Thermoplastic Elastomers via Combined Crystallization and

Dec 29, 2015 - (15) They found that the tensile strength of an iPP-rPP-iPP-rPP-iPP pentablock .... Small-angle X-ray scattering (SAXS) patterns were c...
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Thermoplastic Elastomers via Combined Crystallization and Vitrification from Homogeneous Melts Adam B. Burns and Richard A. Register* Department of Chemical and Biological Engineering, Princeton University, Princeton, New Jersey 08544, United States S Supporting Information *

ABSTRACT: Block copolymers with crystalline, glassy, and rubbery blocks were synthesized by anionic polymerization of butadiene, styrene, and isoprene followed by chlorosilane coupling and hydrogenation. The performance of two pentablock copolymers, with the block sequence crystalline−glassy− rubbery−glassy−crystalline, as thermoplastic elastomers (TPEs) was evaluated against triblock copolymers having either crystalline or glassy end blocks. Judicious choices of block lengths yielded homogeneous melts for both pentablocks; consequently, the low-shear-rate viscosities of the pentablocks were more than 2 orders of magnitude lower than that of the glassy−rubbery−glassy triblock, which remained microphaseseparated in the melt. In the pentablocks, physical cross-linking was achieved by crystallization of the end blocks followed by vitrification of the adjacent glassy blocks, forming composite crystalline−glassy hard domains. All of the polymers studied exhibited desirable mechanical behavior for TPEs, including low Young’s modulus, high extensibility, and low permanent set. Increasing the glassy block fraction (at constant hard block content and molecular weight) systematically improved the mechanical performance by reducing the Young’s modulus and increasing the ultimate strength; however, the strain recovery was still limited by the crystalline component. Taken in the context of prior work on semicrystalline TPEs, this work highlights the influence of the crystalline morphology on the mechanical properties.

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INTRODUCTION

Truly emulating the processability of thermoplastics requires access to single-phase melts. In styrenic ABA triblocks, homogeneous melts can be attained by decreasing the molecular weight (and thus the segregation strength) such that the order−disorder transition temperature (TODT) becomes experimentally accessible. However, decreasing the segregation strength also progressively increases the degree of mixing between PS and the rubbery midblock at all temperatures, plasticizing the PS domains and degrading the mechanical performance at room temperature. Styrene−diene triblocks with thermotropic order−disorder transitions are substantially weaker than their higher molecular weight analogues. PS−PI−PS triblocks with nominally 20 wt % PS, and estimated4 TODT values of 190 and 140 °C, exhibited ultimate strengths which were 60% and 8% of the limiting value obtained at high molecular weights, respectively.5 This trade-off makes simultaneously attaining homogeneous melts and good use-temperature mechanical properties in styrenic TPEs infeasible. Using crystallization, rather than block incompatibility, as the driving force for physical cross-linking can provide pure hard domains (crystals) irrespective of the segregation strength. In such a system, the block chemistries can be chosen to yield

The development of living polymerization mechanisms, led by living anionic polymerization,1 has facilitated the meteoric rise of the field of block copolymers. One of the first and most commercially successful applications of block copolymers is that of ABA triblock copolymer thermoplastic elastomers (TPEs). In the canonical examplecommonly called styrenic TPEsthe A end blocks are glassy polystyrene (PS) and comprise roughly one-quarter of the material, while the majority B block is a rubbery polydiene (polyisoprene, PI, or polybutadiene, PBd) or amorphous hydrogenated polydiene (poly(ethylene-alt-propylene), PEP, or poly(ethylene-co-butylene), PEB).2 While traditional elastomers, such as vulcanized rubber, rely on a network of permanent chemical cross-links, styrenic TPEs employ what are known as physical cross-links derived from the incompatibility between the A and B blocks. At sufficiently high molecular weights this incompatibility causes the blocks to microphase-separate into a network of discrete PS domains embedded in a continuous rubbery matrix.2,3 Below their glass transition temperature (Tg), the PS domains anchor the rubbery chain ends, resulting in elastomeric properties such as high extensibility and low permanent set. These materials can be melt processed above the Tg of the PS domains; however, the persistence of microphase separation leads to highly viscous and elastic melts, ultimately making processing difficult and energy intensive.3 © 2015 American Chemical Society

Received: November 24, 2015 Revised: December 23, 2015 Published: December 29, 2015 269

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are limited by the response of crystals. Bates and co-workers studied two symmetric block copolymer architectures of the general form XBX, where X represents either alternating glassy−crystalline−glassy−crystalline−glassy26 or glassy−crystalline−glassy27,28 block sequences and B is rubbery PEP. In all cases the total hard block (X) content was relatively high (≈50%), leading to relatively high Young’s moduli and permanent sets, but homogeneous melts were obtained and the observed strain-hardening behavior was attributed to the effective reinforcement of the crystals by the adjacent glassy domains. A simpler ABCBA pentablock architecture was studied by Bishop and Register6 using ring-opening metathesis polymerization of various norbornene derivatives, where A is crystalline, B is glassy, and C is rubbery. Incorporating the glassy blocks had a strong beneficial impact on the mechanical properties; a triblock with crystalline end blocks displayed plastic behavior, while the pentablocks were soft elastomers with single-phase melts. This transformation was attributed to the efficacy of the glassy domains in reducing the lateral dimensions of the crystallites and providing mechanical reinforcement. The present report expands on this ABCBA pentablock architecture using anionic polymerization of common monomers: butadiene, styrene, and isoprene.

single-phase melts from which crystallization establishes a physically cross-linked network. Crystallinity can also add a degree of solvent resistance which is not attainable in allamorphous TPEs.6−8 Building off the extensive work on styrenic TPEs, early semicrystalline TPEs employed the same block architecture (ABA), but with crystalline A blocks. Anionic polymerization followed by catalytic hydrogenation was used to prepare ABA triblock copolymers where A is semicrystalline polyethylene (linear low-density polyethylene derived from hydrogenated low-vinyl polybutadiene, PE) and B is either PEP or PEB.7−13 Homogeneous melts and qualitatively elastomeric properties were indeed achieved; however, the crystalline blocks led to higher moduli and permanent sets than their amorphous-end block counterparts. A triblock with higher crystallinity linear polyethylene blocks, studied by Myers and Register, showed elastomeric properties at small strains but significant yielding above 350% strain.14 These adverse effects are a consequence of two factors: (1) crystallization tends to produce interconnected crystals of large lateral extent, which have a greater effect on the modulus than discrete domains, and (2) the crystals are prone to yielding via chain pullout and fragmentation. More complex block architectures have been devised to improve the performance of crystalline−rubbery TPEs. Hotta et al. synthesized block copolymers using crystalline isotactic polypropylene (iPP) blocks and regioirregular polypropylene (rPP) rubbery blocks which displayed both homogeneous melts and elastomeric behavior.15 They found that the tensile strength of an iPP-rPP-iPP-rPP-iPP pentablock copolymer was 4 times that of the analogous iPP-rPP-iPP triblock. The improvement was attributed to the central iPP block inhibiting chain pullout from the crystals; though interestingly, the pentablock architecture did not reduce the permanent set. The natural extension of this approach is to segmented or multiblock copolymers with alternating crystalline and amorphous blocks of the general form (AB)n. Commercial examples include thermoplastic poly(ether esters) and polyolefin block copolymers (OBCs). Poly(ether esters)16 are made by condensation of short (ca. 1000 g/mol) hard and soft segments, and in the case of poly(ether esters), homogeneous melts are attainable with relative ease.17 The high hard block contents necessary to produce sufficiently high melting points give rise to high tensile strength; however, the accompanying high modulus, low yield strain, and appreciable permanent set at modest strains limit their utility. OBCs comprise alternating blocks of ethylene/α-olefin copolymers with low (crystalline blocks) and high (rubbery amorphous blocks) α-olefin content.18−22 Weak block incompatibility allows access to homogeneous melts but imposes little restriction on the polyethylene crystallization. As a result, good ultimate properties can be attained, but the Young’s moduli are relatively high and strongly temperature-dependent. In a similar vein, Koo et al. found that well-defined (PE−PEP)n multiblock copolymers show elastomeric strain hardening only when crystallized from a microphase-separated melt or when n ≥ 10.23 Given that crystallization-driven phase behavior can afford drastically improved processability while glassy domains provide superior mechanical performance, efforts have been made to incorporate both crystalline and glassy components into a single material. The simplest such architecture is an ABC triblock copolymer with one glassy and one crystalline end block.10,24,25 However, microphase separation is still necessary to produce pure glassy domains, and the mechanical properties

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EXPERIMENTAL SECTION

Synthesis. Polymerizations were carried out at 60 °C under vacuum in glass reactors. The reactor, equipped with a Teflon stir bar, was flame-dried under vacuum before adding the initiator in a nitrogen-filled glovebox (MBraun UNILAB, ϕPE ≈ ϕPVCH. It is not obvious a priori whether the PE−PVCH pair (higher χ) or the PVCH−PEP pair (larger N) will have a stronger tendency to microphase-separate, so both must be considered to ensure miscibility. Treating the more strongly interacting PE−PVCH pair and considering the most stringent scenario of a symmetric (ϕPE = 271

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Scheme 1. Synthesis of Block Copolymers by Anionic Polymerization, Chlorosilane Coupling, and Catalytic Hydrogenation, Illustrated with a Pentablock Copolymera

a

Vinyl addition repeat units in the PBd and PI blocks are omitted for clarity.

ϕPVCH = 0.5) diblock indicates that homogeneous melts are possible when the diblock Mn < 13 kg/mol. Using this result and a composite hard block (PE and PVCH) content of 20%, a representative value for TPEs, limits the pentablock Mn to ca. 130 kg/mol. Turning to the corresponding PVCH−PEP pair (ϕPVCH ≈ 0.1), a pentablock with Mn < 160 kg/mol is predicted to be homogeneous in the melt. These simple calculations show that the phase behavior of the pentablocks is more sensitive to the PE−PVCH pair. The other important consideration is the glass transition temperature (Tg) of the glassy domains; in order to achieve the requisite mechanical strength, Tg must be well above the use temperature. The PVCH-domain Tg will decrease with decreasing Mn, both intrinsically37 and as a result of increased mixing with the low-Tg components. Working within these constraints, two pentablock copolymers were synthesized with PVCH block Mn ≥ 5 kg/mol and total Mn = 150 kg/mol. The first has a 1:1 (w/w) ratio of PE and PVCH, and the other has a 2:1 ratio; both have a total hard block content of 20 wt %. Polymers were prepared by sequential living anionic polymerization of butadiene, styrene, and isoprene followed by chlorosilane coupling and catalytic hydrogenation as outlined in Scheme 1. Gel permeation chromatography (GPC) traces after each polymerization step as well as the coupling step are shown in Figure 1. En route to the

the problem and in the worst-case scenario may lead to a population of chains with little or no PS block. Aromatic cosolvents were investigated as a way to improve the crossover rate. The PS content distribution was successfully sharpened (Figure 1 and Figure S1b) using a 1:1 (v/v) benzene/ cyclohexane mixture. The PS block distribution was assessed by GPC as described in the Supporting Information. Using 50 vol % benzene reduced the standard deviation in PS weight fraction by 30% and the spread by 25%. For consistency, all block polymerizations except the initial PBd blocks were conducted in 1:1 (v/v) cyclohexane/benzene regardless of the block sequence. The PBd blocks were polymerized in neat cyclohexane to minimize vinyl addition, thus maximizing the crystallinity and melting point of the eventual PE block. In cyclohexane the 1,2-content was 7.6−7.9% as measured by 1H NMR, corresponding to approximately 20 ethyl branches/1000 backbone carbons in the hydrogenated product. A PBd synthesized in 1:1 (v/v) cyclohexane/benzene had a 1,2content of 9.2 mol %, corresponding to 24 branches/1000 C if hydrogenated. If 1:1 (v/v) cyclohexane/benzene were used as the solvent throughout, the increase in branch content would be expected to reduce the PE melting point by approximately 5 °C.39,40 Sequential polymerization was followed by coupling with dimethyldichlorosilane (DMDCS). This approach was used because it reduces the number of monomer additions and polymerization steps and ensures that the block copolymers are symmetric. The rate of coupling was found to be a strong function of solvent polarity. In cyclohexane, the coupling reaction between polyisoprenyllithium and DMDCS takes days to go to completion, limiting the overall coupling efficiency. In the context of TPEs, uncoupled (monofunctional) material dilutes the bridging fraction, weakening the material. Studies on PS−PI−PS/PS−PI blends have elucidated the impact of monofunctional material on the mechanical properties.41−44 The modulus is minimally affected by uncoupled material44 so long as the uncoupled rubbery block Mn is sufficiently above the entanglement molecular weight (as is the case in this work, Me ≈ 1.5 kg/mol45 for PEP). The ultimate strength σu scales essentially linearly with the coupled fraction.42,44 Although the strain at break ϵb may increase with diblock content,43 the majority of this deformation is likely unrecoverable. Therefore, if a significant portion of uncoupled material remains, fractionation is necessary to remove the residual uncoupled portion if the best possible performance is desired. The living precursor polymers were coupled using a 0.1 M solution of DMDCS in THF (corresponding to 60 equiv of THF to Li). This method provides a robust route to nearly stoichiometric

Figure 1. Normalized gel permeation chromatography traces following each step in the synthesis of Pentablock-2, showing minimal first and second block termination and high coupling efficiency. True molecular weights were calculated from the apparent “polystyrene-equivalent” values by correcting for differences in hydrodynamic volume.29

pentablocks, sequential polymerization of the polybutadiene− polystyrene (PBd−PS) diblocks in cyclohexane yielded broadly distributed PS blocks (Figure S1a). This observation was attributed to the slow rate of crossover from polybutadienyllithium to styrene compared to the rate of styrene propagation.38 The need for short glassy blocks exacerbates 272

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Macromolecules Table 1. Molecular Characteristics of Block Copolymers polymer

Mwa (kg/mol)

Mnb (kg/mol)

Đc

fa,d

% 1,2-PB

% 3,4-PI

wPEf

PVCH−PEP−PVCH Pentablock-1 Pentablock-2 PE−PEP−PE

144 158 164 137

138 154 153 132

1.08 1.07 1.08 1.07

1.95 1.89e 1.95 1.95

7.7 7.8 7.9

7.7 7.1 7.2

0.09 0.13 0.20

wPVCHf 0.20 0.10 0.06

a

Determined by GPC with online light scattering of unsaturated precursors. bDetermined by GPC using dRI detection calibrated with narrowdistribution PS standards in conjunction with hydrodynamic equivalence ratios. cDetermined by GPC using dRI detection. dAverage functionality of coupled material calculated by f = Mw,coupled/Mw,arm, where Mw,arm is determined by light scattering on the uncoupled block copolymer. eMn,arm estimated from a partially coupled sample. fCalculated from 1H NMR after accounting for the appropriate number of hydrogen atoms per repeat unit.

presence of a strong peak (q* = 0.152 nm−1, d = 2π/q* = 41 nm) indicates that the material is microphase-separated; however, the lack of clear higher order peaks precludes an unambiguous assignment of the morphology. Microphase separation persists to at least 200 °C, as evidenced by the persistence of an intense primary peak. In fact, the sharpening of the primary peak and the emergence of a shoulder near q = 0.3 nm−1 (indicated by an arrow in Figure 2a) indicates that the order is improved at 200 °C.47 Based on form factor calculations31 (Figures S2 and S3), the patterns are consistent with cylindrical PVCH domains. Atomic force microscopy on an annealed, 120 nm thick PVCH−PEP−PVCH film (Figure 2b and Figure S4) revealed a disordered collection of circular PVCH domains, embedded in a rubbery matrix. The average domain size is consistent with the size approximated from q* at room temperature assuming a cylindrical morphology (see Supporting Information). Therefore, although the order is poor, the morphology of the PVCH domains is tentatively assigned as cylindrical. The SAXS pattern of PE−PEP−PE at room temperature (Figure 3) reveals a weak first-order peak (q* = 0.128 nm−1, d =

coupling (Figure 1). The residual uncoupled fraction was 5% or less in all cases, and fractionation was deemed unnecessary. Block copolymers synthesized as described above were catalytically hydrogenated over supported Pd0 (5 wt % on CaCO3). Hydrogenation is necessary to produce a crystalline block (PE) and also improves the thermo-oxidative stability of the polydiene blocks. Hydrogenation yields glassy PVCH and rubbery PEP blocks, which have homopolymer Tg values of approximately 14537 and −60 °C,46 respectively. The hydrogenation process was found to regenerate some uncoupled material through breakage of the central Si−C bonds. In PVCH−PEP−PVCH hydrogenation increased the amount of uncoupled material from 40 e−/nm3) to detect such a situation by SAXS. Therefore, the melt is indeed homogeneous. Pentablock-1 shows the same limiting behaviora structured solid and a homogeneous meltwith some added complexity above Tm (Figure 5). Below Tm, the solid-state

Having established the striking difference in phase behavior between triblocks with PVCH and PE end blocks, the SAXS patterns of the pentablock copolymers are examined next. Figure 4a shows the SAXS patterns for Pentablock-2 at room

Figure 5. Small-angle X-ray scattering patterns of Pentablock-1 at room temperature, 110 °C, and 200 °C.

SAXS pattern shows a low-q primary peak (q* = 0.142 nm−1, d* = 44 nm) followed by a broad tail extending to q ≈ 1 nm−1 with no discernible peak. Just above Tm, the primary peak and broad tail give way to a single peak (q = 0.263 nm−1, d = 24 nm). The breadth and temperature dependence (Figure S7) of this peak are consistent with correlation hole scattering, indicating that the melt is disordered but very near to the phase boundary at 110 °C.53,54 Using the same block chemistries, Bates and co-workers observed similar behavior in block copolymers with the general architecture X−PEP−X where X is either PVCH−PE−PVCH−PE−PVCH26 or PVCH−PE−PVCH.27,28 Interestingly, the positions of the correlation hole peaks identified by Bates et al. do correspond to the primary peak positions in the solid state below Tm, whereas this is not true of Pentablock-1. Thus, it appears that PE block crystallization in Pentablock-1 is not constrained to the length scale associated with the PVCH−PE/PEP as is apparently the case in the X−PEP−X polymers. Differential scanning calorimetry (DSC) offers some additional insight into the phase behavior and structure of these materials. DSC traces on dynamic cooling from the melt (Figure 6a) show bimodal crystallization peaks for both pentablocks compared to the usual sharp monomodal peak of PE−PEP−PE. This feature is interpreted as a two-stage crystallization process illustrated in Figure 7. Crystallization initially proceeds from a homogeneous melt. However, due to the connectivity of the blocks, crystallization of the PE blocks concentrates the adjacent PVCH segments in the vicinity of the crystallites, thereby imparting a degree of segregation between the PVCH and PEP blocks, which would otherwise be miscible.55 The crystallization-induced aggregation of the PVCH blocks, in turn, restricts the mobility of the PE segments, retarding crystal growth. The characteristic length scale separating these composite hard domains is on the order of 50 nm, as determined from the primary peaks in the SAXS patterns. As the undercooling is increased further, crystallization resumes within the confines of the PE−PVCH composite domains; the crystal periodicity, also determined

Figure 4. Small-angle X-ray scattering patterns at room temperature and 110 °C (a) and an atomic force microscopy phase image (b) of Pentablock-2.

temperature and just above Tm. At room temperature the most intense feature is the primary peak at low q (q* = 0.105 nm−1, d = 60 nm) followed by a broad hump centered near q = 0.6 nm−1. AFM phase images of Pentablock-2 (Figure 4b and Figure S5) show string-like arrangements of composite hard domains (dark features). The contrast in Figure 4b is inverted from the more typical scenario where hard domains appear bright (cf. Figure 2b). For the semicrystalline polymers the best images were obtained when the sample was minimally engaged by the AFM tip; in this attractive regime the less dissipative (i.e., hard) domains appear dark (see Supporting Information).52 The distance between nodules both along a string and between adjacent strings is consistent with the ≈60 nm long spacing from SAXS, confirming that the primary peak reflects the average spacing between composite hard domains. The string-like structure is attributed to the two-dimensional growth of PE crystallites which has been impeded by the formation of PVCH domains (dark nodules in the image). Pentablocks investigated by Bishop and Register showed evidence of similar structures.6 The broad hump in the SAXS pattern is attributed to the interior structure of the composite domains. Assuming this peak arises from alternating crystalline−amorphous PE lamellae, the crystal thickness (tc) was estimated by tc = ϕcd, where ϕc is the volume fraction of PE which is crystalline, calculated using the weight fraction crystallinity (wc,PE) from DSC. The estimated tc = 4.5 nm is in good agreement with 274

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block fraction was doubled at constant total hard block fraction). Interestingly, this effect appears to be less prominent the present case. In fact, the change in melting enthalpy (ΔHm) measured by DSC apparently scales nonmonotonically with glassy block content (Table 2). Pentablock-2 has the highest apparent ΔH m , which may reflect an extra enthalpic contribution from an underlying PVCH glass transition rather than an actual increase in wc,PE. Similarly, the actual reduction in wc,PE of Pentablock-1 may be partially masked by an underlying PVCH glass transition. It should be noted that although the Tg of the PVCH domains is not apparent in the DSC scans (due to small wPVCH and the hypothesized overlap with the melting endotherms in the pentablocks), the PEP Tg is in good agreement with the homopolymer value,28 indicating that the PVCH and PEP are substantially segregated. In all three semicrystalline polymers the crystallinity is sufficient to render them insoluble to at least 50 °C in cyclohexane, toluene, and tetrahydrofuran, good solvents for all three constituent blocks above Tm. It should also be noted that since Tm and wc are inherently limited by short chain branching, both were found to be largely insensitive to thermal history. The connection between the phase behavior and melt processability was explored using steady-shear and smallamplitude oscillatory shear (SAOS) rheometry at 180 °C. Figure 8 shows viscosity curves for each of the linear polymers constructed from the steady-shear (η) and the complex viscosity (η*) using the Cox−Merz rule.58 The pentablocks and PE−PEP−PE, all of which are homogeneous at 180 °C, show viscosities in the range of (1−3) × 104 Pa·s with clear zero-shear-rate plateaus. The incremental increases in viscosity from PE−PEP−PE to Pentablock-2 to Pentablock-1 reflect the increase in PVCH content. On the other hand, at low shear rates the viscosity of PVCH−PEP−PVCH is more than 2 orders of magnitude larger and shear thinning is observed even at the lowest rates; both observations are consequences of the microphase-separated structure. When probed by SAOS (Figure S8), PVCH−PEP−PVCH is the only linear polymer which does not display terminal behavior (elastic modulus G′ ∼ ω2 and loss modulus G″ ∼ ω1 at low ω). Thus, the pentablocks and PE−PEP−PE display truly liquid-like melts while PVCH− PEP−PVCH remains elastic even at the lowest frequencies tested. Mechanical Properties. With the phase behavior of the four linear polymers established, their performance as elastomers was assessed in uniaxial extension. Figure 9 shows representative stress−strain curves for each of the four polymers. All four polymers display elastomeric behavior characterized by low Young’s modulus E (or modulus ratio E/E0N, where E0N is the plateau modulus of the un-cross-linked rubbery component, EN0 = 3.45 MPa for PEP45), strain

Figure 6. Differential scanning calorimetry thermograms of the three semicrystalline polymers on cooling (a) and all four linear polymers on heating (b) at 10 °C/min.

by SAXS, is approximately 10 nm. In Pentablock-2 the two crystallization peaks are heavily overlapped (peak-to-peak separation of only 4 °C), indicating that the short PVCH blocks provide minimal interference. On the other hand, Pentablock-1 shows two clearly resolved peaks separated by some 24 °C. The large additional undercooling required to fully crystallize Pentablock-1 is ascribed to the proximity to the order−disorder transition, which enhances the segregation between PVCH and PEP, producing higher-Tg domains richer in PVCH. The effect of the PVCH blocks on the PE crystals can also be inferred from the degrees of crystallinity (wc) as well as from the melting points (Tm, Figure 6b and Table 2); the latter depend sensitively on crystal thickness.39,40 In the absence of other constraints, the crystallinity (wc ≈ 0.4) and melting point (Tm ≈ 100−105 °C) of the PE blocks are fundamentally limited by the randomly distributed ethyl branches arising from 1,2-addition in the precursor PBd.39,40,56,57 Pentablock-2 has a melting point 2 °C below that of PE−PEP−PE, meaning that the crystal dimensions are slightly reduced. A more drastic change in crystallite thickness is inferred from the peak melting temperature of Pentablock-1, which is 10 °C below that of PE− PEP−PE. Using the same architecture with different block chemistries, Bishop and Register6 observed a drastic reduction in wc of the crystalline end blocks when the glassy block fraction was increased (from wc = 0.54 to 0.15 when the glassy

Figure 7. Illustration of the phase behavior of the pentablock copolymers. Blue blocks represent PE, thick green blocks are PVCH, and PEP is red. On the left, the material is homogeneous above the melting point (Tm); in the center, crystallization begins to physically cross-link the material below the crystallization temperature (Tc) but is hindered by aggregation of the adjacent PVCH blocks; on the right, further crystallization and glassy block vitrification occurs upon further cooling. 275

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Macromolecules Table 2. Thermal and Mechanical Characteristics of Block Copolymers

a

polymer

Tg,PEP (°C)

PVCH−PEP−PVCH Pentablock-1 Pentablock-2 PE−PEP−PE

−58 −59 −59 −59

Tm (°C)

ΔHm (J/g)

91 99 101

8.3 17.6 21.4

wc,PEa

E (MPa)

σu (MPa)

εb (%)

0.33 0.49 0.39

7.0 6.7 8.4 10.2

9.4b/11.0c 6.9/8.4 7.5/7.7 6.4/6.4

830b/840c 660/740 1000/1070 560/570

Apparent weight fraction crystallinity of the PE blocks. bAverage value of at least three specimens. cValues from the specimen with the highest σu.

a soft elastomer, the effect is modest in this case; the Young’s modulus ratio can be 10 or more6,14 in polymers with higher crystallinity end blocks and comparable block fractions. Incorporating PVCH blocks simultaneously reduces E and increases σu (Table 2), both of which are desirable effects, with longer blocks providing greater improvement. The reduction in modulus is attributed to the effect of the PVCH on the continuity of the hard domains. In the pentablock copolymers aggregation of the PVCH blocks near the growing crystallites limits the lateral dimensions of the crystals. This produces more-discrete hard domains and consequently a lower modulus. A similar argument can be made for the effect of the PVCH content on σu. The presence of the PVCH layer surrounding the PE crystallites inhibits processes like chain pullout and crystal fragmentation which ultimately lead to failure. It should be noted that although the PE end blocks alone qualitatively provide the desired mechanical characteristics, it is impossible to simultaneously reduce E and increase σu without incorporating the adjacent PVCH block (Figure S10).7,8,11 Dynamic mechanical thermal analysis (DMTA) was used to probe the temperature dependence of the small-strain modulus (Figure 10). The upper use temperature of the semicrystalline TPEs is significantly less than that of PVCH−PEP−PVCH, primarily because PE melts some 40 °C below the Tg of the pure PVCH domains. In addition, the moduli of the PEcontaining polymers are more temperature-dependent. This is due, in part, to thermally activated screw dislocations, commonly referred to as the α relaxation.64,65 PE−PEP−PE shows peaks in the storage modulus (E″) and tan δ near 45 °C, consistent with the α relaxation observed in a hydrogenated polybutadiene with similar branch content.64 In the pentablocks, both the intensity and temperature (Tα) of the α relaxation peak decrease with increasing PVCH content (decreasing PE content). The decrease in intensity is attributed to the reduced overall degree of crystallinity, while the reduction in Tα is a result of the reduction in crystal thickness inferred from Tm, measured by DSC.64,65 Unfortunately, the presence of the PVCH domains does not significantly inhibit the α relaxation process. Instead, the reduction in PE crystal thickness dominates, causing Tα to shift to lower temperatures and reducing the storage modulus (E′). In spite of the reduction in Tα, all three semicrystalline TPEs perform relatively well at elevated temperatures. When heated from 25 to 50 °C, PE−PEP−PE and Pentablock-2 show a reduction in E′ of 25%, while Pentablock-1 exhibits only a 16% reduction; at 75 °C, E′ is reduced by 50% in all three polymers. For comparison, the softest polyethylene-based olefin block copolymer studied by Wang et al.18 (OBC H18, E′ = 12 MPa at 25 °C) softens by 50% at 50 °C and 70% at 75 °C. Stiffer OBCs (E′ = 36−520 MPa) fare better but still show stronger temperature dependencies than the TPEs reported herein, softening by 30% at 50 °C and 60% at 75 °C.

Figure 8. Viscosity curves of linear polymers; filled symbols indicate steady-shear data plotted against the shear rate (γ̇) while open symbols indicate complex viscosity plotted against angular frequency (ω).

Figure 9. Representative uniaxial stress−strain curves for each linear polymer. Each specimen of Pentablock-2 tested (3 total) showed a yield point near 800% strain.

hardening, and large strain at break ϵb (Table 2). In the case of PVCH−PEP−PVCH, microphase separation yields welldispersed PVCH domains leading to a modest increase in modulus (E/E0N = 2.0 ± 0.7). The large variability in the modulus of PVCH−PEP−PVCH is consistent with cylindrical PVCH microdomains, where the stress−strain behavior is sensitive to the orientation of the cylinders with respect to the stretching direction.59−61 Melt-pressed films for tensile tests were strongly birefringent, indicative of preferential orientation,62 a well-documented occurrence for strongly segregated melts with anisotropic domains subjected to flow.60,63 Of the six specimens tested (Figure S9), one specimen had a modulus of 11.8 MPa, nearly twice the average of the other five (6.0 ± 0.8 MPa). The high modulus and prominent yield point near 15% strain observed in this sample are consistent with a cylindrical morphology where the cylinder axes are principally coincident with the stretching direction (see Supporting Information for further discussion).59−61 Excluding this specimen due to its apparent alignment (corroborated by a Q test with 95% confidence), the remaining five specimens show the lowest modulus ratios (E/E0N = 1.7 ± 0.2) of all four TPEs. At the other extreme, the crystallization-driven lamellar morphology of PE−PEP−PE is responsible for the larger modulus ratio (E/E0N = 2.9 ± 0.1). While this is undesirable for 276

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Article

Macromolecules

crystallites. At low strains (ϵa < 100%), where the reorientation of hard domains62 causes the recovery to be relatively poor after the first cycle, the pentablocks do outperform PE−PEP−PE. The improvement is again attributed to the limited lateral extent of the composite PE−PVCH domains. However, at large strains (ϵa ≥ 100%), where the recovery response is flat, the PVCH blocks in the pentablocks provide no significant benefit; the recovery is dominated by the PE blocks. Each specimen of Pentablock-2 tested exhibited a yield point (ϵy) near 800% strain as seen in Figure 9. After failure (ϵa = ϵb ≈ 1000%), an estimated 90% of the deformation was recovered. Assuming that 95% of the deformation below ϵy is recovered indicates that only 70% of the postyield deformation is recovered. The location of this feature distinguishes it from the well-studied yielding mechanisms of polyethylene homopolymer,66 which occur at much lower strains. However, similar behavior was observed by Mohajer et al. in a triblock copolymer with 18 wt % PE end blocks and a PEP midblock.11 In addition, a triblock copolymer with linear polyethylene end blocks studied by Myers and Register14 yielded near 370% strain. Thus, the high-strain yield point appears to stem from irreversible deformation of the PE blocks. The influence of the identity of the crystalline end blocks on the mechanical properties is borne out by comparing the results presented here with those of previous reports. An interesting comparison can be made between triblocks with polyethylene end blocks of different crystallinities. Polyethylene (PE) blocks derived from hydrogenated polybutadiene have been studied extensively, including here.7−13 In these systems, relatively low modulus ratios, strain hardening, and good recovery can be achieved with the appropriate block fractions. Increasing the PE fraction in these polymers makes the materials stiffer but does not qualitatively change the shape of the stress−strain curve until ca. 50 wt % PE.7,8,11,13 In contrast, truly linear polyethylene (LPE) end blockswhich are capable of producing much larger (both in terms of thickness and lateral dimensions),67 more interconnected crystals and higher degrees of crystallinityshow qualitatively different behavior. An LPE− hPHN−LPE triblock (where hPHN is rubbery hydrogenated polyhexylnorbornene) with 20 wt % LPE exhibited a larger modulus ratio (E/E0N = 9.1)68 and a prominent yield point near 370% strain.14 Prior to yielding, the triblock exhibited strain hardening and good recovery, consistent with TPE behavior, but it is likely that the postyield deformation is largely unrecoverable. Espinosa et al. investigated an LPE−polyisobutylene−LPE triblock with 40 wt % LPE and found E/E0N > 100, characteristic of a plastic rather than an elastomer.69 A triblock copolymer with hydrogenated polynorbornene (hPN) crystalline end blocks (20 wt %) and an hPHN rubbery midblock showed similar plastic-like behavior with E/E0N = 20, distinct yielding, and correspondingly high permanent set.6 Similar to LPE, hPN is a highly crystalline polymer.70,71 Thus, the properties of linear ABA triblock copolymers with crystalline A blocks cluster into two general classes based on the properties of the end blocks. Inherently low-crystallinity end blocks produce crystallites of limited size and connectivity, which in turn confer elastomeric properties even when the crystalline block content approaches 50%. On the other hand, highly crystalline end blocks produce thicker, more-interconnected crystals, which lead to high moduli and yielding. It should be emphasized that the overall degree of crystallinity is not the controlling parameter. This is most clearly seen by comparing PE-based triblocks11,13 with the LPE-based triblock

Figure 10. Storage modulus E′ (a), loss modulus E″ (b), and tan δ (c) of block copolymer TPEs measured by dynamic mechanical analysis. Insets provide an expanded view of the rubbery region.

Another key property of elastomers is strain recovery. To test the recovery, samples were serially strained and relaxed, increasing the applied strain in each successive cycle. The results, cast in terms of cumulative recovery (R = (1 − ϵs/ϵa) × 100%, where ϵa is the applied strain and ϵs is the total cumulative residual strain), are shown in Figure 11. All four

Figure 11. Strain recovery of all four linear polymers measured 5 min after returning the cross-head to the initial position. For Pentablock-2, the points represent an average of three measurements and error bars represent one standard deviation.

polymers showed elastomeric recovery (>92%) up to at least 400% strain. Strikingly, the three polymers containing crystalline end blocks show similar recoveries, which lie systematically below those of PVCH−PEP−PVCH. The comparatively poorer performance of the materials with PE end blocks is attributed to yielding and fragmentation of the 277

DOI: 10.1021/acs.macromol.5b02546 Macromolecules 2016, 49, 269−279

Article

Macromolecules of Myers and Register.14 At comparable overall degrees of crystallinity (ca. 15 wt %), the PE-based triblocks remain elastomeric to higher strains and have lower modulus enhancement. The key aspect is thus the morphology especially the connectivityof the crystals. With this in mind, it becomes clear why the improvements afforded by the pentablock architecture reported here are less striking than those found by Bishop and Register,6 who used highly crystalline hPN end blocks, the morphology and mechanical response of which were drastically altered by the presence of adjacent glassy blocks. In the present case, the size and connectivity of the crystals are intrinsically limited by the low crystallinity resulting from the short-chain branching, so the addition of the glassy blocks has a comparatively smaller effect on the morphology and performance.

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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.). 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 thank Professor Robert K. Prud’homme and Brian K. Wilson for use of the rheometer.

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CONCLUSIONS Anionic polymerization followed by chlorosilane coupling and catalytic hydrogenation was used to synthesize a series of linear block copolymers. A pentablock architecture, crystalline− glassy−rubbery−glassy−crystalline, was compared against triblock copolymers having either crystalline or glassy end blocks at constant hard block content (20 wt %) and overall molecular weight (Mn ≈ 150 g/mol). The block lengths used in the pentablocks yielded homogeneous melts above the melting point of the crystalline blocks. Compared to the microphaseseparated melt of the all-amorphous PVCH−PEP−PVCH, the viscosity of the pentablocks was reduced by over 2 orders of magnitude. In the materials with PE end blocks, a solid-state network structure was established by crystallization. In the pentablocks, aggregation of the glassy blocks, subsequent to crystallization, was effective in limiting the crystal growth and provided mechanical reinforcement. All of the polymers investigated showed low Young’s moduli, strain hardening, large extensibilities, and good recovery at room temperature, characteristic of elastomeric behavior. Increasing the glassy block content improved both the modulus (reduced) and ultimate strength (increased). Despite these improvements, the glassy blocks had little effect on the strain recovery of the semicrystalline TPEs, which lie systematically below that of the all-amorphous polymer. The results presented here, in conjunction with the report of Bishop and Register,6 highlight the importance of the crystal morphologyspecifically the size and connectivity of the crystalson the mechanical properties. These results also demonstrate the general utility of the pentablock architecturethe simplest architecture capable of producing TPEs which harness the advantages of both crystalline and glassy blocks.

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for all six PVCH−PEP−PVCH specimens tested; and stress−strain curves for PE−PEP−PE triblocks (PDF)

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.5b02546. Procedure for determining molecular weight; molecular weight and styrene distributions in butadiene−styrene diblocks; form factor calculations for PVCH−PEP− PVCH; additional atomic force microscopy images of PVCH−PEP−PVCH, PE−PEP−PE, and pentablocks; additional SAXS patterns of Pentablock-1; smallamplitude oscillatory shear data; stress−strain curves 278

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Macromolecules

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