Where Are the Chain Ends in Semicrystalline Polyethylene

Note Cite This: Macromolecules XXXX, XXX, XXX-XXX

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Where Are the Chain Ends in Semicrystalline Polyethylene? C. Wutz, M. J. Tanner, M. Brookhart, and E. T. Samulski* Venable and Kenan Laboratories, Department of Chemistry CB#3290, The University of North Carolina at Chapel Hill, Chapel Hill, North Carolina 27599-3290, United States

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n the late 1950s, researchers inferred that linear polyethylene (PE) crystallized with its chain contour ostensibly normal to the crystal surface;1−3 this picture was refined using diffraction and electron microscopy of collapsed, (sectored) tent-like crystals obtained from solution which indicated that the alltrans polymer “stems” within crystalline lamellae are tilted relative to the lamella normal.4 These seminal inferences, in turn, implied that the chain transverses the crystal multiple times in a “folded chain crystal habit”. From the outset, especially for melt-crystallized PE, the mechanism of crystallization and the detailed nature of the fold surface was a contentious subject,1,5−7 one that continues to challenge simulators.8 In semicrystalline PE the distinctions between the fold surface and the crystal’s amorphous surroundings are ambiguous in the two-phase “lamellar morphology” composed of crystalline lamellae embedded in a disordered amorphous phase.9 Adjacent and randomly re-entered chain-folded polymer “stems” make up the lamellae and the amorphous phase is widely believed to be composed of topological defects (loops and “tie chains”) and chain ends occluded from the crystalline lamellae.10−14 Herein we show with solid-state deuterium nuclear magnetic resonance (2H NMR) studies of specifically labeled PE that a large fraction of chain end sections (>0.8) are immobilized and therefore must reside in the crystalline lamellae. Our findings corroborate conclusions advanced three decades ago by VanderHart and Perez using carbon-13 NMR15,16 and the recent comprehensive modeling and carbon-13 NMR work of Fritzsching, Mao, and SchmidtRohr.17 The findings may have implications for the mechanism of polymer crystallization generally and give insights into how chains are accommodated in the lamellar morphology of semicrystalline polymers.18,19 Polyethylene has been extensively studied and serves as a prototypical polymer for the investigation of macromolecules in condensed phases. Inferences about macromolecular dispositions in semicrystalline PE are derived from considerations of chain trajectories in solution-crystallized polymers with the recognition that lamellae are held together by “tie chains” chains spanning two lamellaeand that topological defects including entanglements, loose folds (loops), and chain ends reside in the amorphous phase between lamellae (Figure 1).11 In the latter, both theoretical modeling and experimental evidence relegate chain ends to the exterior of the crystal. Keller and Priest20 demonstrated by ozonolysis experiments on monolayer single crystals of PE that 90% of the chain ends were excluded from the crystal. Small-angle X-ray scattering experiments by Cheng et al.21 and Schultz et al.22 using low molecular weight fractions of poly(ethylene oxide) with controlled contour lengths revealed that lamellae thickening and chain diffusion on annealing rapidly leads to crystals having an integral number of folds with the chain ends being fully © XXXX American Chemical Society

Figure 1. Chain topologies in a typical cartoon of the amorphous phase of the lamellar morphology of semicrystalline polymers: (1) loop or loose fold; (2) entanglement; (3) tie molecule; (4) chain end sections in the amorphous interphase.

excluded from the lattice. Theoretical considerations are consonant with this these findings: nonadjacent chain ends in the crystal constitute an energetically costly, so-called row vacancy defect.23 Herein we will use the term “chain end” to refer to the inherently mobile terminal section of a PE chain. Our definition of “chain end” derives from the roughly 35 ethylene units comprising the length of chain between entanglements in the PE melt. In the carbon-13 NMR work referenced above, investigators focused on the explicit terminal segments, i.e., the terminal −CH2CH3 unit resulting from chain transfer or vinyl units resulting from β-elimination.15−17 We are able to explicitly differentiate between chain ends and chain centers using selectively labeled polyethylenes (numberaverage molar mass ⟨Mn⟩ ∼ 20 000 Da). Our labeled PE materials contain a deuterated block containing ∼35 −CD2CD2− units. The labeled block is located either at the chain end (PE-Dend) or in the center of the chain (PE-Dcenter). By means of 2H NMR spectroscopy it is possible to examine the dynamical attributes of the differentiated parts of a PE chain in its semicrystalline morphology. We contrast our findings with observations on perdeuterated PE (PED) and earlier work by Spiess et al.24,25 They demonstrated that the 2H NMR spectrum of semicrystalline PED consists of two components: (i) a rigid part identified with chain stemssegments inside crystalswhich gives rise to a Pake powder pattern; (ii) a mobile component attributed to chain segments in the amorphous phase. Fast librational motion within PE crystals about the stem direction has been observed; its torsional amplitude increases from 5° at 40 °C to 12° at 110 °C.26 Received: September 8, 2017 Revised: October 22, 2017

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DOI: 10.1021/acs.macromol.7b01949 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules Subsequent 2D-exchange NMR spectroscopy revealed that slow segmental motionsdiffusive motion along the c-axis via 180° flips of an all-trans chain segment at rates of ∼105 s−1 also occur within the PE crystals.27

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METHODS

Polymer Synthesis. The PE samples were synthesized by living polymerization using a Co(III) catalyst, C5Me5(P(OMe) 3)Co− CH2CH 3+BAr4−, I, where Ar = 3,5-(CF3)2C6H3. The synthesis of I and the polymerization techniques have been previously described.28,29 The three labeled polymers prepared have average molecular weights of ca. 20K with low Mw/Mn values of 1.13−1.15, characteristic of materials produced in living polymerizations (Table 1). The diblock

Table 1. Properties of Deuterated PE Samples sample PEDend PEDcenter PED a

1st block no. C2H4 units

2nd block no. C2D4 unitsb

646

37

279

34

Figure 2. 2H NMR quad-echo spectra of selectively deuterated PE at different temperatures. The recycling delay was 2 s.

entire polymer ⟨Mn⟩ (MWD)a 19.1K (1.15)

454

544

Determined by GPC analysis. analysis.

3rd block no. C2H4 units

crankshaft motion, and it appears that only a small fraction of the chain ends undergo quasi-isotropic motions at 60 °C or higher temperatures. At 100 °C the line width of the mobile component of PE-Dend is much narrower than for the PED and PE-Dcenter samples. According to the generally accepted picture of the lamellar morphology in semicrystalline PE, the chain ends are thought to be in the amorphous phase and the free chain ends should be unencumbered relative to the labeled segments in the PE-Dcenter chains. But surprisingly, the fraction of mobile segments in the end-labeled chain is lower than in the center-labeled and the perdeuterated chains. Additionally, the 2H NMR spectra of PE-Dcenter display a higher fraction of mobile segments than PE-Dend and PED at all temperatures. In the intermediate motional regime, a quantitative line-shape analysis is difficult due to the influence of reduction factors in the 2H NMR spectra encountered when the time between the two pulses in the quadrupolar echo sequence is comparable to the motional correlation time.31 While the spin−lattice relaxation time measurements have the same deficiency, reliable values for the fraction of material contributing to different motional processes can be evaluated from spin−lattice relaxation measurements more expediently than detailed line shape analyses as the latter require fully relaxed spectra; i.e., these spectra would have to be acquired with very long relaxation delays (3−5 times the longest T1 ∼ 20 s). Consequently, the spectra in Figure 2 are T1weighted and thereby emphasize the contributions from the mobile segments. Figure 3 shows the relative longitudinal magnetization Mz(τw)/M0 of the three PE samples as a function of the logarithm of the waiting time τw at 20 °C. The symbols represent the experimental data, and the curves were obtained by best-fitting of a superposition of exponential functions according to

21.6K (1.14) 17.4K (1.13)

b

Determined by IR spectroscopic

polymer, PE-Dend, was prepared by exposure of I to C2H4 (1 atm, 4 h) followed by removal of C2H4 from solution and introduction of C2D4 (1 atm, 70 min). The triblock, PE-Dcenter, was similarly prepared by sequential exposure of I to C2H4, C2D4, and C2H4 with removal of ethylene from solution following each segment of chain growth. Average block lengths of the deuterated blocks listed in Table 1 were determined by quantitative infrared analysis. Initiation with I results in a −CH2CH3 end group from insertion of ethylene into the Co− CH2CH3 bond. Following chain growth, polymer is cleaved from the Co(III) center by hydrogenolysis which produces a second saturated end group, either −CD2CD2H (PE-Dend, PED) or −CH2CH3 (PEDcenter). Proton NMR analysis was carried out in deuterated odichlorobenzene at ca. 120 °C and shows only a CH2 singlet with no branches evident, indicating highly linear polyethylene which is supported by DSC analysis indicating Tmp = 136 °C typical of linear, high-density PE. 2 H NMR Measurements. The 2H NMR measurements were performed on a Bruker Avance 360 NMR spectrometer at a frequency of 55.28 MHz. The 2H NMR spectra were recorded using the quadrupolar (solid) echo pulse sequence30 with an excitation pulse length of 2.2 μs, a pulse separation time of 20 μs, and a recycling delay of 2 s. Approximately 3000 scans were accumulated for each spectrum. In order to measure the longitudinal relaxation, a progressive saturation pulse train consisting of ten 90x° pulses was applied prior to a variable waiting time, τw (= 1 ms−150 s), and the maximum intensity of the subsequent quadrupolar echo detection period was determined as a function of τw. The temperature dependence of the 2H NMR line shape for the three types of labeled PE is shown in Figure 2. In the first column typical line shapes are observed for PED with the Pake powder pattern dominating the 20 °C spectrum. The sharp edges of the “horns” of the powder pattern with a separation of δ = 120 kHz is indicative of immobilized segments (on a < 10−6 s time scale). The center of the PED spectrum exhibits a smaller, broadened doublet of approximately δ/3= 40 kHz. This feature is more pronounced at 60 °C and is indicative of chain segments constrained on both ends undergoing crankshaft motions;25−27 on increasing the temperature the fraction of such mobile segments grows, and at 100 °C there is evidence of some segments undergoing quasi-isotropic motion (the sharp resonance superimposed on the multicomponent PED line shape). Similar behavior is observed for the spectra of the selectively labeled PE-Dcenter polymer (second column, Figure 2). The line shapes for PE-Dend (third column, Figure 2) reveal little evidence of segments undergoing

n

Mz(τw) = M 0[1 −

∑ xi exp(−τw /T1(i))] i=1

(1)

The magnetization curves clearly demonstrate a two-step relaxation behavior with T1 times of the order of tens of milliseconds for mobile polymer and tens of seconds for the immobilized (rigid) polymer. For the mobile regime it is more appropriate to assume a continuous distribution of relaxation times (correlation times), which reflects the nonuniform distribution of free volume and variable constraints on the segments in the amorphous phase. As proposed by Rössler et al.,32 we assumed a logarithmic normal distribution of correlation times. σ p(ln τ ) = exp[− σ 2(ln τ − ln τ0)] (2) π Relaxation times and correlation times are related to each other by the BPP theory33 which is in most cases also valid for solids:34

⎤ 1 3π 2 2⎡ τ 4τ = C⎢ + 2 2 2 2⎥ ⎣ T1 10 1+ωτ 1 + 4ω τ ⎦ B

(3)

DOI: 10.1021/acs.macromol.7b01949 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules

before performing 2H-relaxation measurements. However, the fractions of mobile and rigid component remained virtually invariant relative to NMR measurements on samples prior to annealing. Normal PE exhibits lamellae long periods of ∼40 nm at ∼70% crystallinity, and the average stem length (immobilized chain length within the crystal) is therefore ∼25 nm.9,10 This dimension is much longer than the contour of the deuterium labeled blocks, ∼9 nm (≈ 35 × 0.25 nm). At 20 °C the global fractions of mobile and rigid PED segments reported in Table 2, xmob = 0.28 and xrig = 0.72, respectively, qualitatively agree with the earlier 2H NMR values reported by Spiess et al.,25,26 but there are apparent differences between their fully relaxed PED line shapes and those reported in Figure 2. We believe that the differences between PED samples⟨Mw⟩ = 100K (Mw/Mn = 10)26 versus the precision PED used here, ⟨Mw⟩ = 20K (Mw/Mn = 1.13) preclude a quantitative comparison with the earlier work of Spiess et al. However, there remains the possibility that signal intensity from segments characterized by motion with correlation times on the order of solid echo pulse separation (20 μs) is under-represented in both our spectral line shapes and our relaxation time measurements. Although this possibility cannot be rigorously eliminated, we do not observe anticipated qualitative changes in the spectra of Figure 2 for the precision-labeled PE end-segments when the temperature is increased from 20 to 100 °C. That is, if many labeled end-segments were in fact in the amorphous interphase, the expected reduction in correlation time for such segments (≪20 μs) should dramatically increase the observed fraction of mobile segments on raising the temperature by 80 °C. Instead, the mobile segment component reported here for PE-Dend account for only 17% of the labeled chain ends at 20 °C and only marginally increases to 22% at 100 °C (Table 2). The corresponding fraction of immobilized PE end-segments at room temperature (0.83) implies that the probability of finding chain end segments in the crystal is 40% higher than that found for the average chain segment in PED (0.72), and this situation is not changed significantly on raising the temperature to 100 °C. We are also cognizant of the tendency for phase separation when PE and PED chains are cocrystallized.35,36 However, the aggregation of deuterated segments would apply to both labeled chain centers as well as labeled chain ends. Aggregation does not appear to be a factor in our observations on PE-Dcenter; i.e., PED and PE-Dcenter exhibit similar behavior, and moreover, aggregation of labeled segments would presumably exacerbate occlusion of labeled chain ends from the crystalline regions of PE-Dend. Yet we find that the labeled chain ends are immobilized. In sum, selective deuterium labeling of polymers using the precision of living polymerization is a valuable tool for ascertaining nano- and mesoscale structural information in the condensed phase of polymers. Our 2H NMR experiments clearly indicate that PE chain ends are preferably incorporated into the crystalline lamellae in agreement with both earlier15,16 and more recent 13C NMR studies.17 But this conclusion conflicts with the generally accepted ideas about the disposition of chain ends in the lamellar morphology of semicrystalline polymers obtained from the melt. Why might polymer chain ends exhibit a proclivity for seeking out an energetically favorable “crystallographic burrow” within the lamellae rather than aimlessly explore the amorphous interphase? Several factors conspire to enable this: (i) chain end sections shorter than an entanglement chain length are inherently more mobile in the melt and rapidly sample a range of configurations near the advancing crystal interfaces; (ii) there is a thermodynamic (entropic) driving force for chain ends to segregate near (growing crystal) surfaces;37 (iii) having the chain’s terminus right at the lamellae surface dilutes packing conflicts within the interphase and such interfacial chain termination together with tilted chain stems away from the interface normal attenuates density anomalies near the crystal−amorphous interface.17 This integral role of chain stem tilt also impacts more subtle aspects of crystalline lamellae, namely, its role in propagating the observed chiral twist of lamellae in spherulites.38 All of these findings recommend that even the most schematic depictions of the morphology of the prototypical semicrystalline polymer, polyethylene, should be amended to show stem tilting and chain ends terminating at the lamellae crystal-interphase boundary, e.g., Figure 4.

Figure 3. Relative longitudinal magnetization Mz/M0 as a function of the waiting time τw for the three PE samples measured at 20 °C (symbols) fitted by calculated curves using eq 1 and the values in Table 2. Table 2 show for both 20 and 100 °C the fractions xi and the relaxation times T1(i) of the dynamically different components of the

Table 2. Relaxation Times and Fractions of the Mobile and Rigid Components (xmob and xrig) in the Three PE-D Samples Evaluated from the Longitudinal Relaxation Curves at 20 and 100 °C T [°C] 20 20 20 100 100 100

sample

xmob

T1,mob [s]

τmob [10−7 s]

xrig

T1,rig [s]

τrig [10−13 s]

PEDcenter PED PE-Dend PEDcenter PED PE-Dend

0.39

0.046

0.6

0.61

31.0

0.85

0.28 0.17 0.38

0.060 0.080 0.18

0.8 1.0 2.3

0.72 0.83 0.62

19.5 17.9 20.8

1.35 1.47 1.26

0.30 0.22

0.29 0.30

3.6 3.7

0.70 0.78

17.5 19.7

1.50 1.33

PE samples obtained from least-squares fits of the relaxation curves using eqs 1−3 with C = 160 kHz, ω = 2π·55.28 MHz, and σ = 1−1.5. At 100 °C the T1 values of the mobile components are increased by a factor of 4−5 relative to the 20 °C values. If we assume that all three samples have the same relaxation time for its rigid component (T1,rig = 25 s) and similarly for the mobile phases identical relaxation times (T1,mob = 0.060 s), the analysis does not significantly change the derived fractions for the dynamically different components; i.e., with this assumption we find mobile fractions of xmob = 0.40, 0.31, and 0.19 for PE-Dcenter, PED, and PE-Dend, respectively. The very short correlation times derived from the multiparameter fits to T1,rig may derive from rapid torsional librations;27 spin diffusion to a relaxation sink is apparently less likely.26 There is some evidence of the increased librational motion as the horns of the Pake powder pattern broaden as the temperature is raised (Figure 2). The evaluation of the fractions xmob and xrig agree with what was inferred from the line shapes in Figure 2. Namely, the end-labeled PE has a lower fraction of mobile segments than the center-labeled PE, and the latter chains have a higher fraction of mobile segments than the perdeuterated PED. All three of the PE samples have an overall degree of crystallization of xc = 0.70 ± 0.02 as determined by differential scanning calorimetry (DSC). The samples measured at ambient temperature were obtained with moderate cooling rates from the melt (approximately 10 °C/min), and it might be argued that the morphology does not represent the thermodynamically stable state. Therefore, samples were annealed under vacuum at 100 °C for 24 h C

DOI: 10.1021/acs.macromol.7b01949 Macromolecules XXXX, XXX, XXX−XXX

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Figure 4. Chain topologies in a corrected cartoon of the amorphous phase of the lamellar morphology of semicrystalline polymers; (1) loop or loose fold; (2) entanglement; (3) tie molecule; (4) chain end sections terminate at the crystal surface.

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected] (E.T.S.). ORCID

E. T. Samulski: 0000-0002-9067-5524 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We are indebted to three very conscientious referees for valuable suggestions and critiques. Both M.B. and E.T.S. thank NSF for support. E.T.S. is grateful for enlightening conversations with Michael Rubinstein and support from the Cary Boshamer Professorship.

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REFERENCES

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DOI: 10.1021/acs.macromol.7b01949 Macromolecules XXXX, XXX, XXX−XXX