Article Cite This: Macromolecules XXXX, XXX, XXX−XXX
Yield Stress Enhancement in Polyethylene−Glassy Diblock Copolymers William D. Mulhearn and Richard A. Register* Department of Chemical and Biological Engineering, Princeton University, Princeton, New Jersey 08544, United States S Supporting Information *
ABSTRACT: The low-strain mechanical properties of linear polyethylene (PE) can be substantially altered by the incorporation of a short block of a polymer with a high glass transition temperature (Tg ) into a majority-PE block copolymer. In particular, the yield stress and the tensile modulus can be sharply increased with the addition of a relatively small fraction of glassy block, especially when combined with a thermal history that promotes high crystallinity and crystal thickness of the PE block. For example, the incorporation of 15 wt % of a hydrogenated poly(norbornyl norbornene) block (Tg = 115 °C) into a PE diblock copolymer, cooled from the melt at ∼1 °C/min, doubles the yield stress (from ∼30 to ∼60 MPa) and tensile modulus (from ∼1.5 to ∼3.5 GPa) relative to the values for a PE homopolymer treated with the same thermal history. These property enhancements are closely associated with the composition of the amorphous layer between crystal lamellae and the spatial distribution of the glassy block within this layer. Finally, the ductility of these polymers at high strains is governed by the presence or absence of tie molecules, which can be correlated with the chain dimensions of the PE block in the melt and the distance between crystal lamellae.
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INTRODUCTION Despite the widespread use of polyethylene (PE) as a low-cost tough plastic, it is not well suited to certain applications due to its relative ease of permanent deformation. Under an applied tensile load, semicrystalline ethylene-based polymers undergo a sequence of deformation mechanisms at progressively higher strains.1,2 First, conformational rearrangements of chains within the interlamellar amorphous layer allow for largely reversible, elastic stretching.1−3 Second, upon reaching a critical yield stress, a fine-slip mechanism allows the crystal stems to reorient parallel to the applied tension.1,2,4,5 This transition, referred to as the first yield point, is easily visible during a tensile test as a local maximum in a plot of stress vs strain. Third, the crystal lamellae fragment as sections are pulled apart by the highly extended tie chains.1,2,6 Crystal fragmentation is often visible as a second local maximum or shoulder in the stress at moderately higher strain and is accordingly referred to as the second yield point. The yield process is also associated with the formation of a “neck” with reduced cross-sectional area, which propagates along the specimen at progressively higher strains.7,8 The latter two processes in this deformation sequence are only partially reversible at ambient temperatures,2 and so the first yielding event defines the limit of a semicrystalline article’s usefulness for many applications. Consequently, routes to enhance the yield stress are of particular interest. For example, high-density PE homopolymers typically exhibit a yield stress (σy) of approximately 30 MPa.9 This value can be manipulated modestly by controlling the crystal thickness5 (Lc) and crystallinity10 (wc), for example, via thermal history. To © XXXX American Chemical Society
Figure 1. Low-strain tensile behavior of perfectly linear polyethylene, Mn = 72 kg/mol.
illustrate, Figure 1 displays the low-strain tensile behavior of two specimens of the same perfectly linear PE (Mn = 72 kg/mol), prepared either by slow cooling from the melt at ∼1 °C/min or by rapid quenching from the melt into roomtemperature water at ∼1000 °C/min. The slow-cooled specimen exhibits an elevated crystal thickness and crystallinity (measured by small-angle X-ray scattering and differential scanning calorimetry) relative to the rapidly quenched Received: November 20, 2017 Revised: December 6, 2017
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DOI: 10.1021/acs.macromol.7b02454 Macromolecules XXXX, XXX, XXX−XXX
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Figure 2. Reaction schematic for the synthesis of hPNbN−PE diblock copolymers. The Schrock-type initiator is represented by a molybdenum atom with associated ligands (Lx), bonded to the neophylidene group that constitutes one of the chain ends in the final polymer. imidoneophylidenemolybdenum(VI) bis(tert-butoxide) (Strem Chemicals), was used as received. Toluene was purified via an MBraun solvent purification system. Trimethylphosphine (PMe3; SigmaAldrich, 97%) was purified by degassing and stirring over sodium prior to vacuum transferring to a collection vessel. Propionaldehyde (Sigma-Aldrich, 97%) was purified by degassing and stirring over 3 Å molecular sieves prior to vacuum transferring to a collection vessel. Diblock Copolymer Preparation. Polymerizations were carried out inside a nitrogen-filled glovebox (MBraun UNIlab, 0) present in the PE−hPNbN system, which is not considered in the aforementioned SCFT treatment,34 wherein the grafted and matrix polymer chains are chemically identical (χ = 0). The other major trend in mechanical properties across the series of G-PE diblock copolymers is the loss of ductility evident at high glassy block contents (and lower molecular weights), particularly in samples cooled slowly from the melt. By taking ⟨R2⟩0/M = 1.25 Å2 mol/g for molten PE,36 the rootmean-square end-to-end distance (simply denoted R0) can be estimated for the PE blocks in the melt for each of the four polymers. If R0 is much less than the measured lamellar spacing (d) for a sample with a particular thermal history, we expect that few chains will be sufficiently extended to span two or more crystals (tie chains), leading to embrittlement.37,38 Figure 9 compares the strain at break (εb) with the ratio R0/d and reveals a rather sharp transition between brittle and ductile behavior at R0/d ≈ 0.81. This observation is consistent with prior work on linear PE in which a polymer with Mw = 71 kg/ mol and Đ = 2.22 could exhibit either brittle or ductile behavior depending on its thermal history.9
and the diblock copolymer. The most significant difference between slow-cooled PE and G-PE 15 is that E′ and E″ are consistently higher for the diblock for all temperatures below the α-transition, and E′ for the diblock is largely independent of temperature below ∼25 °C while E′ for the homopolymer gradually decreases with temperature over that range. Although an amorphous layer composition of wglass,am ≈ 0.4 is accompanied by an abrupt stiffening of the material and a suppression of dislocation motion (fine slip) in the crystallites (Figure 6), it does not correspond to a uniformly vitrified amorphous phase. To demonstrate, Tg was measured for a series of random ethylene/hydrogenated norbornyl norbornene copolymers of varying hNbN contents, shown in Figure 8. The
Figure 8. Glass transition temperatures measured by DSC for random copolymers of ethylene (hydrogenated cyclopentene) and hydrogenated norbornyl norbornene (hNbN) as a function of the weight fraction of hNbN units. The dashed horizontal line denotes room temperature (23 °C), and the solid curve is a quadratic fit to the data to approximately extrapolate to lower hNbN contents.
observed trend is qualitatively consistent with published findings for other PE-containing copolymers,30,31 which indicate that Tg vs composition does not follow the classic Fox equation and that Tg begins to increase significantly only at relatively high comonomer content. While the calorimetric (DSC) glass transitions in homogeneous blends and block copolymers are often observed to be broader than in corresponding random copolymers,32 the mechanical transitions (e.g., steepness of modulus−temperature curve or width of tan(δ) peak) of such homogeneous systems do not typically show a comparable extent of broadening,33 perhaps due to a difference in the length scales probed by DSC and DMTA. The random copolymers in Figure 8 indicate that a uniformly mixed amorphous fraction would require ∼60 wt % hPNbN to become glassy at room temperature; in contrast, the characteristic composition of ∼40 wt % hPNbN, associated in Figure 6 with an abrupt increase in stiffness, corresponds to a Tg near 0 °C. Therefore, it is likely that the high-hPNbN samples have undergone some degree of component segregation within their amorphous layers, such that stiff regions rich in the glassy block have formed to impart the observed enhancements in the tensile modulus. Furthermore, the rise in the materials’ yield stresses implies that these hPNbN-rich regions are preferentially located near the crystal fold surface. This preferential enrichment of hPNbN blocks near the crystals can be explained by the semicrystalline diblock’s chain topology. The PE component of each of the diblock copolymers is highly crystalline (see the legends of Figure 6), F
DOI: 10.1021/acs.macromol.7b02454 Macromolecules XXXX, XXX, XXX−XXX
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within the disordered melt state, and these reduced molecular weights led to a loss of tie molecules. As such, it should be possible to design hPNbN−PE copolymers in such a way as to achieve these enhancements in low-strain mechanical properties with little, if any, loss of ductility. To this end, we propose a few design rules for the preparation of hPNbN−PE copolymers that combine these advantageous low- and high-strain properties and outline an approach for future work to synthesize and test these materials. In order to guarantee a sufficient density of tie molecules to ensure ductility, the molecular weight of the PE component should be at least ∼50 kg/mol (R0/d > 0.81 with d ≈ 30 nm), and in order to reach the critical amorphous layer concentration for rigidity, the total hPNbN content should be ∼20 wt %. Unfortunately, a diblock copolymer of these specifications (50 kg/mol PE, 12 kg/mol hPNbN) would microphase-separate in the melt: (χN)ODT ≈ 21 for this composition according to SCFT,39 while χN ≈ 31 from the measured interaction energy14 (X = 1.39 MPa, the interaction energy corresponding to (χN)ODT = 10.5 without fluctuation corrections for a symmetric diblock copolymer, to facilitate direct comparison with the SCFT prediction). However, the analogous triblock copolymer architecture, 25k−12k−25k PE− hPNbN−PE, is expected to exhibit a disordered melt: taking 2N to be the degree of polymerization, χN ≈ 16 with (χN)ODT ≈ 20 for this composition.40 A triblock architecture with PE end-blocks would be excellently suited to ensuring the presence of tie molecules, as the two PE blocks would sweep out a total distance equal to the full end-to-end length of the macromolecule; higher order multiblocks, or starblock architectures, or mixtures of all of the above would also possess the requisite tie molecules. Thus, judicious choices of chemistry, composition, thermal history, and macromolecular architecture would all play an important role in designing high-performance hPNbN−PE copolymers.
Figure 9. Strain at break vs the ratio of root-mean-square end-to-end distance (R0) of the PE block to the intercrystal repeat spacing (d) for the eight samples described in Tables 1 and 2. Low R0/d samples are brittle, while high R0/d samples are ductile.
In the case of the G-PE polymers, materials with high hPNbN contents needed to be synthesized with lower molecular weights in order to satisfy the criterion of melt miscibility. These lower molecular weights corresponded to PE blocks that spanned a shorter distance (smaller R0) and thus a reduction in the proportion of chains participating in at least two crystallites. While G-PE 15 and G-PE 21 prepared by slow cooling are on the left-hand, “brittle” side of the vertical line in Figure 9, they are both on the right-hand, “ductile” side when quenched (smaller d, hence larger R0/d). Therefore, we believe that there is no intrinsic trade-off between the increased stiffness imparted by glassy block incorporation and the loss of ductility, which is simply due to the reduction in molecular weight, as the two phenomena have different underlying causes; future work will investigate this point directly.
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CONCLUSIONS In this work, we demonstrated that the incorporation of a short block of high-Tg amorphous polymer into a majority-PE diblock architecture can dramatically alter several important mechanical properties relative to those of a PE homopolymer. Within a family of hPNbN-block-PE polymers, with molecular weights chosen to ensure homogeneity in the melt state, the yield stress and the tensile modulus were found to increase with higher glassy block contents, particularly when combined with a slow-cooling thermal history. For example, a diblock copolymer containing 15 wt % hPNbN, cooled slowly from the melt (∼1 °C/min), exhibits a yield stress and tensile modulus approximately double those of PE homopolymer. We attribute this increase in modulus and yield stress to a preferential enrichment of glassy hPNbN blocks near the crystal fold surfaces, resulting in regions of vitrified polymer that stiffen the amorphous layer and immobilize the crystal stems against slipping mechanisms under load. The samples with the highest yield stresses and tensile moduli tended to show reduced breaking strains, but it is important to point out that this need not reflect an intrinsic trade-off between stiffness and ductility in polymers with glassy amorphous layers; for example, nylon-66 and PBT are ductile under a typical Instron test. Rather, the loss of ductility in some of the materials described in this work is a consequence of the need to retain melt miscibility for polymers with increasing hPNbN content. As the hPNbN content was increased, the total molecular weight was concomitantly reduced to remain
ASSOCIATED CONTENT
S Supporting Information *
The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.7b02454. DSC traces, SAXS patterns, full stress−strain curves, and replicate tests for slow-cooled G-PE 15 (PDF)
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AUTHOR INFORMATION
Corresponding Author
*Tel +1 609 258 4691; fax +1 609 258 0211; e-mail register@ princeton.edu (R.A.R.). ORCID
Richard A. Register: 0000-0002-5223-4306 Notes
The authors declare no competing financial interest.
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ACKNOWLEDGMENTS This work was generously supported by the National Science Foundation, Polymers Program (DMR-1402180). The authors gratefully acknowledge Promerus LLC, especially Dr. Andrew Bell, for providing the norbornyl norbornene monomer.
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