Tough and Elastic Thermoplastic Organogels and Elastomers Made of

Article pubs.acs.org/Macromolecules

Tough and Elastic Thermoplastic Organogels and Elastomers Made of Semicrystalline Polyolefin-Based Block Copolymers Fanny Deplace,† Arthur K. Scholz,# Glenn H. Fredrickson,†,∥ Edward J. Kramer,†,#,∥,* Yong-Woo Shin,‡ Fumihiko Shimizu,‡ Feng Zuo,§ Lixia Rong,§ Benjamin S. Hsiao,§ and Geoffrey W. Coates⊥ †

Mitsubishi Chemical Center for Advanced Materials, Departments of #Materials and ∥Chemical Engineering, University of California, Santa Barbara, California 93106, United States ‡ Mitsubishi Science and Technology Research Center 1000, Kamoshida-cho, Aoba-ku, Yokohama 227-8502, Japan § Department of Chemistry, Stony Brook University, Stony Brook, New York 11974-3400, United States ⊥ Department of Chemistry and Chemical Biology, Baker Laboratory, Cornell University, Ithaca, New York 14853-1301, United States S Supporting Information *

ABSTRACT: Development of stereo- and regioselective living alkene polymerization catalysts has led to the capability to produce multiblock copolymers with semicrystalline syndiotactic polypropylene blocks and poly(ethylene-copropylene) rubbery blocks that have excellent elastomeric properties, both undiluted and as gels with as little as 8 wt % copolymer in mineral oil. During step cycle mechanical processing, the crystals can plastically deform and partially transform from lamellae into rod-like fibrils, which align along the tensile direction, giving rise to very large tangent moduli at large strains. The dimensions of these aligned fibrils can be established from small-angle X-ray scattering and used to construct models of the large strain elasticity, as well as its evolution with plastic strain, based on short fiber composite theory. Wide angle X-ray scattering experiments as well as FTIR measurements suggest that the oriented fibrils have the transplanar (form III) crystal structure. The large toughness of the dilute gels can be attributed to the plastic deformation of the crystalline component, which not only evens out the forces on the rubbery chains running from crystal to crystal but also provides a mechanical hysteresis that is effective in absorbing energy at the tip of a growing crack.

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INTRODUCTION Thermoplastic elastomers (TPE) have been the interest of numerous researchers worldwide since last three decades. They are one the most versatile polymers in today’s market and are used in many applications like automotive parts, household goods, adhesives, footwear, and medical devices. The great advantages of thermoplastic elastomers are their ability to be designed to meet specific requirements for each application, for example in terms of elasticity, extensibility, toughness, heat resistance, and ease of processing.1−11 Given the low cost and wide availability of olefinic monomers, much effort is now being directed toward the development of polyolefin TPEs. New developments in polyolefin synthesis give rise to polyolefinbased copolymers that combine the high melting crystallizable segments with elastomeric soft segments suitable for properties typical of the thermoplastic elastomers.6,12−24 The present work is focused on triblock copolymers comprised of syndiotactic polypropylene (sPP) for the hard end-blocks and random amorphous poly(ethylene-co-propylene) (EPR) copolymer blocks for the soft midblock.12 The importance of the distribution of the crystalline domains to achieve good elastomeric properties has also been revealed in © 2012 American Chemical Society

semicrystalline copolymers; better performances are obtained with block copolymers than with random copolymers.3,4,25 In random copolymers, the statistical distribution of crystallizable chain lengths results in a broad crystal size distribution and a low melting temperature, which limits their application at higher temperature. It has also been found that the structure is more oriented and the rubbery soft blocks are more easily stretched to higher strains in the block copolymer than in the random copolymer leading to a higher toughness and extensibility. Here the effect on the mechanical properties of the evolution of the crystalline structure and morphology of semicrystalline block copolymers under deformation is studied. Mechanical studies on semicrystalline materials clearly show that properties such as tensile strength, elongation at break and elasticity are associated with the reorganization and deformation of the hard domains and amorphous domains. The microscopic process of tensile deformation proceeds within several regimes.26−28 Received: April 21, 2012 Revised: June 4, 2012 Published: June 22, 2012 5604

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Scheme 1. Synthesis of sPP-EPR-sPP Triblock with (1) Being the Titanium Precursor Catalyst

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Deformation is accompanied by slip processes within the lamellae. At larger strain, when the stress acting on the crystallites reaches a value at which the crystalline blocks are no longer stable, new oriented crystallites are created whose assembly forms fibrils.26,29−35 In semicrystalline polymers where the crystallinity is low ( 10 GPa, the values of Ec are insensitive to the choice of Ef, we are left with one free parameter Gm since for a rubbery matrix Em ≈ 3Gm. The best fits are obtained with Gm = 2.4 MPa for sPP18_n, Gm = 2.8 MPa for sPP18_dg, and Gm = 2.0 MPa for sPP18_g. The experimental evolution of the maximal tangent modulus as well as the evolution obtained using the short fiber reinforced composite model are shown in Figure 21. The fits, which follow nicely the experimental data, confirm that the elastic deformation occurs in the rubbery matrix while the plastic deformation occurs in the fibrils. Toughness is the ability of a material to absorb energy before fracture. One way to measure toughness is by calculating the area under the stress/strain curve from a tensile test. This value is called the “material toughness”. From the area under the envelope of the curves shown in Figure 3, the material toughness is 44 MJ·m−3, 25 MJ·m−3 and 8 MJ·m−3 for sPP18_n, sPP18_dg, and sPP18_g, respectively. After correction by the dilution, the toughness of the gel is close to 100 MJ·m−3, which is much higher than the toughness of sPP18_n and sPP18_dg. We can also note that the material toughness of sPP18_n is higher than that of sPP18_dg due to the higher stress at break of the neat polymer. This higher stress has already been discussed and has been explained by a higher entanglement density of the sPP18_n. Another type of toughness is the “fracture energy” which is an indication of the amount of stress required to propagate a preexisting flaw. The fracture energy is used to evaluate the ability of a specimen containing a flaw to resist fracture. In the presence of a flaw such as notch, loading induces a triaxial tension stress state adjacent to the flaw. The material develops plastic strains as the yield stress is exceeded in the region near the crack tip. However, the amount of plastic deformation is

so that r = rc/exp(εHpC/2) and the average length should increase so that l = lc exp(εHpC) where εHpC is the plastic true strain after point C and rc and lc are the fibril radius and length at point C. The dashed lines in Figure 17, Figure 18, and Figure 19 are the predicted exponential evolutions. From the fits we can extract the values of rc and lc for each material. We obtained rc (sPP18_n) = 7.5 nm, rc(sPP18_dg) = 11 nm, and rc(sPP18_g) = 8.5 nm for the radii of the fibrils and lc(sPP18_n) = 100 nm, lc(sPP18_dg) = 40 nm, and lc(sPP18_g) = 206 nm for their lengths. The exponential expressions fit the experimental data within the experimental error and this confirms our assumption that the crystalline fibrils plastically deform as the sample is extended and become longer and thinner. The evolution of the aspect ratio of the crystalline fibrils is plotted as a function of the plastic true strain in Figure 20. In

Figure 20. Evolution of the aspect ratio of the crystalline fibrils as a function of the plastic true strain: squares, sPP18_n; triangles, sPP18_dg; circles, sPP18_g.

sPP18_g, the aspect ratio significantly increases from about 60 to 210. The rate at which the aspect ratio increases as a function of the plastic true strain is almost the same in the gel and in the dried gel. The increase of the aspect ratio as a function of the plastic true strain is slower in the case of the neat polymer. We showed that the crystalline fibrils plastically deform as they are extended and this plastic deformation might play a 5614

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Figure 21. Comparison of the evolution of the maximal tangent modulus as a function of the plastic true strain of the previous cycle (filled squares) with the evolution of the modulus predicted from the short fiber reinforced composite model (dotted line). (a) sPP18_n, (b) sPP18_dg, and (c) sPP18_g. Figure 22. Pictures taken during fracture tests: (a) sPP18_n, (b) sPP18_dg, and (c) sPP18_g.

restricted by the surrounding material, which remains elastic. When a material is prevented from deforming plastically, it fails in a brittle manner. On the other hand, in ductile materials, the crack moves slowly and is accompanied by a large amount of plastic deformation. The elastic blunting also plays an important role in soft materials as pointed out by Hui et al.65 In materials where the modulus is much less than the stress required to break bonds, the crack blunts before it can propagate. Within the blunted zone the material is highly stretched and therefore dissipative mechanisms can take place which cause further increases in toughness. In Figure 22 are shown photos of tensile specimens as they are deformed during fracture tests. The pure-shear geometry was used for sPP18_n and sPP18_dg. Since issues like slippage of the sample and early failure caused by the screws of the grips occurred during test on sPP18_g with the pure-shear geometry, the tear geometry has been preferred for this sample. Once the propagation of the crack has been initiated in the neat polymer specimen, it spread very rapidly until failure. The displacement at which the catastrophic failure of the sample occurred is very low (Figure 23a). On the other hand, crack tip blunting, an increase the crack tip radius, is observed in sPP18_dg and sPP18_g. A large portion of the stored elastic energy is consumed which can be quantified by the corresponding hysteresis energy (Figure 23c). The cyclic softening and hysteresis in our sPP-based elastomers and gels

are analogous to the so-called “Mullins effect” in nanoparticlefilled elastomers,66−69 even though the hysteresis and softening in the “Mullins” system must arise from completely different mechanisms. That the toughening effects are so large in both the nanoparticle-filled materials and in these semicrystalline block copolymers is due to the very large hysteretic energy dissipation that a material point experiences as it is loaded as a crack approaches and then is unloaded as the crack advances beyond that point. sPP18_dg showed a remarkable stress whitening zone (Figure 22b), which implies craze and microvoid formation in the EPR matrix, and cavitation between the sPP crystalline phase and the EPR matrix. The large extent of the stresswhitened zone at the crack tip suggests that plastic deformation processes were activated in a significantly large sample volume. The improved toughness from the neat polymer to the dried gel is associated with this stress-whitening.70 Earlier in this paper, we discussed the possibility of void formation during the drawing of our semicrystalline polymers and we concluded that the diffraction streak could not be attributed to the presence of voids. The whitening seen here clearly suggests the presence of voids. However, it is believed that most of these voids scatter X-rays outside of the SAXS detection limit and are not the origin of the diffraction streak. 5615

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before deformation in the deswollen amorphous phase explains its very high extensibility. This high ability to deform stems from its higher toughness compare to the neat polymer. The impressive toughness and extensibility of the gel is explained by the large amount of energy dissipated via the plastic deformation of the highly oriented crystalline fibrils. After the release of the stress, the polymer chains of the amorphous phase easily go back to their random coil conformation thanks to the very high dilution of the sPP crystals. This ability to go back to the initial conformation explains the excellent elasticity of the gel. The dimensions of the aligned fibrils can be established from small-angle X-ray scattering and are used to construct models of the large strain elasticity, as well as its evolution with plastic strain, based on short fiber composite theory. We demonstrated that our system can be considered as a short fiber reinforced composite where nearly all the plastic deformation occurs in the crystalline fibrils once they are formed and nearly all the elastic deformation occurs in the rubbery phase. The increase of the tangent modulus at high strains with the plastic true strain has been quantitatively explained by the increase of the aspect ratio of the plastically deformed crystalline fibrils. Finally, these results suggest that the combination of the very low crystallinity and the low amount of entanglements in the amorphous phase brings an excellent trade-off among the elasticity and the toughness of semicrystalline block copolymers. The key factor to achieve good elastomeric properties is high aspect ratio anchoring crystals with good isolation of the crystalline subunits within a rubbery matrix, similar to what is observed in natural elastomeric materials.

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Figure 23. Force versus displacement curves recorded during the fracture test on sPP18_n (a), sPP18_dg (b), and sPP18_g (c).

ASSOCIATED CONTENT

S Supporting Information *

Details of the correction for the volume fraction of the mineral oil, the determination of the viscous true stress, illustration of the effect of the time allowed to the sample to relax at zero load, details of the identification of the WAXS diffraction from the sPP crystal polymorphs, the determination of the fibril radius from the SAXS patterns and the details of the Ruland analysis of the SAXS streak from the fibrils giving the length and misorientation of the fibrils. This material is available free of charge via the Internet at http://pubs.acs.org/.

The failure of the sPP18_dg sample occurred in the grips of the tensile machine. The higher toughness of the dried gel compared to the neat polymer is also associated with its higher extensibility caused by a reduction in trapped entanglements in the rubbery phase, and by the compact conformation of the network strands with small end-to-end distances before deformation.42,48 The resistance to fracture of sPP18_g is also impressive. During the “tear test” the initial crack does not propagate. Instead, the arms of the sample clamped in the grips elongated along the tensile direction until failure which occurred in the grips. No stress whitening is observed during the fracture test of the gel. In this network, the energy dissipation can occur by two mechanisms: the plastic deformation of the crystals as in the dried gel and the viscous dissipation due to the friction of the small molecules of the mineral oil. No quantitative values of the fracture toughness are reported in this paper. But from the fact that the crack does not propagate we can safely say that sPP18_dg and sPP18_g are extremely tough. On the other hand, sPP18_n is “notch-brittle” and does not strongly resist crack propagation.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors gratefully acknowledge support from the Mitsubishi Chemical Center for Advanced Materials at UCSB. We made use of the UCSB MRL Central Facilities funded through the NSF MRSEC Program (DMR11-21053; a member of the NSF-funded Materials Research Facilities Network (www.mrfn.org)) and the Cornell Center for Materials Research Shared Experimental Facilities funded through the NSF MRSEC Program (DMR 0520404). Use of the National Synchrotron Light Source, Brookhaven National Laboratory, was supported by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences, under

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CONCLUSION Simultaneous SAXS/WAXS/tensile stretching experiments allowed us to investigate the origin of the toughness and elasticity of three materials prepared from the same triblock copolymer sPP-EPR-sPP. In the dried gel, the removal of a significant amount of trapped entanglements (topological constraints) and the compact conformation of the chains 5616

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Contract No. DE-AC02-98CH10886. We thank Hervé Elettro for his assistance in some of the experiments.

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