/mol (⟨nmax⟩ = 13 and single crystal thickness of 8.5 nm) among the various semicrystalline polymers.40 Form III shows orthorhombic packing structure leading to magnetically inequivalent doublet 13C peaks for each carbon (Figure 3A). Thus, dipolar networks at two sites in the single stems were used to recognize either 2D or 3D conformation for the folded chain. Under the assumption of any 2D conformation, simulated DQ curves at two sites, could not simultaneously reproduce the experimental data at both sites. On the other hand, 3D conformation (cluster) models, including 14 stems in two rows and 12 stems in four rows, could reproduce the experimental data by adjusting the ⟨F⟩ value (Figure 3B).40 Moreover, rapid quenching to ∼0 °C also led to the same 3D conformation. The observed 3D cluster and the T c independence of the cluster size supported (i) the formation of the baby nuclei induced by self-folding as the initial step and (ii) the subsequent depositions of the baby nuclei on the growth front as the second step (Figure 3D). The proposed mechanism provides a falsification of secondary nucleation theory,3,4 but is consistent with the theoretical bundle5,6 and aggregation models.8 In the condensed melt state, polymer chains are highly entangled. Therefore, individual chain-folding events may interfere with each other due to topological constraints. In ref 38, chain-folding structure of iPB1 form I crystallized at a wide range of Tcs of 95 to ∼0 °C was investigated. Figure 2A represents the DQ curves for iPB1 in the melt-grown crystals (black and green open triangles at Tc = 95 and ∼0 °C, respectively). Through the detailed analysis, three important findings were obtained: (i) The chains in the melt-grown crystals adopt the same zigzag pattern as determined in the solution-grown crystals. (ii) Under the assumption of ⟨F⟩ = 100%, ⟨n⟩ is ∼2 within the wide Tc range. (iii) The ⟨n⟩ value is much shorter than that in the solution-grown crystals. The last two indicate that the concentration and entanglement of polymer significantly influence adjacent re-entry sequence length. Possible chain-folding patterns in the melt-grown crystals obtained by changing ⟨F⟩ and ⟨n⟩ values were schematically illustrated in Figure 2C. Furthermore, the kinetic effect on chain-folding structure was investigated for iPP (⟨Mw⟩ = 172 K g/mol) melt-grown α crystals at Tc = 100 and 150 °C.36 iPP α crystals show Tc dependence on packing structures: Low Tc of ∼100 °C induces disordered α1 crystals, while high Tc of 150 °C induces ordered α2-rich crystals. It was expected that α1 and α2 crystals have largely different chain-folding numbers and fractions. However, DQ experiments clearly demonstrated that the iPP chains adopt ⟨n⟩ = 5−7 under the assumption of ⟨F⟩ = 100% in both α1 and α2 crystals. These two examples in the melt-grown crystals indicate that the chain-folding number, under fixed ⟨F⟩ values, does not change as a function of Tc. These newly obtained results contradicted expectations from the secondary nucleation theory.3,4 Recently, molecular dynamics (MD) simulation of
Figure 4. (A) 13C DQ buildup curve for 15% 13C CH3-labeled iPP in the blend sample with nonlabeled iPP with a mixing ratio of 1:9 in mesomorphic form. Green and red DQ curves were simulated based on (B) solidification model and (C) cluster folding model with ⟨n⟩ = 4 and ⟨F⟩ = 100%.37
DQ buildup curve for 15% 13C CH3-labeled iPP blends with nonlabeled iPP in the mesomorphic form and simulated curves based on solidification model (green curve) and stem clusters formed via self-folding (red curve). This comparison supported that the chains self-fold 3−4 times during mesomorphic formation.37 Thus, it was concluded that self-folding of the iPP chains initiated nucleation of mesomorphic form preceding crystal formation. On the basis of chain mobility in the crystalline regions, Hu and Schmidt-Rohr categorized semicrystalline polymers as either fixed or mobile crystals.46 The latter shows large deformability at temperatures above crystal relaxation temperature. The effect of deformation on crystallographic structure47 and Rg48,49 of individual chains has been well studied. However, structural evolution of locally folded chains under stretching has not been detected due to experimental difficulty. Kang et al. applied DQ NMR to understand the deformation mechanism of iPP chains under uniaxial stretching at 100 °C, where iPP stems perform helical jump motions.42,46 It was found that folded chains are gradually unfolded with increasing engineering strain (e), and that chain conformation is locally extended (⟨n⟩ = 0) for e ≥ 10 at 100 °C. This result confirmed that stem dynamics plays an important role for large deformability of iPP.46 Additional experiments at various temperatures and 357
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different stretching rates will provide systematic understanding of deformation of semicrystalline polymers at the molecular level. In summary, we highlighted the recent progress on understanding the chain-level structures of semicrystalline polymers under crystallization as well as under deformation using ssNMR spectroscopy. It was demonstrated that (i) the concentration and entanglement of polymers play significant roles for chain-folding selection while available kinetics does not, (ii) polymer chains self-fold in dilute solution and in the early stage of bulk crystallization. The important findings were useful to validate theories,5−7 falsify the well established secondary nucleation,3,4 and understand crystallization mechanisms9−12 at the molecule level. Further studies on the chainfolding structure in solution-grown crystals from dilute to highly concentrated solutions is helpful to understand how concentration and entanglement influence chain-folding events from entire molecules to a part of molecules. From a kinetics view, it is necessary to characterize chain-level structure in quenched glass. Furthermore, this novel approach will be applied to study chain-level structures in related areas including confined crystallization,50 crystallization and self-assembly at liquid/liquid interface,51,52 and stereocomplex of chiral polymers,53 and other processing such as shear flow and electrospinning.54
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AUTHOR INFORMATION
Corresponding Authors
*E-mail:
[email protected]. *E-mail:
[email protected]. Present Address †
RIKEN CLST-JEOL Collaboration Center, Yokohama, Kanagawa 230-0045, Japan (Y.-l.H.).
Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This work was financially supported by the National Science Foundation (Grant Nos. DMR-1105829 and 1408855) and a UA start-up fund.
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REFERENCES
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