Mimicking (Linear) Low-Density Polyethylenes Using Modified

Article pubs.acs.org/Macromolecules

Mimicking (Linear) Low-Density Polyethylenes Using Modified Polymacrolactones Mark P. F. Pepels,† Ronald A. C. Koeken,† Sjoerd J. J. van der Linden,† Andreas Heise,†,‡ and Rob Duchateau*,†,§ †

Laboratory of Polymer Materials, Department of Chemical Engineering and Chemistry, Eindhoven University of Technology, P.O. Box 513, Eindhoven 5600 MB, The Netherlands ‡ School of Chemical Sciences, Dublin City University, Dublin 9, Ireland § SABIC T&I, STC-Geleen, SABIC Europe B.V., Urmonderbaan 22, 6160 AH Geleen, The Netherlands S Supporting Information *

ABSTRACT: This paper presents a new approach toward the introduction of both short- (SCB) and long-chain branching (LCB) in polyethylene-like polyesters via the ring-opening polymerization of macrolactones. Macrolactones containing an alkyl (S1) or alcohol (S2) branch were obtained using radical thiol−ene chemistry of ambrettolide (Amb). Kinetic studies revealed the need for an excess of thiol to achieve a high conversion of the double bond. Even though homopolymerization of the three monomers Amb, S1, and S2 revealed comparable reactivities, the molecular weight buildup during polymerization of S2 differs drastically from that of Amb and S1. Instead of the linear increase of Mn with conversion observed for Amb and S1, the molecular weight buildup for the ring-opening polymerization of S2 resembles that of a step-growth polymerizationslow buildup at low and moderate conversion followed by a rapid increase in molecular weight at high conversions. This disparity was attributed to the possibility of S2 to function as both an initiator and a monomer, leading to oligomers during the first part of the reaction that are subsequently connected to each other at the final stage of the reaction. Copolymerization of pentadecalactone (PDL) with various ratios of Amb, S1, and S2 in bulk led to the associated random copolymers containing double bonds, short-chain branches, and long-chain branches. The trans-double bonds in poly(PDL-coAmb) are included in the crystal lattice, leading to a slight decrease in the melting temperature, melting enthalpy and yield stress, while up to 20 double bonds/1000 backbone atoms the crystallinity and lamellar thickness remain similar to those of polypentadecalactone. In contrast, SCBs are fully excluded from the crystal lattice, leading to a more significant decrease in melting temperature and enthalpy as well as crystallinity and lamellar thickness with increasing branching density. The stiffness of these SCB-copolymers exponentially decreases as a function of branching content, effectively changing the mechanical behavior from semicrystalline to elastomeric. The LCB-containing polymers show an even larger linear decrease in melting temperature with increasing branching density than their SCB equivalents, likely due to the particular topology of the polymers consisting of a brush to a hyperbranched structure. However, a rapid decrease of molecular weight as was observed upon increasing the S2 content is also likely to play a role. The observed low molecular weight can be ascribed to both the fact that (macrocycles of) S2 can function as initiator, effectively increasing the amount of polymer chains, and the change of molecular weight buildup.

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INTRODUCTION

ring-opening metathesis polymerization (ROMP) of unsaturated renewable macrolactones with cyclic olefins6 and ringopening polymerization (ROP) of macrolactones.12−14 The ROP of macrolactones, viz. ω-pentadecalactone (PDL), can be catalyzed using enzymes,14,15 organic catalysts,13,16 and organometallic catalysts.12,17−20 Polypentadecalactone (PPDL) has been compared with linear low-density polyethylene (LLDPE) due to the similar stiffness and molecular structure, in which the ester groups are

Polymers derived from bio-based fatty acid feedstock have received considerable attention over the past years.1,2 Especially aliphatic long-chain polyesters (ALCPEs) are interesting, since they exhibit ductile properties resembling polyethylene.3,4 This is generally ascribed to the long methylene sequences in the backbone of the polymer, which dominate the crystallization behavior. Some small counits, like ester groups, can be incorporated in the crystals without significantly changing the crystallinity.5,6 Several routes have been explored for the synthesis of these types of polyesters using (combinations of) acyclic diene metathesis (ADMET),7,8 polycondensation,3,9,10 and thiol−ene chemistry11 of (modified) fatty acids as well as © XXXX American Chemical Society

Received: April 19, 2015 Revised: June 8, 2015

A

DOI: 10.1021/acs.macromol.5b00820 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules

Scheme 1. Modification of Ambrettolide Using Radical Thiol−Ene Chemistry To Obtain the Pentyl Branched (S1) and 1Hydroxyhexyl Branched (S2) Macrolactone; Subsequent Copolymerization of Amb, S1, and S2 with PDL Leads to the Corresponding Double Bond-Containing Copolymers (DB-Copolymers), Short-Chain-Branched Copolymers (SCBCopolymers), and Long-Chain-Branched Copolymers (LCB-Copolymers)

compared with short-chain branches (SCBs).13,21 However, in contrast to SCBs in LLDPE, the ester groups in PPDL are included in the crystal lattice and do not significantly hinder the crystallization, only leading to slightly larger values of the a and b parameter of the unit cell compared to high-density polyethylene (HDPE).22 Even though PPDL has a lower stiffness and yield stress than HDPE, the deformation behavior of the two is similar, manifesting itself in considerable localization of strain as a result of the high crystallinity. Conversely, the decrease in stiffness for LLDPE is related to the decrease in crystallinity due to the incapability of the shortchain branches to crystallize in the crystal lattice. 23 Furthermore, the accompanied decrease in lamellar thickness of LLDPE is the primary reason for the lower yield stress compared to HDPE.24 The presence of long-chain branching (LCB), which leads to an increased melt strength compared to HDPE and LLDPErequired for processes like film blowing and foamingforms the main characteristic of low-density polyethylene (LDPE).23 In order to introduce LLDPE- and LDPE-like topological features in ALCPEs and to elucidate their effect on the physical and mechanical properties of the products, SCBs or LCBs have to be introduced in the otherwise linear ALCPE molecular structure. A few reports describe the introduction of SCBs in ALCPEs, for example in which (derivatives of) ricinoleic acid are polymerized in combination with unbranched comonomers.25−28 However, the branch of this hydroxy-fatty acid is located adjacent to the alcohol group, which increases steric

hindrance and leads to a low polymerization reactivity (both polycondensation and ROP of the corresponding macrolactone). A similar effect of the branch position was also observed in the ROP of PDL with ε-decalactone, a sevenmembered ring with comparable ring strain as the highly reactive ε-caprolactone but with a butyl branch at the αmethylene carbon, for which the branch led to significant reduction in activity of the used phenoxy-imine zinc catalyst toward incorporation of the strainless PDL during the copolymerization. As a result, block copolymers rather than random copolymers were obtained that required additional transesterification using a different transesterification catalyst to actually produce random copolymers.29,30 On the other hand, long-chain branches have been introduced in polyesters using various strategies, e.g., polymerization from multifunctional initiators,31,32 postmodification with chain extenders,33 polymerization using difunctional lactone monomers,34 or using lactones containing an alcohol functionality.35 Long-chain branches are defined as branches of sufficient molecular weight to form entanglements. Similar to LDPE, branch-on-branch topologies are in principle also possible for ALCPEs and would result in a drastic increase of the melt strength. To our knowledge, however, no attempts to synthesize LDPE-like polyesters using the ring-opening copolymerization of macrolactones with hydroxyl-functionalized macrolactones have been reported. Here we describe the modification of macrolactones by introducing alkyl and alkanol branches using radical thiol−ene B


Article

Macromolecules chemistry. Subsequent copolymerization of PDL with ambrettolide or ambrettolide-based branched comonomers leads to alkenyl functionalities or short- and long-chain branching, respectively, in the molecular architecture of the corresponding copolymers. Furthermore, the relation between the molecular architecture, crystal structure, and mechanical properties of these polymers and copolymers containing double bonds is elucidated.

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RESULTS AND DISCUSSION Introduction of Alkyl and Alkanol Branches into Macrolactones Using Thiol−Ene Chemistry. The radical thiol−ene modification of ambrettolide (Amb) can be performed before or after polymerization. Postpolymerization modification has the advantage that excess or unreacted thiol can be easily removed by precipitation of the polymer in a proper solvent. However, dilute conditions are needed during the reactionespecially for high molecular weight polymers in order to obtain a homogeneous mixture of the reactants, which decreases the respective concentrations and results in a lower coupling yield.36−38 Controlling the amount of branching introduced in this way into (co)polymers is therefore difficult.36 Moreover, this method can only be used to introduce shortchain branches. In contrast, thiol−ene coupling of Amb and the thiol prior to polymerization allows the reaction to be performed in bulk due to the low viscosity of both reactants and the product. This approach allows the branching content in the polymer to be fine-tuned by copolymerizing the purified branched macrolactone with for example PDL in the desired ratio. Furthermore, it is expected that copolymerizing a macrolactone monomer containing an alcohol branch will lead to long-chain branching as a consequence of concurrent ring-opening polymerization and chain transfer, where the alcohol functionalities act as chain transfer agents. Hence, the prepolymerization modification of Amb using both 1pentanethiol (PT) and 6-mercapto-1-hexanol (MH) was investigated (Scheme 1). As previous attempts to modify Amb using thiol−Michael additions proved to be unsuccessful due to the lack of electron-withdrawing groups near the double bond of Amb and the low acidity of the used thiols,39 it was decided to use radical mediated thiol−ene chemistry to introduce the branches. A recent study showed that for globalide, a 16-membered macrolactone with an internal double bond similar to Amb, the coupling yield strongly depends on the type of thiol used and fine-tuning of the thiol:olefin ratios was required to obtain good yields.36 Preliminary enzymatic polymerization attempts with the substituted globalide produced only low molecular weight polymers.37 In order to gain more insight in the influence of the reactants on the thiol−ene coupling, bulk reactions of Amb with PT and MH using AIBN as radical source were performed at 80 °C. The 1H NMR spectra of the reaction mixtures revealed two distinct peaks, which can be used to determine the conversion, viz. the double bond signal of Amb at 5.35 ppm and the signal of one of the protons adjacent to the sulfide bond after addition of the thiol at 2.65 ppm (Figure 1 for S2 and Figure S1 in the Supporting Information for S1). The conversion was determined by normalization of the integral of the respective peaks (Ix) relative to the integral of the α-methylene protons at 4.12 ppm. Subsequently, the average of both conversions was taken:

Figure 1. 1H NMR spectrum of S2 after purification. The depicted chemical structure is one of the four possible regio- and stereoisomers.

⎡⎛ I ⎞ 2I ⎤ conversion = ⎢⎜1 − 5.35 ⎟ + 2.65 ⎥ /2 ⎢⎣⎝ I4.12 ⎠ I4.12 ⎥⎦

(1)

Since it was expected that Amb and its thiol adducts are difficult to separate, the focus of the reactions was on full Amb conversion. The influence of the radical concentration on the bulk reaction kinetics was investigated by increasing the ratio of AIBN/Amb stepwise from 0.01 to 0.1, resulting in only a minor increase in rate and final conversion (Figure 2a). Furthermore, an increase of the polarity by the addition of THF (30% v/v of the total reaction mixture) only has a negative effect on the reaction rate as a consequence of the dilution of all the components. Moreover, longer reaction times at 70 °C resulted in a lower final conversion compared to the reactions at 80 °C. An increase of the PT/Amb ratio from 1 to 5 led to a significant increase in rate, resulting in full conversion of Amb after 3 h for a PT/Amb = 3 (Figure 2b). For all reactions, the disappearance of the double bond signal at 5.35 ppm in 1H NMR was inversely proportional to the increase of the thiol adduct signal at 2.65 ppm, indicating that side reactions involving Amb were negligible. The addition reaction of MH to Amb shows similar reaction kinetics as PT to Amb (Figure S3, Supporting Information). Both S1 and S2 were synthesized in larger quantities (15 g) using the bulk reaction of Amb with a 3:1 ratio of thiol to Amb at 80 °C for 3 h. Subsequently, the reaction mixtures were subjected to vacuum at elevated temperatures to evaporate the excess thiol. The residual liquid product was analyzed for purity using 1H NMR. For the synthesis of both S1 and S2, negligible amounts of residual Amb and thiol were detected. The occurrence of multiple signals from the α-methylene protons at 4.12 ppm (Figure 1, signal a), compared to the triplet originally present in Amb, shows the formation of multiple regio- and stereoisomers, resulting from the thiol addition at either one or the other side of the double bond. Homopolymerization of Amb, S1, and S2. First, the homopolymerization of S1 and S2 was investigated. The aluminum salen complex (PDO-Al[salen] (1′); PDO = pentadecanoxy) was used as ROP catalyst. Therefore, monomer solutions ([Amb], [S1], [S2] = 1 M) in toluene were polymerized using 0.01 M of 1′ at 100 °C. No precipitation of polymer was observed during the reactions, showing that the polymerizations were completely homogeneous. As a consequence of the overlapping signals in the αmethylene regime in 1H NMR, the standard method for conversion determination of macrolactones using 1H NMR could not be used. C


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Figure 2. Conversion of Amb into S1 vs time for the bulk reaction of Amb with PT at 80 °C, using different amounts of AIBN (a) and using different ratios of Amb: PT ([AIBN] = 0.026 M) (b). The legend represents the molar ratios of the reactants.

Figure 3. Monomer conversion determined using low-MW-SEC vs time for Amb, S1, and S2 (a). Theoretical conversions taken from ref 12. Mn determined using SEC (RI detector) relative to PS standards vs the monomer conversion (b). The dotted lines represent a linear fit (for Amb and S1) and a fit using eq 2 (S2) (A = 5.8 kg mol−1, B = 0.32 kg mol−1).

Mn values stay very low up to high conversions. This can be explained by the fact that all monomers also contain an initiating group, viz. the alcohol side chain. Rapid exchange of the aluminum−alkoxide species with free alcohols will lead to a situation in which the monomers will actually act as initiators, ring-opening other S2 monomers. Besides the ability to grow themselves, the formed oligomers can also be incorporated in other growing chains by the reaction of the terminal S2 ring, a process that becomes more likely with increasing monomer conversion. Therefore, the buildup of the Mn of this polymer is a combination of both chain-growth and step-growth polymerization. This suggestion is strengthened by the excellent fit of a function of the form

In earlier work it was shown that the conversion of Amb could quantitatively be determined using low-MW-SEC analysis.40 Therefore, here we followed a similar approach for the polymerization of Amb, S1, and S2. Using this method, the conversion is defined as the disappearance of the monomer and not as the fraction of opened macrolactone rings. This discrepancy is especially important for S2, which can function both as monomer and as initiator. The conversion vs time profile of Amb follows the expected pseudo-first-order reaction path based on its kinetic parameters (Figure 3a).12 Both S1 and S2 follow a similar conversion profile as Amb. The fact that the aluminum salen catalyst is still capable of polymerizing S2 at similar rates compared to Amb means that the presence of excess alcohol does not deactivate the catalyst. The molecular weight (Mn) buildup of both Amb and S1 follows a linear trend relative to the conversion of the monomers, as is expected for this type of chain-growth polymerization. However, the ROP of S2 follows a completely different pattern (Figure 3b), in which

Mn =

A·X

chain growth

+

B 1−X

step growth

D

(2) DOI: 10.1021/acs.macromol.5b00820 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules

Figure 4. MALDI-ToF-MS spectrum of P(S1) (top) and P(S2) (bottom).

Table 1. Synthesized Copolymers with Their Thermal and Crystalline Propertiesa molecular weightc

mole ratio sample PPDL P(PDL-Amb P(PDL-Amb P(PDL-Amb P(PDL-Amb P(PDL-Amb P(Amb)

20-1) 10-1) 4-1) 2-1) 1-1)

P(PDL-S1 P(PDL-S1 P(PDL-S1 P(PDL-S1 P(PDL-S1 P(S1)g

20-1) 10-1) 4-1) 2-1) 1-1)

P(PDL-S2 P(PDL-S2 P(PDL-S2 P(PDL-S2 P(PDL-S2 P(S2)g

20-1) 10-1) 4-1) 2-1) 1-1)

ĐM

Tm (°C)

Tc (°C)

ΔHm (J g−1)

Xcryse

0.0 3.0 5.6 12.3 20.4 30.3 58.8

118 199 187 149 122 106 133

2.1 2.3 2.5 2.6 2.6 2.2 2.7

97.4 95.3 94.2 91.0 83.2 75.9 60.8

78.6 75.2 75.8 71.1 67.1 57.3 40.8

126.0 123.1 121.6 113.8 110.8 106.8 93.8

0.55 0.55 0.53 0.53 0.51 0.47 0.44

98 97 91 91 88 92 88

3.0 5.6 12.3 20.4 30.3 58.8

176 334 346 193 227 114

3.0 3.9 3.7 3.0 4.7 4.0

90.8 86.4 76.0 61.0 50.7 −h

75.6 68.4 57.2 42.1 29.1 −h

120.0 100.1 82.3 62.8 52.4 −h

0.50 0.46 0.38 0.25 0.16 n.d.

82 71 56 38 28 n.d.

3.0 5.6 12.3 20.4 30.3 58.8

130 85 43 20 10 3

5.5 5.3 4.7 5.0 4.3 2.0

88.7 84.2 73.0 59.4 42.9 −h

74.5 70.9 60.5 45.7 27.7 −h

125.5 119.4 101.2 79.3 51.7 −h

n.d. n.d. n.d. n.d. n.d. n.d.

n.d. n.d. n.d. n.d. n.d. n.d.

Amb

no. of defects

1 20 10 4 2 1 0 PDL 20 10 4 2 1 0 PDL 20 10 4 2 1 0

0 1 1 1 1 1 1 S1 1 1 1 1 1 1 S2 1 1 1 1 1 1

−1

thermal propertiesd

Mw (kg mol )

PDL

b

lcf (Å)

Polymers synthesized in bulk with [Mon]/[1]/[PDOH] = 500/1/0.5 at 100 °C for 16 h. bDefects/1000 backbone atoms, calculated using Xamb/ (XAmb/S1/S2·17 + XPDL·16)·1000, where Xi is the mole fraction of component i. Ester groups do not count as defects. cDetermined using HT-SEC. d Determined using DSC at a heating/cooling rate of 5 °C min−1. eCrystallinity determined using WAXD. fLamellar thickness calculated from the Long period (determined by SAXS) × crystallinity. gSynthesized using a ratio of [Mon]/[1]/[PDOH] = 100/1/1. hNo melting exotherm present. n.d. = not determined. a

in which A and B are fit factors and X is the conversion. The

implications with respect to the mechanical and rheological properties. MALDI-ToF-MS analysis provided additional insight into the molecular structure of the polymers of S1 (P(S1)) and S2 (P(S2)) (Figure 4), which should be considered nonquantitative. In the spectrum of P(S1), the main cyclic distribution, which has a repeating unit corresponding to the

practical implication for the ROP of S2 is that reactions need to be continued up to high conversions in order to obtain decently high molecular weights. It should also be kept in mind that the polymer microstructure will be somewhere between a brush and a hyper branched polymer, which will have severe E


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Macromolecules mass of S1 (356 g mol−1), can be identified, confirming that the carbon−sulfur bond is stable during polymerization and MALDI-ToF-MS analysis. The absence of PDOH-initiated chains in the MALDI-ToF-MS spectrum is the consequence of the high molecular weight of the sample. Consequently, the vast majority of the chains in the low molecular weight regime will be cyclics.40 As a result of the lower average molecular weight of P(S2), besides cyclics, a number of distributions corresponding to linear chains can be identified. All distributions have a repeating unit of 386 g mol−1 equivalent to the molecular weight of S2. Noncyclic chains were identified that were initiated with PDOH, 6-mercapto-1-hexanol, and its disulfide (Figure 4). The latter two originate from the incomplete removal of residual reactants after the thiol−ene radical reaction. Furthermore, the three homopolymers were analyzed for their thermal stability by performing thermogravimetric analysis. All polymers appeared to be stable up to temperatures of at least 200 °C (Supporting Information). Surprisingly, P(Amb) shows more weight loss than P(S1) and P(S2), which suggests that the double bond is thermally less stable than the carbon−sulfur bond. Copolymerizations of PDL with Amb, S1, and S2. The copolymerization of PDL with Amb, S1, and S2 will lead to polyethylene-like polyesters containing double bonds, shortchain branches (LLDPE mimic), and a combination of shortand long-chain branches (LDPE mimic), respectively. For the remainder of this paper double bonds (DBs), short-chain branches (SCBs), and long-chain branches (LCBs) can be referred to as defects, while ester groups will not be considered as such. Similar to the homopolymerizations, all copolymerization reactions were homogeneous (no precipitation of polymer). 1. Double-Bond-Containing Polymacrolactones. The copolymerization of Amb with PDL leads to polymacrolactones in which the fraction of DBs can be determined by the initial ratio of the two monomers. Copolymer of PDL or CL and Amb have been reported before,15,38,41 but only a single composition was investigated for its thermal properties. Therefore, various molar ratios of Amb and PDL were copolymerized in bulk at 100 °C, using an [Amb + PDL]/[1]/[PDOH] ratio of 500/1/ 0.5. These concentrations were selected based on previous experience in obtaining sufficiently high molecular weight for good mechanical properties (ductile deformation). The ratio of Amb to PDL was chosen to vary in a range where the resulting copolymer has 0−59 double bonds per 1000 backbone atoms (from PPDL to P(Amb)). Polymerization of the reaction mixtures led to DB-copolymers with high Mw values ranging from 119 to 199 kg mol−1 (Table 1). After precipitation and washing of the copolymers in methanol, 1H NMR analysis of the double bond protons at 5.35 ppm and α-methylene protons at 4.05 ppm revealed that the ratios of incorporated Amb to PDL were similar as the initial monomer ratios. Unlike for random copolymers of PDL and ε-caprolactone,13 13C NMR did not differentiate between homopolymer and copolymer diads, resulting from the very similar chemical shifts of the PDL−PDL and Amb−Amb diads. However, additional MALDI-ToF-MS analysis of the copolymers did prove the presence of copolymers of PDL and Amb (Figure 5). For P(PDL-Amb 2-1), a closer look at the isotope distribution of the chains consisting of five monomer units reveals the signals corresponding to each of the possible combinations of PDL and Amb (with the exception of 5 × Amb) for which the isotope

Figure 5. MALDI-ToF-MS spectrum of P(PDL-Amb 2-1) showing the distributions relating to cyclic copolymers of Amb and PDL (+ Na+).

patterns of 1 × Amb + 4 × PDL and 2 × Amb + 3 × PDL have the highest intensity. This is in agreement with the expectations for a sample consisting of twice as much PDL as Amb and indicates that fully random copolymers have been formed, as expected for the copolymerization of strainless rings.40 The influence of the double bond on the thermal properties of the DB-copolymers reveals a gradual decrease of the melting temperature from 97.4 °C for PPDL to 60.8 °C for P(Amb) (Figure 6a). The crystallization temperature shows a similar decreasing trend with increasing double bond content. Interestingly, the isomorphic copolymers stay crystalline over the whole composition range, and only a slight decrease in the melting enthalpy is observed (Figure 6c). WAXD and SAXS analyses provided further insight into the effect the double bonds have on the crystalline properties. The DB-copolymers with the exception of P(Amb) all have similar WAXD patterns with the typical main reflections at 2θ = 21.4° and 23.8°, corresponding to the 110 and 200 planes of the orthorhombic unit cell of PPDL (Supporting Information). Especially noteworthy is the fact that the calculated crystallinity from the WAXD spectra is constant for PPDL and DB-copolymers up to P(PDL-Amb 2-1) (Figure 6b,c). Furthermore, the lamellar thickness of these polymers, as obtained by SAXS, does not significantly decrease by the introduction of double bonds. Generally, two situations can occur when a counit is introduced in a crystalline polymer. Either this unit is not able to crystallize in the unit cell and is therefore excluded from the crystal lattice, which generally leads to a decrease in crystallinity while the unit cell dimensions are maintained (the typical example being butyl or hexyl branches in LLDPE), or the counits are included in the crystal lattice, typically leading to an expansion of the unit cell (often accompanied by a decrease in crystallinity).42 Ester functionalities form an exception to this rule. Typically, copolyesters show isomorphism over the whole composition range.5,6,43 It appears that the DB-containing copolymers also show this type of behavior, i.e. a similar unit cell and no decrease in crystallinity up to P(PDL-Amb 2-1), which leads to the conclusion that the double bonds are included in the crystal lattice without seriously affecting the crystalline structure. Similar to ester groups,5,6 the inclusion of the double bond functionality in the crystal lattice does lead to an enthalpic F


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Figure 6. DSC thermograms of the DB-copolymers (a). WAXD patterns of copolymers having 0, 30, and 59 double bonds/1000 backbone atoms (b). Influence of the amount of double bonds on the crystallinity (from WAXD), lamellar thickness (from SAXS), and melting enthalpy (from DSC) (c).

of the copolymers. P(Amb) shows a different WAXD pattern compared to the other DB-copolymers with distinct reflections at 19.21°, 21.22°, and 23.49°, indicating the presence of a second crystal structure. The polymorphous nature of P(Amb) was reported before,46 and it was suggested that this second crystal structure is comparable to the monoclinic structure of PE. Polymorphism of trans-polyalkenamers was reported to yield specific crystal structures in which the conformation of the double bond is out of the zigzag plane of the methylenes.47 It is likely that P(Amb) chains crystallize both with the double bond in and out of the zigzag plane, leading to the presence of both crystal structures as can be observed in Figure 6b. The introduction of PDL units in this structure probably forces the double bonds to crystallize in-plane, leading to the typical reflection of the PPDL crystal structure for DB-copolymers. Even though P(PDL-Amb 1-1), which consists of equimolar amounts of PDL and Amb, shows a similar WAXD pattern as the other DB-copolymers, a small decrease in crystallinity from 59% to 49% is observed. It is likely that the crystallization is impeded in this concentration range of double bonds. The mechanical behavior of the DB-copolymers was investigated by means of tensile testing on die-cut dumbbell samples. All polymers show ductile deformation behavior typical for PE-like polymers (Figure 8). The value of the yield stress is generally considered to be dependent on the intrinsic polymer properties in combination with its lamellar thickness, being independent of molecular weight and spherulitic superstructure.24,48 Except for the polymorphous P(Amb), all the DB-copolymers are considered as ALCPEs with double bond defects, since all polymers have identical lattice parameters as PPDL. The yield stress decreases linearly as a function of defect content from 17.9 MPa for PPDL to 10.0 MPa for P(PDL-Amb 1-1), which is far more than can be explained by the minor decrease in lamellar thickness, from 10.7 nm for PPDL to 9.7 nm for P(PDL-Amb 1-1). Therefore, the decrease is likely caused by the double bond defects in the crystal lattice, leading to plastic deformation of the crystals at lower stress levels. A more thorough investigation and explanation regarding the effect of the inclusion of defects in the crystal structure in relation to the mechanical properties fall out of the scope of this work. With regards to the current

penalty for the crystal lattice, resulting in a decrease in melting temperature and enthalpy. Since kinks in the molecular conformation caused by cisdouble bonds are known to disturb crystallization more significantly than their trans-counterparts,44 it is likely that the used ambrettolide is the trans-isomer and not the natural occurring cis-hexadec-7-enolide found in the vegetable oil of the ambrette seed. To elucidate on the molecular conformation of the polymers, a closer look at the double bond isomerism was taken using FTIR. The IR spectrum of PPDL shows the characteristic vibrations for the C−H bonds of the methylene group at 2922 and 2853 cm−1, the singlet at 1732 cm−1 characteristic for the carbonyl group, the fingerprint region which includes the vibrations corresponding to the methylene C−C bonds and the ester bond, and the CH2−rock vibration at 723 cm−1 (Figure 7). Introducing Amb into the polymer chain leads to an additional vibration at 965 cm−1 corresponding to the trans-CC vibration.45 The absence of a signal at 694 cm−1 indicates that both the used Amb and its polymers exclusively contain trans-double bonds, explaining the crystalline properties

Figure 7. FTIR spectra of PPDL, P(PDL-Amb 1-1), P(Amb), and Amb. G


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Figure 8. Stress vs strain curves of the DB-copolymers (PDL-Amb copolymers) (a). Yield stress vs the fraction of double bonds in the polymer (b).

Figure 9. DSC thermograms (vertically stacked) of the SCB-copolymers (PDL-S1 copolymers) (a). Melting temperatures for the SCB- and DBcopolymers as a function of units (either SCB or DB) present in the backbone (b).

of the melting temperature with increasing branching content from 97.4 °C for PPDL to 50.7 °C for P(PDL-S1 1-1) is far steeper than the decrease for the double bond content (Figure 9b). P(S1) does not exhibit a melting temperature and is fully amorphous over the whole temperature range. A glass transition temperature can be observed at −69 °C. It should be noted that P(S1), being the homopolymer of S1, is the only SCB-copolymer that has a regular distribution of the shortchain branches. In comparison, the precision polyethylene containing 30 butyl branches/1000 carbons has a melting temperature of 60.0 °C, which is slightly higher than its P(PDLS1 1-1) counterpart (being 50.7 °C).50 Compared to commercial LLDPE, a relatively narrow melting peak for the SCB-copolymers with low branching content (decyl) in some cases cocrystallize with the main polymer chain.49 Therefore, it is expected that the pentylsulfide branch, being comparable in length with a hexyl branch, will be excluded from the crystal lattice and will not participate in any crystallization events. Copolymerization of PDL and S1 in bulk at 100 °C leads to random short-chain branched-copolymers (SCB-copolymers) having a certain number of pentylsulfide branches, corresponding to the initial ratio of the two monomers (Table 1). The Mw of the polymers was sufficiently high to obtain plateau values, independent of molecular weight, for the physical properties.42 The melting behavior of the SCBcopolymers (Figure 9a) clearly show a similar trend as was observed for the DB-copolymers; however, the linear decrease H


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Figure 10. Lorentz-corrected normalized SAXS profile (a) and WAXD profile (b) of the SCB-copolymers. Crystallinity, lamellar thickness, and melting enthalpy as a function of the amount of pentyl sulfide branches/1000 backbone atoms (c).

Figure 11. TEM images of PPDL, P(PDL-S1 4-1), and P(PDL-S1 1-1) showing the decrease in lamellar thickness and crystallinity with increasing branch content.

Figure 12. Stress−strain curves for the DB-copolymers (a). E-modulus of the polymers as a function of branching content (b).

excluded from the crystal lattice. All semicrystalline SCBcopolymers were subjected to further investigation using X-ray analysis. Upon increased branching content, the Lorentzcorrected SAXS profiles show a broadening of the main reflection at 0.02−0.03 Å−1, corresponding to a broadening of the lamellar thickness distribution. This is the consequence of the exclusion of the branches from the crystal lattice and the fact that random copolymers of PDL and S1 have a distribution of lengths between the branches and thus crystallizable

multiple active sites in the Ziegler−Natta catalysts used to produce these copolymers. The DSC thermograms of the SCB-copolymers have welldefined crystallization exotherms showing a similar decreasing trend as a function of the branching content as the melting temperatures. Next to the decrease in melting temperature, the DSC results also show a clear decrease and simultaneous broadening of the melting enthalpy (Figure 9a), which is expected for a pseudo-random distribution of branches that are I


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increase of the modulus of the soft phase as well as an increase in hard filler content (the crystals). The combination of these two phenomena leads to a rapid increase in modulus. Below a certain branching content the lamellae become closely packed, which results in a more cocontinuous system, rapidly decreasing the mobility of the amorphous phase. In order to obtain further insight into the elastomeric behavior of P(PDL-S1 1-1), hysteresis tests were performed on this sample as well as on a reference rubber band sample, by consecutively stretching the samples to 50% strain and relaxing them back to 0 MPa, and subsequently repeating this process for 100% strain (Figure 13). For the first cycle up to 50% strain,

segments. All WAXD patterns show the same reflections typical for the PPDL orthorhombic unit cell at 21.4° and 23.8° (Figure 10b). Evidently, an increase in branching content leads to a decrease in crystallinity as a result of the decrease in crystallizable segments and concomitant increasing difficulty of chains to crystallize. In agreement with this, the lamellar thicknesses, obtained from SAXS, show a rapid decrease upon introduction of branching. Already 3 branches/1000 backbone atoms cause a decrease from 10.7 nm for PPDL to 8.2 nm for P(PDL-S1 20-1). The lamellar thickness decreases further down to 3.3 nm for 30 branches/1000 backbone atoms. For this case theoretical lamellae, having all 32 backbone atoms between the two branches, would have a length of 4 nm. Polymers generally have lamellar thicknesses that are significantly lower than the theoretical maximum due to kinetic effects during crystallization. The observation of relatively thick lamellae for the SCB-copolymers with high branching can be explained by the fact that the distribution of S1 over the PPDL chain is random, viz. the chain contains sections with more sequential PDL units than the ratio suggests. The lamellar morphology, as visualized using TEM, furthermore confirms the decrease in lamellar thickness (Figure 11). Additionally, the increase in amorphous fraction can clearly be seen by the decrease in contrast, especially for P(PDL-S1 11). Most important is the observation that the lamellae seem to be distributed homogeneously, since no large amorphous domains could be found. The consequence of introducing short-chain branches in PPDL on the mechanical properties becomes evident from the tensile test results (Figure 12). The yield stress for PPDL (17.9 MPa) rapidly decreases to 13.3 MPa for P(PDL-S1 20-1) and decreases further becoming nondistinguishable from P(PDL-S1 4-1) on, due to the absence of localized plastic deformation, similar to LLDPE. This is a direct result of the decrease in lamellar thickness, being the most important property in relation to the yield stress.48 This decrease in yield stress originates from a different structural property than the decrease observed for the DB-copolymers. The double bonds of the DBpolymers are positioned in both the amorphous phase and the crystalline lamellae. Even though the lamellar thickness is not dependent on the amount of double bonds, the lamellae themselves exhibit an increased mobility due to the defects introduced by the double bonds. This is in contrast to the SCBcopolymers, whose crystal structures do not contain branches and the lamellar thicknesses decrease by exclusion of the branches. Next to a decrease in yield stress, also an exponential decrease of the E-modulus is observed with increasing branch content (Figure 12b) from 480 to 10.8 MPa. Indeed, the deformation behavior changes from a typical semicrystalline polymer (PPDL) to an elastomer (P(PDL-S1 1-1)). This change accompanied by the decrease in modulus can be explained considering the two extremes. PPDL has closedpacked lamellae which results in a high stiffness and a clear yield stress connected to plastic flow of these crystals. P(PDLS1 1-1) has thinner lamellae which seem to be clearly separated by an amorphous phase. Therefore, similar to thermoplastic elastomers, the small crystalline regions act as physical crosslinks for the amorphous domains in between. Furthermore, the crystalline domains act as hard fillers, which increase the stiffness of the polymer. Considering these two aspects, the decrease in branching content has two consequences, viz. an increase in the content of physical cross-links resulting in an

Figure 13. Loading (solid lines) and relaxation (dashed lines) stress− strain curves for P(PDL-S1 1-1) and cross-linked natural rubber bands up to 50% and 100%.

P(PDL-S1 1-1) mostly shows elastic deformation having a residual strain of merely 5.4%, which is equal to a standard rubber band. The energy recovery, calculated using the ratio of the integral of the relaxation and the stretching curve, for P(PDL-S1 1-1) is also similar to that of a rubber band. Deformation and relaxation of P(PDL-S1 1-1) up to 100% leads to a less elastic response, with 16.7% residual strain and only 58.8% energy recovery. This suggests that upon increased strains (and thus stresses) part of the lamellae start to exhibit plastic deformation, leading to a reduction in elastic behavior and an increase in energy loss. These observations furthermore strengthen the hypothesis that at high branching content the lamellae function as physical cross-links for the amorphous matrix, and thus these polymers can be called thermoplastic elastomers. 3. LCB-Copolymers as LDPE-like Polymacrolactones. Copolymers of PDL and S2 are expected to result in polymacrolactones containing long-chain branches, shortchain branches, and branches on branches, depending on the incorporation of S2, making them the topological equivalents of LDPE. To investigate the influence of the LCB and possible SCB introduced by S2, PDL-S2 copolymers (LCB-copolymers) containing various monomer ratios were synthesized under similar conditions as reported for the DB- and SCB-copolymers (Table 1). The thermal properties of the LCB-copolymers show the same trend as the SCB-copolymers, i.e., a linear decrease in melting temperature from 97.4 °C for PPDL to 42.9 °C for P(PDL-S2 1-1) (Figure 14). Furthermore, the melting peaks J


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Figure 14. DSC thermograms (vertically stacked) of the LCB-copolymers (PDL-S2 copolymers) (a). Melting temperature of LCB-copolymers as a function of the amount of branches in the backbone (b).

become increasingly broader for LCB-copolymers containing more branching. Similar to P(S1), P(S2) is fully amorphous and has a Tg,P(S2) at −57 °C. The increase in Tg for P(S2) compared to P(S1) (Tg,P(S1) = −69 °C) can be explained by the increased amount of alcohol end-groups, leading to polar interaction restricting chain mobility. Compared to the SCBcopolymers, the melting temperatures of the LCB-copolymers decreases slightly more as a function of increasing branch content (Figure 14b). There are several factors responsible for this effect. First of all, the branching points of long-chain branches are (similar to short-chain-branching) excluded from the crystal lattice, but furthermore also impose a larger decrease in chain mobility, which leads to a decrease in the ability to crystallize. Second, the increased amount of alcohol end-groups for the LCB-copolymers changes the surface energy between the crystalline and amorphous phase, which leads to melting at lower temperatures compared to the SCB-copolymers. Furthermore, the reduction of the molecular weight (vide inf ra) also contributes to the decrease in melting temperature. The latter two effects are likely the cause for the observed decline in melting temperature, since the first effect is (partly) compensated by the decrease in molecular weight, which leads to higher chain mobility. During the copolymerization reactions, S2 functions not only as a monomer but also as an initiator. Since chain transfer of a growing chain and an alcohol is known to be orders of magnitude faster than polymerization,51 theoretically, during the initial phase of the polymerization mainly S2-initiated chains are formed. These chains subsequently act as macromonomers, which in following steps are incorporated into other polymer chains. Furthermore, also ring-closing reactions in which S2 units are involved will lead to similar structures. To confirm that S2 acts as both an initiator and a monomer, 13C NMR analysis was performed on the S2-rich copolymers (Figure 15). Since P(PDL-S2 2-1), P(PDL-S2 1-1), and P(S2) have high concentrations of end-groups, next to the ester αmethylene carbon (64.0−64.5 ppm) also the α-methylene carbons next to the alcohol end-groups have a distinct signal (62.7−63.2 ppm). S2 itself only has one signal in the latter regime, which originates from the MH side chain. After polymerization, P(S2) reveals two signals corresponding to αmethylene carbons next to an alcohol end-group: one signal at

Figure 15. 13C NMR of S2 and its copolymers with PDL.

62.9 ppm from the MH side chain (S2b) and one signal at 62.8 ppm from the alcohol originating from the lactone (S2a). Upon copolymerization with PDL a third signal appears at 63.0 ppm, resulting from the PDL end-group. Even though the integral values can only be considered semiquantitatively, the ratio of the PDL/S2a/S2b signals has values as expected from statistics, e.g., a ratio of 1/1/1 for P(PDL-S2 1-1). This shows both that alcohols of S2 are active in the ROP process and that S2 does not solely function as an initiator. Attempts to quantify the effect of long-chain branching on the mechanical properties were unsuccessful due to brittle failure for LCB-copolymers with more S2 incorporated than in P(PDL-S2 20-1) as a result of the rapid decrease in molecular weight with increasing branching content (Table 1). Since all polymerizations were allowed to react over prolonged times, it appears that this decrease in molecular weight is an intrinsic property and does not originate from a low conversion in combination with the typical Mn buildup observed for these S2 polymers (Figure 3b). Three factors were identified contributing to the decrease in molecular weight upon increased S2 incorporation into the polymer: (i) The apparent decrease in molecular weight caused by the composition of the polymera branched polymer generally has a smaller hydrodynamic volume compared to a linear polymer with the same molecular K


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containing up to 20 double bonds/1000 backbone atoms have a similar crystal structure, crystallinity, and lamellar thickness as PPDL, showing the isomorphic properties of this system. The effect of the decrease in melting temperature as well as melting enthalpy is a result of an enthalpic penalty introduced by inclusion of the double bonds in the crystal lattice. The defects introduced into the crystal lattice also lead to a decrease of the mechanical yield stress of the copolymers. For the SCB-copolymers the branches are excluded from the crystal lattice, leading to a decrease in crystallinity and lamellar thickness with increasing branching content, similar as observed in LLDPE with respect to HDPE. This is also reflected in the mechanical properties, for which the yield stress rapidly decreases as a consequence of the thinner lamellae. The modulus exhibits an exponential decrease, effectively changing the behavior from typical semicrystalline to elastomeric. Indeed, P(PDL-S1 1-1), having 30 branches/1000 backbone atoms, could be considered as a thermoplastic elastomer in which the crystalline lamellae function as physical cross-links for the amorphous phase. Copolymerization of PDL with S2 also resulted in polymers showing a linear decrease in melting temperature as a function of branch content. However, in these copolymerizations S2 does not only cause long-chain branching but also functions together with its cyclic co-oligomer as an additional source of initiator. This effect results in a decrease of molecular weight for increasing S2 concentrations. To summarize, these results demonstrate the implications, possibilities, and limitations of introducing double bonds, shortchain branches, and long-chain branches in the backbone bone of polyethylene-like polyesters synthesized via ROP on the polymer architecture, crystal structure, and mechanical properties.

weight. Furthermore, the increase in the amount of OH endgroups upon introduction of S2 might lead to a lower affinity of the polymer with the used solvent (trichlorobenzene) for SEC analysis, further decreasing the hydrodynamic volume of S2rich polymers. (ii) The second contribution is related to the residual 6-mercapto-1-hexanol and its disulfide, which are present as impurities in S2 acting as additional initiator. However, this should lead to a linear decrease in Mn versus the S2 (and thus impurity) concentration, which is not the case, suggesting that the reduction in molecular weight has an additional source. (iii) It is known that a significant part of the original monomer is present in the low molecular weight cyclic fraction of the macrolactone product.40 For PDL homopolymer, this leads to an additional low molecular weight distribution of cyclic polymer next to the high molecular weight linear chains. Even though this “bimodal” distribution considerably lowers the actual Mn, it does not significantly influence the molecular weight distribution of the linear chains, which is generally measured using SEC (Supporting Information, Figure S6a). However, for systems in which the formed cyclic polymers have side groups with OH-functionalities, additional initiation and transesterification occur at these moieties, leading to linear chains with the cyclic chains incorporated.35 In SEC, this results in the disappearance of the bimodality and shifts the main distribution to lower molecular weights (Supporting Information, Figure S6b). It furthermore should be taken into account that S2 can form two monomeric cyclesthe structure depicted in Figure 1 and the structure in which the hexyl sulfide is part of the monomeric ring and the branch is heptanoland therefore likely has a significantly higher cyclization constant than macrolactones without initiating side groups. In summary, the observed molecular weight decreases for three reasons: hydrodynamic volume effects, residual impurities, and additional initiation from the (increased amount of) cyclic fraction. Since the latter is an intrinsic property of this type of (poly)macrolactone, it is a factor which limits the molecular weight that can be achieved for copolymers containing S2 or equivalent macrolactones.

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EXPERIMENTAL SECTION

Reagents and Methods. All solvents and reagents were purchased from commercial sources (Sigma-Aldrich, BioSolve) unless stated otherwise. Toluene and THF were dried over an alumina column prior to use. Amb was kindly received from Symrise and freshly distilled from CaH2 under nitrogen prior to use. Irganox1010 and Irgafos 168 were kindly received from BASF. The aluminum Schiff base complex, [N,N′-bis(salicylidene-1,2-ethylenediimine)aluminum ethyl (1), was synthesized using a literature procedure.52 The corresponding pentadecanolate derivative 1′ was prepared in situ by reacting equimolar amounts of 1 and pentadecanol in toluene overnight. 1-Pentanethiol and 6-mercapto-1-hexanol were degassed by purging argon through the solution for 30 min. All reactions and preparations were either carried out in an MBraun MB-150 GI glovebox or using proper Schlenk techniques. 1 H NMR and 13C NMR spectra were recorded in 5 mm tubes on a Varian Mercury 400 MHz spectrometer equipped with an autosampler at ambient probe temperature in CDCl3. Chemical shifts are reported in ppm vs tetramethylsilane. A Polymer Laboratories PL XT-220 robotic sample handling system was used as autosampler. Low molecular weight size exclusion chromatography (low-MWSEC) for the conversion determination was measured on a system equipped with a Waters 1515 isocratic HPLC pump, a Waters 2707 autosampler, a Waters 2414 refractive index detector (35 °C), a Waters 2996 photodiode array detector, and a PSS SDV 5 m guard column followed by two SDV 5 m, 500 Å (8 × 300 mm) columns in series at 40 °C. Tetrahydrofuran (THF stabilized with BHT, Biosolve) with 1% v/v acetic acid was used as eluent at a flow rate of 1.0 mL min−1. The molecular weights were calculated against polystyrene standards (Polymer Laboratories, Mp = 580 Da up to Mp = 21 000 Da).

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CONCLUSION Radical thiol−ene chemistry was successfully applied to obtain macrolactones with a pentyl sulfide or 1-hexanol sulfide side group. The kinetic investigation of this reaction revealed that an excess of thiol was required to obtain full alkene conversion. Homopolymerization of the synthesized lactones S1 and S2 as well as Amb was successful and yielded similar reactivities for the three monomers. While Amb and S1 exhibit the typical linear molecular weight buildup versus conversion, polymerization of S2 results in low molecular weights at low conversions and steeply increased molecular weight at high conversions. This effect is caused by the fact that S2 can function as both an initiator and a monomer. Since chain transfer is much faster than propagation, the molecular weight buildup behaves as a combination of step-growth and chaingrowth polymerization. High molecular weight copolymers containing varying amounts of trans-double bonds and short-chain branches were successfully synthesized by copolymerization of PDL with Amb or S1, respectively. Both the introduction of DBs and SCBs into the backbone of the polymers leads to a linear decrease in melting temperature, which is steeper for SCBcopolymers than for the DB-copolymers. DB-copolymers L


Article

Macromolecules Size exclusion chromatography (SEC) for determination of the molecular weights of the Amb, S1, and S2 homopolymers was measured on a Waters Alliance system equipped with a Waters 2695 separation module, a Waters 2414 refractive index detector (35 °C), a Waters 2487 dual absorbance detector, and a PSS SDV 5 m guard column followed by two PSS SDV linearXL columns in series of 5 m, 500 Å (8 × 300 mm) at 40 °C. Tetrahydrofuran (THF stabilized with BHT, Biosolve) with 1% v/v acetic acid was used as eluent at a flow rate of 1.0 mL min−1. The molecular weights were calculated with respect to polystyrene standards (Polymer Laboratories, Mp = 580 Da up to Mp = 2.25 × 106 Da). Before SEC analysis was performed, the samples were filtered through a 0.2 μm PTFE filter (13 mm, PP housing, Alltech). High temperature size exclusion chromatography (HT-SEC) of the DB-, SCB-, and LCB-containing copolymers was performed at 160 °C using a Polymer Laboratories PLXT-20 Rapid GPC Polymer Analysis System (refractive index detector and viscosity detector) with three PLgel Olexis (7.5 × 300 mm, Polymer Laboratories) columns in series. 1,2,4-Trichlorobenzene was used as eluent at a flow rate of 1 mL min−1. The molecular weights were calculated with respect to polyethylene standards (Polymer Laboratories, Mp = 433 Da up to Mp = 1.08 × 106 Da). MALDI-ToF-MS analysis was performed on a Voyager DE-STR from Applied Biosystems equipped with a 337 nm nitrogen laser. An acceleration voltage of 25 kV was applied. Mass spectra of 1000 shots were accumulated. The polymer samples were dissolved in CHCl3 at a concentration of 1 mg mL−1. The matrix used was trans-2-[3-(4-tertbutylphenyl)-2-methyl-2-propenylidene]malononitrile (DCTB) (Fluka) and was dissolved in THF at a concentration of 40 mg mL−1. Solutions of matrix, sodium/potassium salt, and polymer were mixed in a volume ratio of 4:1:4, respectively. The mixed solution was hand-spotted on a stainless steel MALDI target and left to dry. The spectra were recorded in the reflection mode. Differential scanning calorimetry (DSC) analyses of the polymers were performed on a DSC Q100 from TA Instruments at a heating and cooling rate of 5 °C min−1 in the desired temperature range. The reported thermograms and crystallization as well as melting temperatures correspond to the peak temperatures of the first cooling run and second heating run, respectively. Wide- and small-angle X-ray scattering (WAXD/SAXS) measurements were performed on a Ganesha laboratory instrument equipped with a GeniX-Cu ultralow divergence source producing X-ray photons with a wavelength of 1.54 Å and a flux of 1 × 108 photons s−1. Scattering patterns were collected using a Pilatus 300 K silicon pixel detector with 487 × 619 pixels of 172 μm2 in size placed at a sampleto-detector distance of 480 and 1080 mm, respectively. The beam center and the q range were calibrated using the diffraction peaks of silver behenate. The mass fraction of the crystal phase, calculated using the integrals from peaks in the WAXD spectra of the crystalline phase divided by the total area, is used as the crystallinity. Since the exact lamellar thickness requires knowledge of the volume percentage of crystals and since the density of the amorphous and crystalline phase of the copolymers are unknown, the mass fraction of crystal (derived from WAXD) × the long period (derived from SAXS) has been used to calculate the average lamellar thickness. Therefore, the calculated values are affected by an absolute error, and they must only be considered for comparison between the different samples. TEM analysis was done using the following procedure. Before analysis, the samples (obtained using compression molding, vide inf ra) were trimmed at low T (−130 °C) and subsequently stained for 20 h. with a RuO4 solution prepared according to Montezinos et al.53 Ultrathin sections (70 nm) were obtained at −100 °C using a Leica Ultracut S/FCS microtome. The sections were put on a 200 mesh copper grid with a carbon support layer. The sections were examined in a Tecnai 20 transmission electron microscope, operated at 200 kV. Tensile tests were performed with a Zwick 100 tensile tester equipped with a 100 N load cell at room temperature. The test samples were prepared by compression molding the dry polymer into plates of 40 mm × 40 mm × 1 mm for 10 min at 170 °C, from which dog-bone-shaped tensile bars were stamped. A grip-to-grip separation

of 20 mm was used with a gauge length (LE) of 12.5 mm. The samples were prestressed to 1 N and then loaded until fracture with a constant cross-head speed of 10 mm min−1. For the hysteresis experiment, films having a width of 13 mm and thickness of 1 mm (20 mm grip-to-grip distance) were loaded and relaxed with 50 mm min−1. Radical Thiol−Ene Reactions. In a typical thiol−ene kinetic reaction, azobis(isobutyronitrile) (10 mg, 0.061 mmol, AIBN), Amb (1.54 g, 6.09 mmol), and PT (0.635 g, 6.09 mmol) were mixed under an inert atmosphere and distributed over seven small crimp cap vials. Subsequently, these vials were capped, taken out of the glovebox, and put in a carousel reactor at 80 °C (t = 0). At predetermined time intervals, vials were taken out of the reactor and uncapped, and aliquots of the crude reaction mixture were analyzed using 1H NMR. The average Amb conversion was determined using eq 1. For the synthesis of S1 and S2, a similar approach was taken. AIBN (0.11 g, 0.65 mmol), Amb (10.0 g, 39.7 mmol), and PT (12.4 g, 119 mmol)/MH (16.1 g, 120 mmol) were mixed together and allowed to react for 3 h at 80 °C. After this, the residual PT/MH was evaporated using vacuum distillation at 100 °C. Subsequently, molecular sieves were added to the monomers to remove trace amounts of water. General Polymerization Procedure. For the polymerization of Amb, S1, and S2, similar procedures were followed. Before polymerization, a catalyst stock solution was prepared by dissolving an equimolar amount of 1 and pentadecanol in toluene and stirring it for ±12 h at 100 °C to in situ produce 1′. In a typical polymerization, S2 (401 mg, 1.04 mmol), the catalyst stock solution (containing 0.0107 mmol of 1 and 0.0104 mmol of PDOH) and toluene (490 mg) were added to a glass crimp cap resulting in concentrations of [S2]: [1]:[PDOH] 1:0.01:0.01. The resulting clear solution was distributed over seven crimp cap vials, which were capped, taken out of the glovebox, and put in a carousel reactor at 100 °C (t = 0). At predetermined time intervals, vials were taken out of the reactor and uncapped, and aliquots were taken for low-MW SEC, SEC, and 1H NMR. Bulk Copolymerization Procedure. For the copolymers of PDL with Amb, S1, and S2, a similar synthetic procedure was followed having approximately a [monomers]:[catalyst]:[initiator] ratio of 500:1:0.5. In a typical bulk copolymerization S1 (1.34 g, 3.77 mmol), PDL (3.63 g, 15.08 mmol), 1 (13 mg, 0.041 mmol), and PDOH (4 mg, 0.016 mmol) were added to a 20 mL crimp-cap vial, which was capped, taken out of the glovebox, and stirred in an oil bath at 100 °C for 16 h. Subsequently, the vial was uncapped, and aliquots were taken for 1H NMR and HT-SEC. The remaining polymer mixture was dissolved in CHCl3, precipitated in cold methanol, and washed three times with methanol. The polymer was dried, and an acetone solution containing 0.5 wt % Irganox 1010 and 0.5 wt % Irgafos 168 relative to the polymer weight was added to the product. The acetone was allowed to evaporate (open air), after which the polymer was dried at 40 °C under vacuum.

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ASSOCIATED CONTENT

S Supporting Information *

Figures S1−S6 and Tables S1 and S2. The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.5b00820.

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AUTHOR INFORMATION

Corresponding Author

*Phone +31 40 247 4918; Fax +31 40 246 3966; e-mail r. [email protected] (R.D.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS Financial support by SABIC for this work is gratefully acknowledged. The authors thank Ilja Voets (Eindhoven University of Technology) for her help with X-ray analysis M


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and Anne Spoelstra (Eindhoven University of Technology) for her help with TEM analysis.

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