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C Solid State NMR Characterization of Structure and Orientation Development in the Narrow and Broad Molar Mass Disentangled UHMWPE
Yefeng Yao,†,* Songzi Jiang,† and Sanjay Rastogi‡,§,* †
Department of Physics & Shanghai Key Laboratory of Magnetic Resonance, East China Normal University, North Zhongshan Road 3663, 200062 Shanghai, P. R. China ‡ Department of Materials, Loughborough University, Leicestershire, LE11 3TU, England, U.K. § Research Institute, Teijin Aramid, Velperweg 76, 6802 ED Arnhem, The Netherlands S Supporting Information *
ABSTRACT: The ultimate draw ratio in semicrystalline polymers, and the resultant mechanical properties of the polymer, strongly depends on entanglement density in the amorphous region of the polymer. The influence of entanglement density becomes more pronounced with the increasing molar mass, for an example in the ultrahigh molecular weight polyethylene (UHMW-PE, having weightaverage molar mass greater than a million g/mol) the solidstate deformation (draw ratio >7) is feasible on crystallization of the polymer from dilute solution or during controlled polymerization using a single-site catalytic system. Here we address the influence of the molar mass distribution, associated with the polymerization conditions, on structural changes during solid-state deformation in the crystalline and the noncrystalline regions of UHMW-PE. With the help of various solid state NMR methods, differences in the deformation behavior of the mobile-amorphous, rigid-amorphous, and crystalline polymorphs have been followed in the broad and narrow molar mass UHMW-PEs. Orientation parameters arising at the segmental length scales, within different regions of the semicrystalline polymers, have been addressed. 2D 13C exchange NMR methods have been employed to follow the spatial proximity between the methylene segments of the noncrystalline regions (mobile- and rigid-amorphous phases) with the crystalline regions (crystalline core and crystal surface) during deformation of the two polymers. Distinct differences in the orientation parameters of the methylene segments in the noncrystalline and the crystalline regions, arising with the deformation of the broad and the narrow molar mass distributed UHMW-PE, have been observed and addressed.
1.0. INTRODUCTION High modulus and high strength in a uniaxially oriented semicrystalline polymer strongly depends on the crystalline component and the presence of entanglement density in the amorphous region. It has been shown that solid-state deformation of ultrahigh molecular weight polyethylene (having molar mass greater than a million g/mol) is feasible either by dissolution of commercially available entangled polymer (normally synthesized using a Z−N catalyst) close to the chain overlap concentration and its subsequent crystallization from the solution or by controlled synthesis using a single-site catalytic system. The solution route results into the viscous solution that can be spun to obtain fibers, whereas tapes are obtained on the uniaxial deformation of the polymer synthesized using the single-site catalytic system. Both approaches result in high tensile strength and high modulus light weighted materials either in the form of fibers or in tapes.1−4 Consequently, ultradrawn PE becomes a material of choice for many demanding applications, for instance, the hard ballistic for body armors, the high strength ropes and cables, © 2014 American Chemical Society
etc. A theoretical study predicts that the highly oriented methylene units in the ultradrawn PE fibers should have a Young’s modulus of ∼300 GPa and a tensile strength of ∼30 GPa, respectively.5,6 However, empirical values obtained are much lower than the theoretical prediction. Such discrepancies have been attributed to the incomplete orientation of chain segments in the ordered structure or the presence of the partially ordered component between the crystalline domains.7,8 In principle the partially ordered structures consist of chain segments that are not fully oriented,8 for example in the interphase between the crystalline domains. The presence of such less ordered structures influences the stress distribution in the ultradrawn PE and thus deteriorates the mechanical properties of the final products.9 Wide angle X-ray diffraction (WAXD) is widely used to determine chain orientation and crystallinity in the ultradrawn Received: October 28, 2013 Revised: January 20, 2014 Published: February 3, 2014 1371
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fibers due to its good resolution to detect the ordered structures.10−14 However, partially ordered and disordered structures possess much smaller correlation length compared to the perfectly ordered structures making it difficult to monitor using WAXD technique.14−16 Moreover, the information on spatial correlation of the observed structural components cannot be elucidated by WAXD, which hinders the understanding of stress transfer mechanism and its relation to the mechanical properties in the fibers.4,6,9 In this work, solid-state NMR technique is used to reveal structural changes and orientation development during uniaxial deformation of the solution-cast films of the ultrahigh molecular weight polyethylenes (UHMW-PEs) having narrow and broad molar mass distribution. First, with the help of 13C isotropic chemical shift in the 13C spectra structural differences in the crystalline region such as orthorhombic, monoclinic, crystal defects induced on shearing of orthorhombic crystals, and in the noncrystalline region such as the mobile amorphous domains mainly away from the crystalline region and the rigid amorphous domains near the crystalline region, are identified. The results obtained from the 2D 13C exchange NMR are used to identify the spatial proximity of the different structural components.17,18 With the knowledge of the signal assignment and the spatial proximity of the different structures, structural changes during deformation within the crystalline and the noncrystalline regions are followed by the 13C spectra. The study mainly focuses on the structural changes in the amorphous region, during deformation, and correlates to the molar mass distribution and the polymerization conditions. The 13C chemical shift anisotropy (CSA) of the methylene groups is examined to obtain the degree of chain orientation with the increasing draw ratio (DR) in different segments of the noncrystalline region of polyethylenes having narrow and broad molar mass distributions. The 13C CSA recoupling method19 is used to detect the chain orientation in the different structural regions, such as the crystalline region (i.e., orthorhombic, monoclinic, crystal defects, and crystal surface) and the two noncrystalline regions (mobile amorphous and rigid amorphous), in the drawn samples having different draw ratios. The influence from molar mass distribution on the chain orientation in the drawn samples is discussed thereafter.
determined by measuring marked points at 0.5 cm apart prior and after the deformation. Two commercially available PE fibres and tapes, namely, Dyneema SK75 (having a tensile strength of 3.5 GPa and a modulus of 130 GPa) and tapes made from disentangled UHMWPE22 (having a tensile strength of 4.0 GPa and a modulus of 200 GPa) respectively, were investigated for the comparison. The Dyneema fiber (abbreviated as UD-Fiber), which is considered as the ultra-drawn fiber, is made from the polymer synthesized using a Z−N catalyst. The ultra-drawn tape made from a disentangled UHMWPE (abbreviated as UD-Tape) is made from the polymer, having a narrow molar mass distribution, synthesized using a postmetallocene single-site catalytic system.22,23 The 13C melt-state NMR experiments show that all of the samples investigated in this work have very low branching degree (less than 1 branch per 10,000 CH2 backbone units). For the NMR study, samples of the solution cast films at different draw ratios were packed in the rotor either in the isotropic state or in the oriented state. The isotropic samples were prepared by packing the small pieces of the uniaxially drawn samples, whereas the oriented samples were prepared by using an adhesive to adhere the uniaxially drawn samples together. Schematic drawings shown in Figure 1 depict packing of the oriented and the isotropic samples investigated in this publication.
Figure 1. Packing of the oriented (left) and the isotropic (right) samples in a rotor. 2.2. General Settings in Solid-State NMR Experiments. The solid-state NMR experiments were performed with Bruker Avance III spectrometer operating at 600 MHz 1H Larmor-frequency. A 4.0 mm MAS double-resonance probe was used in most of the experiments. In the case of 13C CSA recoupling experiment the spinning speed of the rotor was set at 3 kHz, whereas in all other experiments the spinning speed of rotor was set at 10 kHz. The 90° pulse length in the experiments was varied between 2.5 and 3.0 μs on both channel, corresponding to ω/2π = 83−100 kHz. The CW or TPPM24 schemes were applied for dipolar decoupling with the decoupling frequency of ω/2π = 83−100 kHz. Ramped-CP was used for the experiments with the cross-polarization step. The 13C chemical shifts were determined from the carbonyl carbon signal (δ = 176.0 ppm) of glycine relative to tetramethylsilane (TMS). The temperature of the bearing gas was varied for the temperature dependent experiments. Signal deconvolution was carried out by using the free software program DMFIT.25 2.3. Distinguishing the Noncrystalline Components of Semicrystalline PE Using the 13C Single Pulse Excitation and Hahn Echo Filter. An amorphous signal was often observed in the 13 C single pulse excitation spectrum under MAS of solid PE at the slightly elevated temperature, e.g., 340 K. The width of the signal was related to many factors such as 1H decoupling efficiency during acquisition, the intrinsic T2 relaxation, the chemical shift distribution, and so on. By using Hahn echo as a filter,26 it was possible to select the amorphous signals having a long intrinsic T2 relaxation. Figure 2 shows the pulse sequence scheme of the 13C single pulse excitation with a Hahn echo filter. In this pulse sequence, by adjusting the echo time (ntR, tR is the rotor period.), it was possible to select the amorphous signals having a long intrinsic T2 relaxation while suppressing the signals having a short intrinsic T2 relaxation. 2.4. Following Conformation Exchange between the Noncrystalline and the Crystalline Regions of PE Using the 2D 13C
2.0. EXPERIMENTAL SECTION 2.1. Materials. The polyethylene (PE) samples studied in this work were prepared from two different grades of ultrahigh molar mass polyethylenes, namely, PE-1 and PE-2. The molecular characteristics of the two polyethylenes, synthesized using Z−N (PE-1) and postmetallocene single-site catalytic system (PE-2) are listed in Table 1. Solution cast films, used for tensile deformation from these
Table 1. Molecular Characteristics of the Two UHMWPE Samples sample
Mw (g·mol‑1)
Mw/Mn
Mn (g·mol‑1)
PE-1 (Z−N catalyst) PE-2 (single-site catalytic system)
4.5 × 10 4.5 × 106
8.0 3.0
0.56 × 106 1.50 × 106
6
polymers, were prepared by following the procedure described in detail elsewhere.20,21 To recall, UHMW-PE powder was first dissolved in xylene having 1.0 wt % of the nascent polymer and the viscous solution was casted in the form of a film. The samples for our study were prepared by solid-state deformation of the films obtained after evaporation of the solvent over several weeks. The samples having the draw ratios of 5, 10, and 30 were prepared, where the draw ratio was 1372
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Figure 2. Pulse sequence scheme of the with a Hahn echo filter.
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patterns of the different phase components, the SUPER pulse sequence was modified to have the initial source of 13C magnetization created via the cross-polarization (CP) transfer (CP-SUPER, Figure 4a), and the direct excitation using a short recycle delay (SP-SUPER, Figure 4b). More details about the pulse sequences can be found in the Supporting Information. Simulations of the 13C CSA pattern of powder or partially oriented samples were achieved by the NMR simulation software SIMPSON.27 The details about the simulation are given in the Supporting Information.
C single pulse excitation
3.0. RESULTS AND DISCUSSION With the development of the advanced experimental tools, it is becoming evident that the structural hierarchy exists in the semicrystalline polymers, where the simplest of all is the crystalline core having the well-defined density and chain packing specific to the polymer configuration and the associated conformation. The abstract nature of the noncrystalline region is strongly influenced by the crystallization history, molar mass, and molar mass distribution. In the utmost simplicity, in the nascent powder obtained directly from the reactor, the crystallization kinetics will be influenced by the chain growth during polymerization and the active sites concentration. These conditions will strongly influence the entanglement formation in the amorphous region with the implications on the dissolution of the polymer. For example, the experimental observation by us is that the nascent powder obtained from the Z−N catalytic system, having entangled amorphous region, requires slow stirring process for dissolution whereas the same concentration of the polymer synthesized using the single-site catalytic system requires larger stirring rate. The control in the stirring rate is required to avoid swelling and Weissenberg effect. This difference in the stirring rate requirement for the dissolution process can be associated with the differences in the number-average molar mass. However, the extent to disentangle an entangled chain prior to its crystallization will be also dependent on the entangled state of the nascent powder prior to its dissolution. To have some structural or topological information in the noncrystalline region of the semicrystalline polymer a series of NMR techniques have been applied. What follows are the structural
Exchange NMR. The conformation exchange between the noncrystalline and the crystalline regions of PE can be monitored by the 2D 13C exchange NMR.18 Figure 3 shows the pulse sequence scheme
Figure 3. Pulse sequence scheme of the single-pulse based 2D exchange NMR (SPEX).
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C
of the single-pulse based 2D 13C exchange NMR (SPEX). In this method, by using a relatively short recycle delay, structural components having a fast 13C T1 relaxation were polarized. In the semicrystalline UHMW-PE, structural components having the fast 13C T1 relaxation are the methylene units outside the crystalline core, including the mobile and the rigid amorphous components. The conformation exchange between the mobile/rigid amorphous components and the crystalline components can be followed by monitoring the signal exchange process in the 2D exchange spectra acquired by using different exchange times. 2.5. 13C CSA Analysis and Chain Orientation Determination. 13 C CSA of the semicrystalline PEs was site-selectively recorded by the two-dimensional NMR experiment, SUPER (separation of undistorted powder patterns by effortless recoupling).19 To obtain the 13C CSA
Figure 4. Pulse sequence of SUPER (separation of undistorted powder patterns by effortless recoupling)19 experiment: (a) the 13C magnetization is created via the CP step (CP-SUPER); (b) the 13C magnetization is created via the T1 relaxation during the preparation time (SP-SUPER). For more details on the figure and the methods used please see the Supporting Information. 1373
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Figure 5. (a) 13C CP/MAS spectrum of UD-Fiber. (b) 13C SP/MAS spectrum of UD-Fiber. Both experiments were performed at 340 K. For the 13C SP/MAS, the recycle delay was set at 3 s.
Figure 6. Left: Illustration of chain conformations in a semicrystalline polyethylene; Right: The 2D 13C single pulse based exchange spectra of the UD-Fiber. The experimental temperature is set at 340 K, and the applied exchange time is 1 s. Schematics on the left are used for explaining the exchange process between different domains of the noncrystalline phase.
supplementary section and during discussion of the results. In Figure 5a the noncrystalline signal (centered at ∼31.5 ppm) could be also resolved though the signal intensity is low. The low signal intensity is attributed to the smaller fraction of the amorphous phase and to the low 1H→13C CP efficiency.18 Figure 5b shows the 13C spectrum acquired from a single pulse excitation experiment under MAS (SP/MAS) using a recycle delay of 3 s. In this spectrum, the orthorhombic signal and the monoclinic signal are significantly suppressed due to the applied excitation method. Moreover, the crystalline signals are much broader compared to those in Figure 5a. The origin of the crystalline signals in Figure 5b can be attributed to chain diffusion between the crystalline and the noncrystalline regions that transfers the 13C polarization from the noncrystalline to the crystalline regions.18 Detailed mechanism involved in the chain diffusion process and the associated 13C polarization transfer is described elsewhere.17,38,39 Compared to Figure 5a, in Figure 5b the noncrystalline signals are strongly enhanced. This could be attributed to the short 13C T1 relaxation time of the noncrystalline component and the transient 1H→13C NOE enhancement because of the short recycle delay. An interesting asymmetric line shape is observed for the noncrystalline signal in Figure 5b. Considering the 13C 2D exchange spectrum of the same sample (Figure 6), where the two noncrystalline signals at ∼31.5 ppm and ∼30.8 ppm can be resolved along the diagonal line, a tentative decomposition of the noncrystalline signal in Figure 5b is performed. The results from the signal
variations observed in the two set of samples, depicted in Table 1, crystallized from solution. 3.1. Identification of Structural Variations in the Crystalline and the Noncrystalline Regions by Solid State 13C NMR. The isotropic 13C chemical shift of a semicrystalline polyethylene is sensitive to the local packing arrangement of the chain segments. Van der Hart et al.28−31 have conclusively demonstrated relationship between the 13C chemical shift of chain segments packed in the crystalline phases of polyethylene. On the basis of their study, we have assigned peaks in the 13C NMR spectra of the samples. Here, we exemplarily demonstrate the 13C CP/MAS (Figure 5a) and SP/MAS spectra (Figure 5b) of the ultradrawn PE fiber (UDFiber). These experiments were performed at 340 K. In Figure 5a, the dominant signal at ∼32.8 ppm is assigned to all-trans conformations in the orthorhombic crystal packing, and the small signal at ∼34.2 ppm is assigned to the all-trans conformations in the monoclinic phase.28,32,33 Between the orthorhombic and the monoclinic crystals, a small peak (∼33.4 ppm) is observed. In literature, this signal has been reported many times, and is assigned to intermediate/interfacial component,31,34 to the crystalline−amorphous interphase35 or to the second orthorhombic phase.32,36,37 In this study, the small peak is assigned to crystal defects, residing inside the crystalline core that is induced on shearing of the orthorhombic crystals during the uniaxial deformation. More details on the reasoning for the assignment of this signal can be found in the 1374
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Figure 7. (a) 13C CP/MAS and SP/MAS spectra of the solution crystallized PE-1 grade samples: the undrawn film, the uniaxially drawn samples having the draw ratios 5, 10 and 30, the ultradrawn sample (UD-Fiber); b) The 13C CP/MAS and SP/MAS spectra of the PE-2 grade samples: the undrawn film, the uniaxially drawn samples having draw ratios 5, 10 and 30, and the ultradrawn sample (UD-Tape). For comparison, the spectra of ultradrawn samples are taken on UD-Fiber and UD-Tape made from polymers having broad and narrow molar mass distribution. All experiments are performed at 340 K. The recycle delay in the 13C SP/MAS experiments was set at 3 s.
decomposition are shown in the table in Figure 5b. The broad noncrystalline signal at ∼31.5 ppm is assigned to the rigid noncrystalline components such as the taut-tie molecules and/ or the noncrystalline component close to the crystalline regions. The narrow noncrystalline signal at ∼30.8 ppm is assigned to the mobile noncrystalline phase away from the crystalline region, e.g., the amorphous bulk, the chain segments on the surface of voids between the microfibrils.8,40 From the decomposition of the noncrystalline signal, the content ratio between the broad and the narrow noncrystalline components in the UD-Fiber is approximately 8:3. With the knowledge of the signal assignment in the 13C NMR spectra, using the 2D 13C single pulse based exchange experiment; the spatial proximity of the different structural components of the semicrystalline polyethylenes is investigated. The UD-Fiber is used as an example to demonstrate the applied scheme. Figure 6 shows the 2D 13C single pulse based exchange spectrum of the UD-Fiber. In the spectrum, the clear ridges indicate the presence of the exchange process between the crystalline signal and the rigid noncrystalline signals. From the schematics shown in Figure 6, this exchange process can be associated between the all-trans conformations in the crystalline region and the gauche conformations in the rigid noncrystalline region. These findings are in accordance with our earlier studies on a nascent UHMWPE powder having folded chain crystals (see Figure 11b in ref 18). Considering the relatively short exchange time in the experiment, that will allow conformations on the crystalline surface to exchange with the conformation adjacent to them in the noncrystalline area, we assign the crystalline signal in Figure 6 to the structures on the crystal surface and the wide noncrystalline signal to the noncrystalline structures adjacent to the crystal surface. It is interesting to find that no exchange process is observed between the crystalline signal and the narrow noncrystalline signal centered at 30.8 ppm (see also Figure 11a in ref 18). These NMR studies indicate that the mobile amorphous component (∼30.8 ppm) is not adjacent to the crystalline regions, as shown schematically in Figure 6. However, the presence of disordered chains having mobile segments in the proximity of crystallites (but not close
to the fold surface) and do not contribute in the NMR exchange process cannot be ignored. To better demonstrate the difference in the conformation exchange processes between the two noncrystalline signals and the crystalline signal, two 1D slices have been extracted from the 2D spectrum. The slice extracted at 31.5 ppm in the one dimension shows two peaks having chemical shifts corresponding to the crystalline signal and the rigid noncrystalline signal. This indicates conformation exchange between the crystalline structures and the rigid noncrystalline structures (∼31.5 ppm) confirming proximity between the two structures. In contrast, the slice extracted at 30.8 ppm in the one dimension shows only one narrow peak at 30.8 ppm, indicating that no clear exchange process occurs between the mobile noncrystalline signal and the crystalline signal. To sum up, the above results show that the combination of the 1D solid state 13C spectrum and the 2D 13C exchange NMR provides a possibility for detecting spatial connectivity between the different structural components in the material, which is of relevance in our understanding of the stress transfer mechanism and its relation to the mechanical properties in the drawn materials. For more details on the conclusions made from the 2D 13C exchange NMR please see the Supporting Information. To monitor the development of the structural components during uniaxial deformation, what follows is the solid state 13C NMR on the uniaxially drawn samples of the solution cast UHMWPE films made from the two different molar masses, PE-1 and PE-2 at different draw ratios. 3.2. Structure Development in the Solution Cast Films of UHMWPE during Uniaxial Deformation. Figure 7 shows the 13C CP/MAS and SP/MAS spectra of the drawn samples of the PE-1 and PE-2 grades. In the 13C CP/MAS spectra of the samples, the orthorhombic signals do not change with the increasing draw ratio, while the intensity of the monoclinic signal changes with the increasing draw ratio. The intensity of the monoclinic signal strengthens with the increasing draw ratio. From the 13C SP/MAS spectra of Figure 7, it is apparent that the noncrystalline signals of the PE-1 and PE-2 samples develop differently with the increasing draw ratio. For example, 1375
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in the sample PE-1 with the increasing draw ratio the noncrystalline peak evolves into an asymmetric peak, which is indicative of the evolution of the noncrystalline mobile component. The observed asymmetric peak in the sample PE-1 drawn to the draw-ratio of 30 is similar to that of the solution-spun ultradrawn Dyneema fiber, UD-Fiber (see Figure 5b also). In contrast to the PE-1 grade samples, with the increasing draw ratio, the PE-2 grade samples show gradual broadening of the noncrystalline peak where the line shape remains almost symmetric. For an example, the noncrystalline peak of the disentangled UHMWPE sample drawn in the solid-state also shows an almost symmetric line-shape, but the peak width after deformation is nearly two times wider compared to that of the undrawn sample. The broadening of the peak width indicates increase in the conformation distribution, and presence of the restricted molecular dynamics in the noncrystalline regions. The restricted dynamics in the noncrystalline region causes slowing down of the exchange process between the conformations in the region. Unlike the PE-1 grade samples, no clear noncrystalline mobile signal is observed in the PE-2 grade samples, including the undrawn and the drawn samples. This indicates that no detectable mobile chain segments are present in the undrawn sample and, meanwhile, no mobile chain segments develop during deformation of this material. To have further insight on the structural variations in the mobile component of the two samples (PE-1 and PE-2), during the solid-state deformation, more detailed experiments and analysis have been performed. Comparison of the uniaxially drawn solution cast films of the two samples have been made with the solution spun Dyneema fiber (commercially available SK75) obtained from the polymer synthesized using a Z−N catalytic system, and solid-state drawn UD tape obtained from a polymer synthesized using the post metallocene single-site catalytic system.22 3.3. Monitoring the Mobile Noncrystalline Phase of the Semicrystalline Solution Crystallized Films during Uniaxial Solid-State Deformation. From the width of the noncrystalline signals in Figure 7, parts a and b, it is anticipated that the mobile noncrystalline phase will have a longer T2 relaxation than the rigid component. This difference in the T2 relaxation provides the possibility to selectively follow changes in the mobile noncrystalline signals while the rigid component using the T2 filter is suppressed. Figure 2 shows the pulse sequence scheme of the 13C single pulse excitation with a Hahn-echo block inserted before the reading pulse. The spectra acquired by applying this pulse sequence on the PE-1 grade samples, before and after the deformation, are shown in Figure 8. For comparison, the commercial sample (Dyneema SK75) having broad molar mass distribution, UD-fiber, is also investigated. It is evident that with the increasing Hahn echo time (i.e., the filter strength), the narrow noncrystalline signal, depicting the mobility, becomes more and more evident. The coexistence of the narrow and the wide noncrystalline signals in the spectra of all samples is strongly indicative of the inhomogeneous noncrystalline domain in these samples. The same pulse sequence is also applied on the PE-2 grade samples, before and after the uniaxial deformation, and the acquired spectra are shown in Figure 9. It is interesting to find that with the increasing Hahn echo time, the rigid noncrystalline signal becomes weaker, but no distinct narrow peak representing the mobile noncrystalline phase appears. Compared to the spectra in Figure 8, this clearly indicates that the
Figure 8. Filtered 13C SP/MAS spectra of the PE-1 grade samples with increasing Hahn echo times: (a) the undrawn film; (b) the drawn sample (DR = 30); (c) UD-Fiber. The recycle delay in the filtered 13C SP/MAS experiments was set at 3 s.
Figure 9. Filtered 13C SP/MAS spectra of the PE-2 grade samples with the different Hahn echo times: (a) the undrawn film; (b) the drawn sample (DR = 30); (c) UD-Tape. The recycle delay in the filtered 13C SP/MAS experiments was set at 3 s.
noncrystalline structures in these PE2 samples, having narrow molar mass distribution, is much more homogeneous than those in the PE-1 grade samples, both in the undrawn state and the drawn state. It is also possible to insert the Hahn echo filter in the 2D 13C exchange pulse sequence. This pulse sequence provides a way to selectively monitor the exchange process between the mobile noncrystalline signal and the crystalline signal. Figure 10 shows the 2D 13C exchange spectra of the two PE1-grade samples (drawn and undrawn) with and without the Hahn echo filter. For comparison the spectra of the UD-Fiber are also presented. In the 2D spectra acquired without the Hahn echo filter (Figure 10a−c), the presence of clear ridges indicates the exchange process between the crystalline and the amorphous signals (most likely, the rigid noncrystalline signal). In contrast, in the 2D spectra acquired with the Hahn echo filter (Figure 10d−f), only the diagonal signals, i.e., the narrow noncrystalline signal and the crystalline signals, can be observed. The spectra in Figure 10 thus further confirm the inhomogeneity of the 1376
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Figure 10. 2D 13C exchange spectra: (a) the undrawn PE-1 sample; (b) the drawn PE-1 sample (DR = 30), and (c) UD-Fiber. Filtered 2D 13C exchange spectra: (d) the undrawn PE-1 sample; (e) the drawn PE-1 sample (DR = 30), and (f) the UD-Fiber. The chosen echo time for measurement is 10 ms. In order to reduce the influence of 13C T1 relaxation on the signal intensity, an exchange time of 0.1 s is used in these experiments. The recycle delay in these experiments was set at 3 s.
Figure 11. (a) Illustration of the 13C CSA orientation in the all-trans conformation and the simulated 13C CSA patterns; (b, c) the 2D CP-SUPER spectra and the 13C CSA patterns of the orthorhombic structures of the PE-1 and PE-2 grade samples, respectively. The drawn samples are oriented and the undrawn sample (DR = 0) is isotropic. The sample packing in the rotor is illustrated schematically in Figure 1. In the extracted 1D 13C CSA patterns, the gray dash lines are the best-fit simulated patterns. The principle z-axis of the 13C CSA tensor of the methylene group is set to be along the chain axis with a Gaussian distribution having a width (fwhm) of ϕ (in deg). The lower the ϕ (in deg) the higher the chain orientation.
as shown schematically in Figure 11a. Motional narrowing of the 13C CSA interaction can provide detailed information about geometry of the local dynamics. For example, rotation by 180° in the crystalline regions does not change the CSA tensor. However, effective axial motion of extended all-trans segments around their local chain axis, leads to an average axially symmetric CSA tensor, where σ33 does not change while σ11 and σ22 are averaged. Precisely this kind of motional averaging is observed for the noncrystalline regions of the solutioncrystallized UHMW-PE sample. This axial motion of the
amorphous structures in the samples having broad molar mass distribution. To investigate the influence of the inhomogeneity in the orientation factor of the aligned chains along the fiber axis, further analysis has been performed in the following section. 3.4. Detecting Chain Orientation in the Different Structures of Semicrystalline PEs by 13C CSA Analysis. It has been shown that the 13C CSA analysis can be used to study the local chain dynamics in semicrystalline PEs.8,38,41 For extended chains, having all-trans conformations, the principal axes of the 13C CSA interaction are perpendicular to each other 1377
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Figure 12. Experimental 13C CSA patterns of the oriented drawn PE-1 grade samples: (a) the 13C CSA patterns of the monoclinic structures; (b) the 13 C CSA patterns of the signal between the monoclinic and the orthorhombic phases i.e. crystal defect phase on the surface of the crystalline core; (c) the 13C CSA patterns of the noncrystalline phase (extracted at 31.5 ppm). The gray dashed lines are the best-fit simulated patterns. The principle z-axis of the 13C CSA tensor of the methylene group is set to be along the chain axis with a Gaussian distribution having a width (fwhm) of ϕ (in deg). The lower the ϕ (in deg), the higher the chain orientation.
degree of chain orientation. The comparison shows that the chain orientation in the orthorhombic crystals of the drawn samples are ϕ = 16° for the sample having DR = 5, ϕ = 8.0° for the sample having DR = 10, ϕ = 7.0° for the sample having DR = 30, and ϕ = 5.0° for the ultradrawn commercial sample (UDFiber).8 Figure 11c shows the experimental 13C CSA patterns of the all-trans conformers in the orthorhombic crystals in the samples having the narrow molar mass distribution, i.e. crystallized during polymerization using the single-site catalytic system, PE2 grade samples. Similar to the spectra in Figure 11b, with the increasing draw ratio, the line shape of the 13C CSA pattern changes from the wide pattern to the narrow peak with two horns on the top. Simulation shows that the chain orientation in the orthorhombic structures of the drawn samples are ϕ = 20° for the sample having DR = 5, ϕ = 5.0° for the sample having DR = 10, ϕ = 4.0° for the sample having DR = 30, and ϕ =2.0° for the ultradrawn sample having the narrow molar mass distribution., For the same draw-ratio, compared to the Z−N synthesized samples (PE-1 grade), the metallocene synthesized samples (PE-2 grade) having the narrower molar mass distribution depicts higher chain orientationin the orthorhombic crystalline phase. Taking advantage of the signal resolution in the one dimension of the 2D SUPER spectrum, the 13C CSA patterns of the monoclinic structure, the crystal defect and the rigid amorphous structure of the oriented drawn samples can be measured at the same time. The 13C CSA patterns of these phase structures are shown in Figure 12. For comparison, the 13 C CSA patterns of the isotropic samples are also given. From parts a and b of Figure 12, it is clear that the 13C CSA patterns of the monoclinic structure and the crystal defect in the oriented drawn samples show narrowing of the line shape with the increasing draw ratio, indicating increase in the chain orientation of the drawn samples. The simulation shows that
segments is considered to be very effective for the chain diffusion. In this study, 13C CSA analysis is used to characterize the degree of chain orientation in different phases, crystalline and noncrystalline, of the drawn samples. The oriented and isotropic samples are prepared and filled in the rotor as depicted in Figure 1. To monitor the 13C CSA patterns of the different phases, the 2D CSA recoupling technique (namely, the SUPER experiment) is applied to the samples. A simulation program27 is used to derive the chain orientation. In the simulation, the crystal file for the isotropic pattern composes 3240 2-angle sets averaging over a hemisphere. The orientation is introduced by changing the weighting factors in the corresponding powder angle files. Figure 11a shows the simulated 13C CSA patterns of UHMW-PE. In the simulation, the principle z-axis of the 13C CSA tensor of the methylene group is set to be along the chain axis with a Gaussian distribution having a width (fwhm) of ϕ (in degree). The three principle values of the 13C CSA tensor used in the simulation are σ11 = 50.0 ppm, σ22 = 36.5 ppm, and σ33 = 12.5 ppm. The geometry of the 13C CSA tensor is illustrated schematically in Figure 11.30,38 From Figure 11a, it is apparent that with the increasing orientation (i.e., reducing ϕ), the patterns show a gradual change from the wide tensor-line-shape to the sharp tensor-line-shape with two horns on the top.8,28 Figure 11b presents the experimental 13C CSA patterns of the all-trans conformers in the orthorhombic crystals in the PE1-grade samples. It is clear that with the increasing draw ratio, the line shape of the 13C CSA patterns changes from the wide pattern to the narrow peak with two horns on the top. The observed change is similar to that depicted by simulation in Figure 11a, indicating that the degree of the chain orientation in the orthorhombic crystals of the drawn samples increases with the draw ratio. On comparing the experimental patterns with the simulated patterns, it is possible to determine the 1378
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Figure 13. Experimental 13C CSA patterns of the drawn polyethylene samples having narrow molar mass distribution (the PE-2 grade samples): (a) the 13C CSA patterns of the monoclinic structure; (b) the 13C CSA patterns of the crystal defect; (c) the 13C CSA patterns of the rigid amorphous phase (extracted at 31.5 ppm). The principle z-axis of the 13C CSA tensor of the methylene group is set to be along the chain axis with a Gaussian distribution having a width (fwhm) of ϕ (in deg). The lower the ϕ (in deg), the higher the chain orientation.
may cause a significant change in the chain conformation and thus changes in the local orientation of the 13C CSA tensor in the molecular frame, making difficult to estimate the local orientation in the noncrystalline region. Without the detailed information about the 13C CSA tensor orientation and local dynamics, empirically it is not possible to quantitatively analyze the orientation of the chain segments in the noncrystalline region of the semicrystalline polymer. We have also analyzed the SUPER spectra of the oriented drawn narrow molar mass samples, PE-2. Parts a and b of Figure 13 show the extracted 13C CSA patterns of the monoclinic phase and the crystal defects, respectively. It is clear that the 13C CSA patterns of the monoclinic phase and the crystal defect in the oriented drawn samples show narrowing of the line shape with the increasing draw ratio, indicating the increase of the chain orientation in the drawn samples. Simulation shows that chain orientation in the monoclinic phase and the crystal defects is higher in the ultradrawn PE-2 grade sample (UD-Tape) than in the ultradrawn PE-1 grade sample (UD-Fiber). Figure 13c shows the extracted 13C CSA patterns of the rigid amorphous structures. Similar to Figure 12c, the pattern of the ultradrawn sample indicates the possible chain orientation in the crystalline regions of the semicrystalline polymer, while the patterns of the drawn samples (DR = 5, 10, 30) indicate that the chain segments in the noncrystalline region of these samples do not gain any significant orientation 3.5. Chain Orientation on the Crystalline Surface. In the above-mentioned 13C 2D exchange NMR, fast conformation exchange between the crystalline and the noncrystalline regions of the drawn samples is observed. Taking advantage of this dynamic process, it is possible to selectively excite signal from the crystal surface by choosing a suitable exchange time in the exchange experiment.18 We have combined the exchange pulse sequence with the CSA recoupling pulse sequence and the pulse sequence named as SP-SUPER shown in Figure 3b. By applying this pulse sequence on the oriented sample, it is
the chain orientation in the monoclinic crystals of the samples are ϕ = 16° for the sample having the draw ratio 5, ϕ = 12° for the sample having the draw ratio 10, and ϕ = 8.0° for the sample having the draw ratio 30, and ϕ = 6.0 0 for the ultradrawn sample (UD-Fiber). The simulation assumes that the principle z-axis of the 13C CSA tensor of the monoclinic structure is along the chain axis. For the crystal defects, residing inside the crystalline core, the local orientation of the 13C CSA tensor in the molecular frame is not reported in literature. Considering that the line shape of the isotropic pattern is similar to that of the orthorhombic pattern, the local orientation of the 13C CSA tensor of the crystal defect is similar to that of the orthorhombic structure. With this assumption, we have fitted the patterns and determined chain orientation in the crystal defect region of the drawn samples, that is, ϕ = 20° for the sample having draw ratio 5, ϕ = 16° for the sample having draw ratio 10, ϕ = 10° for the sample having draw ratio 30, and ϕ = 8.0° for the ultradrawn sample (UD-Fiber). Figure 12c demonstrates the patterns of the noncrystalline structures of the oriented drawn samples. It is found that for the samples having DR = 5, 10, and 30, the pattern line shape is similar to that of the isotropic pattern (DR = 0), and does not change with the increasing draw ratio. The similarity in the line shape indicates that the noncrystalline structures in these samples do not gain orientation during uniaxial deformation. For the ultradrawn sample, having broad molar mass distribution, the pattern shows a much narrower line shape, indicating the presence of the chain orientation in this sample. In literature, such a line shape is reported for the cold-drawn high density PE by Mowery and Schmidt-Rohr.8 From the line shape of this pattern, with the help of simulation, it is possible to determine the chain orientation. However, the pattern simulation requires detailed information about the local orientation of the 13C CSA tensor at the molecular level. At the same time, chain dynamics in the noncrystalline regions 1379
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Figure 14. 2D SP-SUPER spectra: (a) the drawn PE1 sample (DR = 30) and (b) the drawn PE2 sample (DR = 30). The 13C CSA patterns are extracted at 33.0 ppm. Each sample is packed in the rotor and measured in both the oriented state and in the isotropic state. The exchange time is kept at 3 s, and the experimental temperature is maintained at 340 K. The principle z-axis of the 13C CSA tensor of the methylene group is set to be along the chain axis with a Gaussian distribution having a width (fwhm) of ϕ (in deg). The lower the ϕ (in degrees), the higher the chain orientation.
Figure 15. 2D SP-SUPER spectra of the drawn PE-1 sample (DR = 30) acquired by using the exchange times of 3 and 12 s. The crystalline 13C CSA patterns are extracted at 33.0 ppm. The pattern of the “isotropic crystal surface” acquired by using the exchange time of 3 s, and the pattern from the CP-SUPER spectrum are shown for comparison. The samples were measured in the oriented state in the rotor at 340 K. The principle z-axis of the 13 C CSA tensor of the methylene group is set to be along the chain axis with a Gaussian distribution having a width (fwhm) of ϕ (in degree). The lower the ϕ (in degrees), the higher the chain orientation.
possible to monitor the 13C CSA pattern on the crystal surface. Figure 14 shows the 2D SP-SUPER spectra of the drawn PE-1 and PE-2 grade samples. Both samples have a draw ratio of 30 (DR = 30). Each sample has been packed in both the oriented and isotropic states. The exchange time in the experiments was set to 3 s to ensure that the obtained crystalline signal is mainly from the crystal surface. The 1D 13C CSA patterns are extracted at 33.0 ppm. It is clear that the crystalline patterns of the two “isotropic” samples show a typical tensorial line shape, whereas the crystalline patterns of the same “oriented” samples show a narrow Lorentzian shape. The clear difference in the line shape indicates that the chain segments on the crystal surface have gained a certain degree of orientation. Simulation shows that the chain orientation on the crystal surface of the PE1 sample (ϕ = 20°), is distinctly less than that of the PE-2 sample (ϕ = 10°). With increasing the exchange time, the 13C polarization created in the amorphous regions progressively transfers into the crystal via the chain diffusion. This mechanism provides a way to detect chain orientation on the crystal surface at
different depths. Figure 15 presents the 2D SP-SUPER spectra of the drawn PE1 sample (DR = 30) acquired by using the exchange time of 3 and 12 s, and the 1D CSA patterns extracted at 33.0 ppm. It is clear that the pattern acquired by using the exchange time of 3 s is wider than that acquired by using the exchange time of 12 s. On simulating the patterns chain orientation of ϕ = 20° and ϕ = 12° is observed for the low and the high exchange times, respectively. The measured increase in the chain orientation with the increase in the exchange time indicates that the chain orientation increases with increasing depth from the crystal surface. But compared to the pattern extracted from the CP-SUPER spectrum, it is realized that the chain orientation at and below the crystal surface is still lower than that in the crystal core. Therefore, the CSA patterns in Figure 15 clearly demonstrate development of the chain orientation from the crystal surface to the crystal core. The same experiments have been performed on the drawn PE-2 sample (DR = 30). The SP-SUPER spectra and the extracted patterns are depicted in Figure 16. Similar to Figure 15, the pattern acquired by using the exchange time of 3 s is 1380
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Figure 16. 2D SP-SUPER spectra of the drawn PE-2 sample (DR = 30) acquired by using the exchange times of 3 and 12 s. The crystalline 13C CSA patterns are extracted at 33.0 ppm. The pattern of the “isotropic sample” acquired by using the exchange time of 3 s, and the patter from the CPSUPER spectrum are shown for comparison. The sample is measured in the oriented state at 340 K. The principle z-axis of the 13C CSA tensor of the methylene group is set to be along the chain axis with a Gaussian distribution having a width (fwhm) of ϕ (in degree). The lower the ϕ (in degrees), the higher the chain orientation.
wider than that acquired by using the exchange time of 12 s, indicating increase in the chain orientation with increasing distance from the crystal surface. Simulation yields the chain orientation of ϕ = 10° and ϕ = 5.0°, for the signals acquired using the exchange times 3 and 12 s, respectively. Compared to the pattern extracted from the CP-SUPER spectrum, it can be realized that the chain orientation below the crystal surface is still lower than that in the crystalline core (ϕ = 4.0°). However, when compared to the patterns of the drawn PE-1 sample, Figure 15, it can be concluded that for the same draw ratio the chain orientation on the surface of the PE2 sample is higher than that of the PE-1 sample.
Chain orientation in the different phases, crystalline and noncrystalline, of drawn UHMW-PEs can be quantified using the 13C CSA analysis. Crystalline core of the broad molar mass polyethylene shows lower chain orientation compared to the polyethylene having the low molar mass distribution. This difference in the chain orientation in the crystalline core can be attributed to smaller crystallite size in the uniaxially drawn broad molar mass polyethylene (Figure 11), where these crystallites are connected by mobile component having poor chain orientation. The chain segments in the crystalline surface of the drawn UHMW-PE show a clear orientation (Figure 14), whereas the rigid noncrystalline components, which are in close proximity to the crystalline surface (indicated by the 2D exchange NMR), do not show any clear orientation (Figures 12 and 13). This indicates that there is a sudden “gap” in the chain orientation between the rigid noncrystalline component and the crystalline surface area. But inside the crystalline area, the chain orientation gradually increases from the outside crystalline surface to the crystalline core (Figures 15 and 16). The chain orientation in the crystalline surface of the polyethylene sample having the narrow molar mass distribution is somewhat higher than that of the polyethylene having the broad molar mass distribution. The rigid noncrystalline phase, located between the crystal surface and the mobile noncrystalline phase, shows similar poor chain orientation independent of the draw ratio and the molar mass distribution (Figures 12c and 13c). The oriented rigid noncrystalline phase is only observed in the ultradrawn samples (UD-Fiber and UD-Tape). From these findings, it can be anticipated that the ultradrawn materials made from narrow molar mass distribution (or high number-average molar mass) will be mechanically superior than those made from broad molar mass polyethylene (or low number-average molar mass).
4.0. CONCLUSIONS In summary, the 13C solid state NMR studies have been carried out on the solution-cast films of the two UHMW-PEs, having broad and narrow molar mass distribution (or low and high number-average molar mass distribution, respectively). These studies have been performed on the samples drawn uniaxially at different draw ratios. Methylene segments orientation at different length scales ranging from crystalline (crystal core and crystal surface) to noncrystalline regions (rigid and mobile amorphous) as a function of draw ratio has been followed. By using 13C solid state NMR, it has been conclusively shown that the distinction between the noncrystalline structures, i.e., mobile noncrystalline and rigid noncrystalline structures can be made from different isotropic chemical shift and the signal width. The spatial connection between the noncrystalline structures and the crystalline structure can be achieved with the help of 2D 13C exchange NMR. The influence of the inhomogeneity in the noncrystalline region (mobile and rigid amorphous) on the exchange process becomes evident on using the desired filtration (Figure 10). The mobile noncrystalline phase is found to be nearly absent in the polyethylene having narrow molar mass distribution, i.e., the high number-average molar mass (Figure 9), whereas, in the polyethylene having the broad molar mass distribution (having the low number-average molar mass), this phase becomes prominent with the increasing draw ratio (Figure 8).
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ASSOCIATED CONTENT
S Supporting Information *
Monitoring the crystal defects in different UHMW-PE samples and 13C chemical shift anisotropy of methylene segments, simulation of 13C CSA pattern, and 13C 2D exchange NMR. 1381
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This material is available free of charge via the Internet at http://pubs.acs.org.
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AUTHOR INFORMATION
Corresponding Authors
*(Y.Y.) E-mail:
[email protected]. *(S.R.) E-mail:
[email protected]. Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS The authors wish to thank Teijin Aramid B.V., Arnhem, The Netherlands, for financial support. Y.Y. also acknowledges financial support from NSFC Grant No. 21174039.
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