A Comparison of the Rheological and Mechanical Properties of

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A Comparison of the Rheological and Mechanical Properties of Isotactic, Syndiotactic, and Heterotactic Poly(lactide) Love-Ese Chile,†,‡ Parisa Mehrkhodavandi,*,† and Savvas G. Hatzikiriakos*,‡ †

Department of Chemistry and ‡Department of Chemical and Biological Engineering, University of British Columbia, 2036 Main Mall, Vancouver, BC V6T 1Z4, Canada S Supporting Information *

ABSTRACT: A series of poly(lactide) (PLA) samples, exhibiting various levels of syndiotactic enrichment, were formed via the ring-opening polymerization of meso-lactide using two families of dinuclear indium catalysts: (RR/RR)-[(NNO)InCl]2 (μ-Cl)(μ-OEt) (1) and (RR/RR)[(ONNO)In(μ-OEt)]2 (2). Isotactic and heterotactic PLAs were also synthesized using known methodologies, and the thermal and rheological behaviors of these PLAs with different microstructures were compared. Solution rheological studies showed that the values of intrinsic viscosities and hydrodynamic radii as functions of molecular weight (Mw) were highest for iso-PLAs, followed by hetero and then syndio-PLAs. The viscosities of the heterotactically enriched PLAs were in agreement with literature values reported for atactic PLAs. The molecular weight between entanglements (Me) was greatest for the syndiotactically enriched PLAs, giving rise to the lowest zero-shear viscosity. In addition, hetero- and isotactically enriched PLA had higher flow activation energies (Ea,flow) than syndiotactic variants, implying the inclusion of transient aggregate regions within these polymers due to enhanced L- and D-interactions. Although strain hardening was observed for all types of PLAs, it was more dominant for isotactic PLAs due to stronger L- and D-interactions possibly leading to a small degree of stereocomplex microcrystallites.

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INTRODUCTION Biodegradable aliphatic polyesters such as poly(lactide) (PLA) have been the subject of intense attention in the past few years due to the possibilities of a range of potential applications for these bio-derived materials.1 PLA has some mechanical properties (such as tensile strength) which are superior to poly(styrene) and poly(ethylene terephalate) and which make it a promising alternative to petroleum-based polymers.2 Despite this promise, PLA is currently not a viable replacement for these materials, in part due to low thermal stability, poor melt strength, and poor elasticity.3 A convenient and efficient method of synthesizing PLAs with high molecular weight and narrow molecular weight distributions is through the ring-opening polymerization (ROP) of the chiral molecule lactide (LA).4 Various microstructures of PLA have been reported for the polymerization of a mixture of Dand L-lactide (rac-LA) and meso-LA.5 These subtle changes in tacticity are often stated to have a significant impact on polymer properties parallel to what has been observed for poly(styrene)6 and poly(propylene).7 The focus of many research groups has been on designing discrete organometallic complexes which impart control of PLA microstructure.4f,8 Many examples of catalysts for the formation of highly isotactic PLA (i-PLA), which can be generated from polymerization of L-LA or from the selective polymerization of rac-LA, have been reported.9 Comprehensive studies of the viscoelastic behaviors and mechanical properties of isotactic, stereocomplex, and atactic PLA have also been reported.10 Isotactic © XXXX American Chemical Society

PLA has a highly ordered microstructure, with plateau moduli, G0N, ranging around 0.9 ± 0.2 MPa and entanglement molecular weights (Me) of approximately 4400 g mol−1.10a,b,e These properties impart chain stiffness and strength to the polymer but leave the material brittle, with an elongation at break of just 3%. In contrast, reports of syndiotactic PLA (s-PLA) are scarce in the literature, with only two groups reporting highly syndioselective organometallic catalysts for the polymerization of meso-LA.11 Although examples of highly heterotactic PLA (hPLA) have been reported,12 there are currently no investigations into the thermorheological behaviors of these materials. A recent report comparing rheological properties of syndiotactic and isotactic poly(styrene) (s- and i-PS)6b showed that Me is dependent on tacticity, with s-PS having the lowest Me (14 500 g mol−1) giving rise to the highest zero-shear viscosity and energy barrier to flow. These results were in line with earlier studies exploring the viscoelastic properties of poly(propylene) homo-7a,b,g and copolymers.7f These accounts concluded that syndiotactic polymers exhibit higher activation energies of flow (Ea,flow) compared to their isotactic and atactic counterparts, implying a higher sensitivity of rheological properties on temperature. The tacticity dependence on Received: November 26, 2015 Revised: January 8, 2016

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DOI: 10.1021/acs.macromol.5b02568 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules Chart 1. Chiral Indium-Based Catalysts [(NNO)InCl]2(μ-Cl)(μ-OEt) (1) and [(ONNO)In(μ-OEt)]2 (2)

Figure 1. Plots of observed PLA Mn (blue ■) and Đ (red ●) as functions of added meso-LA for (A) [(NNO)InCl]2(μ-Cl)(μ-OEt) (±)-1, (B) (RR/ RR)-1, (C) [(ONNO)In(μ-OEt)]2 (±)-2, and (D) (RR/RR)-2 (Mn = number-averaged molecular weight, Đ = dispersity). The line represents theoretical Mn values based on the LA:initiator ratio at 100% conversion. All reactions were carried out with 2.8 mM catalysts for 16 h (catalyst 1) or 4 h (catalyst 2) at room temperature in CH2Cl2. All polymer samples were isolated at >90% monomer conversion.

rheological properties was also found in solutions of poly(vinyl alcohol),13 indicating that increasing degrees of syndiotacticity enhances chain stiffness as well as the entanglement density of the polymer chain. We have recently reported two families of active indium catalysts for the ROP of racemic lactide. The first-generation catalysts are dinuclear asymmetrically bridged diaminophenolate systems [(NNO)InCl]2(μ-Cl)(μ-OEt) (1) (Chart 1), which exerts moderate isoselectivity (Pm ∼ 0.65) for the ROP of rac-LA.14 The second-generation catalysts [(ONNO)In(OEt)]2 (2) are more active and show greater control of the polymerization of rac-LA, forming iso-rich gradient PLA (Pm ∼ 0.75) through enantiomorphic site control.15 We aim to generate a comprehensive work on the rheological properties of PLA with various microstructures (tacticities) in the context of the similar work with polystyrene. Thus, we first investigate the ring-opening polymerization of meso-LA by chiral catalysts 1 and 2 and examine the resulting syndiotactically enriched polymers. Additionally, we utilize other stereoselective systems to produce PLAs with varying levels of stereoregular microstructures with the aim of correlating

tacticity with the thermorheological behavior of the resulting materials.

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RESULTS AND DISCUSSION Polymerization of meso-Lactide with Catalysts 1 and 2. The racemic and enantiopure catalysts [(NNO)InCl]2(μCl)(μ-OEt) (±)- and (RR/RR)-(1)14a−c and [(ONNO)In(μOEt)]2 (±)- and (RR/RR)-(2)14e,15c were prepared according to known methods. Polymerization reactions with catalysts 1 and 2 and meso-LA were carried out in CH2Cl2 at room temperature for 16 and 4 h, respectively (Tables S1 and S2). Conversions and selectivities were determined by 1H and 1 H{1H} NMR spectroscopy (CDCl3, room temperature) and molecular weight and dispersity measurements were carried out by gel permeation chromatography (GPC). The polymers formed from the reaction of (±)- or (RR/RR)1 with meso-LA show good agreement between the calculated and experimental molecular weights (Figure 1, Tables S1 and S2). Low dispersity values (Đ < 1.05) indicate that the polymerization of meso-LA is living, as was observed for racLA.14a,b Polymerization reactions with (±)- or (RR/RR)-2 B


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Figure 2. Plot of ln[LA] versus time for polymerization of meso-LA catalyzed with (a) left, (±)-1 (blue ●) and (RR/RR)-1 (green ◆) and (b) right, (±)-2 (blue ●) and (RR/RR)-2 (green ◆). Reactions were carried out in an NMR tube at 25 °C and followed to 90% conversion. 1,3,5Trimethoxybenzene (TMB) was used as internal standard. All reactions were carried out with 200 equiv of LA in CDCl3 at 25 °C and followed to 90% conversion by 1H NMR spectroscopy [1] = 0.0023 M, [LA] = 0.46 M. [2] = 0.0018 M, [LA] = 0.36 M. The value of kobs was determined from the slope of ln[LA] vs time, averaged from at least three experiments.

show less control over molecular weight at higher monomer loadings than those carried out with 1 due to depolymerization reactions, catalyzed by 2 at the longer reaction times necessary for higher molecular weights, as discussed previously.15c In Situ Studies and Polymer Tacticity. Polymerization of meso-LA using (±)- and (RR/RR)-1 and -2 was performed at room temperature in CDCl3 with a monomer:initiator ratio of 200 and was monitored to >90% conversion using 1H NMR spectroscopy. The plots of ln[LA] vs time for all catalysts are linear after a short initiation period for catalyst 1 and a longer period for 2 (Figure 2, Table S3). The long initiation period for 2 is due to dissociation of the dimer prior to polymerization.15c Catalysts (±)- and (RR/RR)-2 exhibit higher propagation rates for the ROP of meso-LA than (±)- and (RR/RR)-1, which is consistent with their relative rates of polymerization of rac-LA (Figure 2).14b,15c The kobs value for polymerization of meso-LA with (RR/RR)1 is over 5 times lower than that of the racemic analogue (±)-1 (Figure 2a, Table S3, entries 1−3). The 1H{1H} NMR analysis of the methine region of PLA from the polymerization of mesoLA with the two catalysts shows syndiotactic enrichment for both (±)-1 (Ps ∼ 0.83) and (RR/RR)-1 (Ps ∼ 0.86) (Figures S2 and S3). Enantiopure and racemic 2 display similar kobs values leading to similar syndioselectivities (Ps ∼ 0.78) toward mesoLA (Table S3). Synthesis of Varied Microstructure PLAs. In the current study we compare (a) syndio-rich PLAs generated from polymerization of meso-LA with catalyst (RR/RR)-1, (b) isorich gradient PLAs generated from polymerization of rac-LA with catalyst (RR/RR)-2, (c) essentially atactic PLA (with a slight heterotactic bias) generated from polymerization of racLA with tin octanoate, and (d) highly heterotactic PLA generated from polymerization of rac-LA with an InCl3/amine/ alcohol system (Table 1 and Table S4).12k,l The polymers are designated labels based on their molecular weight and tacticity; for example, 155-het-60 denotes heterotactic polymers with Mn of 155 kg mol−1 and Pr of 0.60 We also attempted to synthesize highly syndiotactic PLA with a system as closely related to ours as possible: the enantiopure aluminum salen binaphthylamine alkoxide complexes to generate highly syndiotactic PLA.11a,16 While these catalysts form highly syndiotactic PLA at low molecular weights (Mn ∼ 12 kg mol−1) in our hands, at the high molecular weights required for rheological measurements (Mn > 100 kg mol−1)

Table 1. PLA with Varied Microstructure Used in Comparative Rheological Studies entry

sample

Mn,expta (g mol−1)

tacticityb

Đa

1 2 3 4

155-het-60c 182-het-96 147-syn-86 119-iso-71

154600 182100 146700 118600

Pr = 0.60 Pr = 0.96 Ps = 0.86 Pm = 0.71

1.12 1.05 1.02 1.04

a

Molecular weights were determined by GPC-LALLS (gel permeation chromatography−low angle laser light scattering) to the polystyrene standard calibration curve via the Mark−Houwink equation in THF at 25 °C ([η] = KMa, while [η] = intrinsic viscosity, M = molecular weight, and K and a are Mark−Houwink parameters K = 1.832 × 10−4 dL g−1, and a = 0.69; dn/dc = 0.044 mL g−1). THF 4 mg mL−1 and flow rate = 0.5 mL min−1. bCalculated from 1H{1H} NMR spectra. c 155-het-60 denotes heterotactic polymers with a molecular weight of 155 kg mol−1 and a Pr value of 0.60.

only moderate syndio-enrichment was attained (Ps = 0.89); thus, these materials were not further analyzed (Figure S9). We attribute this decrease in syndiotacticity to the greater extent of polymer exchange at the long reaction times required to polymerize large monomer loadings for generation of high molecular weight material.16c During the characterization of these materials, some literature discrepancies in Ps calculations were noted (Supporting Information section D). Purely syndiotactic PLA is attained by ring-opening at one of the two enantiotopic acyl-oxygen sites. Selectivity is determined by the rate at which the catalyst ring-opens at one position over the other (Figure 3).16d This ratio can be quantified by the α-value and is determined experimentally by comparing the ratio of the integrations for the [sss] and [sis] tetrads expected from pure meso-LA (Table S5). However, as industrially produced meso-lactide is a byproduct of lactide production from epimerized L-lactic acid, even after careful separation it can be contaminated with significant quantities of D- and L-LA.4a These impurities can be incorporated into the growing polymer, leading to stereodefects beyond the [sis] tetrad (Table S5).8g These defect peaks also arise if extensive transesterification occurs during polymerization. The syndiotactic units within the polymer chain can be quantified by the Ps value and are determined by comparing the ratio of [sss] tetrad sequences to stereodefect tetrad sequences (Supporting Information section D). The inclusion of these C


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Figure 3. Formation of syndiotactic PLA and quantification of the α-value.

defects decreases the calculated Ps.16d Inconsistencies arise when peak intensities are used for calculations or if defect peaks are not taken into consideration during calculations. The authors recently remarked upon the need for consistency when reporting Pm values for rac-LA;15c we believe that a consistent method should be used when determining and comparing similar values for meso-LA. We have included the integration of all defect peaks in our calculations (Figures S2−S9). Thermal Study of Polymers. Thermogravimetric analysis (TGA) and differential scanning calorimetery (DSC) were used to determine the thermal properties for PLA samples with different degrees of stereoregularity (Table 2, Figures S12 and S13).

polymers show that the syndiotactic chains can pack well to form crystalline domains.6h,7a Thus, one can presume that extended syndiotactic domains should allow the formation of aggregates increasing the Tg of syndiotactic PLA compared to atactic PLA. Interestingly, syn-86 has a Tg comparable to atactic PLA (het-60), indicating that such aggregates are scarce within the polymer. Highly syndiotactic PLA reported by Coates in 1999 shows a glass transition temperature lower than that of heterotactic and isotactic PLAs.11a Isotactically enriched iso-71 exhibits the largest Tg of the polymers studied due to the strong interactions between PLLA and PDLA domains within the stereogradient polymer backbone (Table 2, entry 4). As a comparison, values were obtained for purely isotactic PLA18 (Table 2, entry 5) which displays a high glass transition temperature, supporting the rationalization above. Heterotactic PLA, het-96, exhibits the second highest Tg (37.9 °C), implying that a highly ordered heterotactic polymer has fewer interchain interactions than a moderately isotactic material (iso-71) (Table 2, entry 2). Interestingly, although heterotactic PLAs are prevalent in the literature,8e,g there is only one report of a melting temperature (Pr = 0.95, Tg = ∼46 °C, Tm = ∼120 °C).12p Solution Rheology. The average values of the intrinsic viscosities and the weight-average molecular weight were obtained from light scattering GPC in THF and are plotted in Figure 4. The slope of the linear regression line (exponent of the Mark−Houwink equation) for the polymers in this study is 0.73, which agrees with the values reported by Dorgan et al. for PLAs in THF.10b,e It is noted that for linear polymers with random coil conformation the exponent of the Mark−Houwink equation has a value from 0.5 in a poor solvent to 0.8 in a good

Table 2. Thermal Properties of Polymers with Varying Tacticities entry 1 2 3 4 5 6

het-60 het-96 syn-86 syn-96c iso-71 iso-100d

T5%a (°C)

Tgb (°C)

258(11) 290(5) 323(2)

31.9 37.9 30.8 34 46.2 55

319(6) 310

Tmb (°C)

151 175

a

Thermogravimetric analysis was performed on approximately 20 mg of material. Samples were heated to 500 °C at a rate of 20 °C/min to determine the degradation onset temperature (temperature at which there is 5% weight loss, T5%). Calculated standard errors given in parentheses. bThermal analysis of samples was performed by using a differential scanning calorimeter (DSC) with ca. 2 mg of sample. Samples heated to 170 °C at 10 °C/min and cooled to 25 °C at 5 °C/ min to determine Tg and Tm. Glass transition and melting temperatures calculated from first heating scans. cReference 11a. d Reference 18.

The TGA data of the varied PLAs in this study show similar degradation profiles (Figure S12). These results are in line with studies on poly(styrene) which showed that degradation mechanisms are dependent on repeat unit rather than microstructure. The authors noted a weak tacticity effect on the thermal stability of PS where atactic PS displayed low stability, while isotactic PS showed the highest thermal stability.6k However, contrary to the studies with PS, the syn86 polymers show the highest degradation onset temperature (T5%) and thus better stability compared to the other microstructures (Table 2). All of the polymers in this study are amorphous and only show glass transitions (Tg) in their DSC isotherms (Figure S13). Reports of stereoregular vinyl polymers such as poly(methyl methacryalate) (PMMA) show a tacticity dependence of the glass transition temperatures.6h,17 In the current study a significant effect of tacticity on glass transition temperatures is also observed (Table 2). We can rationalize these differences by invoking the presence of strong interchain interactions. Reports of other syndiotactic

Figure 4. A log−log plot of intrinsic viscosity of various types of PLAs as a function of the molecular weight. The intrinsic viscosities of isoPLAs are the highest compared to PLAs possessing other stereocomplex microstructure. D


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Macromolecules solvent.19 Figure 4 shows that the intrinsic viscosity of the isotactic (iso-71) is higher compared to heterotactic polymers (het-96), which in turn are higher than those of the syndiotactic counterpart with the latter possessing the lowest intrinsic viscosity of all different types of PLAs. This is due to different chain conformations, although the slopes of 0.75−0.76 for all the different families imply good solvent conditions and random coil conformation.19 Figure 5 shows the hydrodynamic radii dependence on the weight-averaged molecular weight for the range of polymers

molecular weight and a broadening of dispersity values were observed after the test.20 Dynamic linear viscoelastic measurements focused on the comprehensive characterization of syndiotactic and heterotactic PLAs and comparison with isotactic and atactic PLAs. The rheological behaviors of isotactic and atactic PLAs have been reported previously.10a−c,e,21 These measurements were conducted within the linear viscoelastic regime at temperatures in the range of 70−190 °C and at angular frequencies ranging from 0.01 to 100 rad/s with a constant strain of 2%. A gap of 0.5 mm was used to minimize edge effects. The isothermal frequency sweep measurements for each microstructure show an expected decrease in complex viscosity with temperature, and the materials display shear-thinning behavior. All PLA polymers studied were found to be thermorheologically simple, allowing for the application of the time−temperature superposition to generate master curves. Figure 6 shows representative master curves for the linear

Figure 5. Characteristic hydrodynamic radius (Rh) as a function of molecular weight and tacticity. The radii of the iso- PLA are the highest while those of syndio-PLA are the lowest, in agreement with the intrinsic viscosity data plotted in Figure 4.

studied. The scaling relation for heterotactic (het-96) and atactic (het-60) PLAs is Rh = 0.017Mw0.55, in agreement with trends for amorphous PLAs reported by Othman and coworkers.10e The radii of iso-71 polymers are higher than those expected by this relation while the radii of syn-86 polymers are lower. This result correlates well with the calculated mass dependence of the intrinsic viscosity plotted in Figure 4. Linear Viscoelasticity. An important aspect of polymer rheology is the study of how viscoelastic behavior changes with respect to temperature, molecular weight, and polymer microstructure. In particular, it is significant to determine structure−property relationships, i.e., zero-shear viscosity versus molecular weight, which is useful in assessing the flow and processing properties of materials. Currently there are no reports of these properties for heterotactic or syndiotactic PLA. The thermal stability of the various microstructured PLAs under shear stress was probed in isothermal time sweep experiments at 180 °C over 60 min. A parallel plate rotational rheometer was used at a constant frequency of 0.5 rad/s and strain amplitude of 2%. The complex modulus, |G*|, a measure of resistance to deformation sensitive to structural changes, is plotted against time in Figures S14−S17. Heterotactic (het-96) and atactic (het-60) PLAs show a steady decrease in complex modulus and associated melt strength over the experimental time frame. Subsequent GPC analysis conducted on the collected material showed a 20% decrease in molecular weight consistent with the fundamental relationship of η0 ∝ Mw3.4 for linear macrostructure. Isotactically enriched PLA (iso-71) and syndiotactic (syn-86) samples showed stable melt strength over the experiment time frame; however, a small decrease in

Figure 6. Master curve of the linear viscoelastic moduli, G′ and G″, and complex viscosity, |η*| (Tref = 150 °C), for (a) 182-het-96 and (b) 147-syn-86 polymers.

viscoelastic moduli, G′ and G″, as well as the complex viscosity at the reference temperature of 150 °C for het-96 (Figure 6a) and syn-86 (Figure 6b) polymers. Trends for the storage modulus are similar for both heterotactic and syndiotactically enriched PLAs where the storage moduli exhibit a clear plateau value (plateau modulus) at high frequencies. Similar trends are observed for the loss modulus which increases with increase of frequency, reaching a maximum and a minimum value at higher frequencies typical of linear monodisperse polymers. The E


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Macromolecules characteristic slopes of 1 and 2 for G′ and G″, respectively, show that the terminal zone at small deformation frequencies has been reached (Figure 6). The linear viscoelastic moduli of syn-86 and het-96 polymers with varying molecular weights are shown in Figures 7 and 8,

Figure 8. Master curves of the linear viscoelastic moduli for syn-86 polymers (Tref = 150 °C): (a) loss modulus vs angular frequency, (b) storage modulus vs angular frequency, and (c) complex viscosity vs angular frequency.

respectively. Trends seen in these plots are in line with those reported for isotactic PLAs.10e By increasing the molecular weight, the terminal zone shifts to lower frequencies (Figures 7a,b and 8a,b), the degree of shear thinning increases (Figures 7c and 8c), the plateau modulus, G0N is independent of molecular weight, and the response of the material becomes

Figure 7. Master curves of the linear viscoelastic moduli for het-96 polymers (Tref = 150 °C): (a) loss modulus vs angular frequency, (b) storage modulus vs angular frequency, and (c) complex viscosity vs angular frequency. F


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Macromolecules nearly independent of Mw at higher frequencies. All these observations are typical of linear monodisperse polymers. Reports comparing the viscoelastic behaviors of syndiotactic and isotactic poly(propylene) homo- and copolymers7a,b,f showed that increasing the degree of syndiotacticity increases chain stiffness as well entanglement density, imparting strength to syndiotactic polymers by increasing the plateau modulus, G0N. The plateau modulus for the PLAs of various microstructure studied were obtained from the minima of van Gurp−Palmen plots22 (Figures S18−S20). From the G0N, the molecular weight between entanglements, Me, can be calculated using Me =

ρRT GN0

(1)

where ρ is the melt density (g/cm ) at the chosen reference temperature, obtained using the relationship10e 3

Figure 9. Zero-shear viscosity, η0, versus weight-averaged molecular weight for PLAs with different microstructures.

4

ρ(t ) = 1.2836e(−7.7 × 10 T )

(2)

Moderately syndiotactic PLA, displaying the lowest plateau modulus and largest Me, exhibits lower chain stiffness compared to heterotactic and atactic polymers, possibly due to a decrease in the number of PLLA and PLDA interactions (Table 3). Heterotactic and atactic PLA with similar plateau moduli and Me thus are expected to have comparable chain stiffness.

the reference temperature (Figure S21). The obtained aT values were plotted against the reciprocal of temperature, and this was used to calculate the flow activation energy for all the microstructures studied (Table 3). The activation barrier of flow (Ea) is highest for the isotactically enriched polymer, iso-71 (Table 3, entry 4). This is evidence that strong chain interactions between regions of Land D-domains form aggregates, which need more energy to break apart and allow the material to flow. The small flow activation energy displayed by the syndiotactically enriched material, syn-86 (Table 3, entry 3), suggests that lasting aggregates are not formed between the syndiotactic units within the polymer chains. The high energy barrier observed for atactic PLA, het-60 (Table 3, entry 1), can be explained by the statistical distribution of the different tacticities within its structure. Depending on the synthetic conditions for this material, it may contain enough regions of isotactic enrichment to form transient aggregates (enhanced PLLA and PDLA interactions) which cause the higher activation energy compared to the heterotactic PLA, het-96 (Table 3, entry 2). Uniaxial Extensional Rheology. The extensional viscosity of polymeric materials is greatly influenced by the presence of chain interactions at different chain scales and thus can be used as a measure to probe the effect of microstructure on the melt properties of PLA. Strain hardening for linear polymers is unusual; however, Palade10a showed that high molecular weight isotactic PLA exhibits strain hardening at relatively low strain rates. In 2004, Yamane reported that the formation of stereocomplex microcrystallites in PLLA/PDLA blends resulted in temporary cross-links between polymer chains allowing for strain hardening.24 In addition, the time scales of the deformation compared with the terminal relaxation times of the polymers might also favor the occurrence of strain hardening. Extensional tests at low temperatures are useful in illustrating the effect of microstructure on the melt strength of PLA as any strain hardening will be amplified. Uniaxial extensional tests were carried out at 70, 90, and 110 °C at Hencky strain rates of 0.01−10 s−1. Strain hardening was most significantly observed at strain rates of 10 s−1. Figure 10 shows the log−log plots of extensional viscosity versus time at the three temperatures. Strain hardening behavior can be observed when the molecular relaxation time exceeds the characteristic time of

Table 3. Comparison of Rheological Properties of PLA, PS, and PP in This Work and Obtained from the Literature6b,7g,23 entry

polymer

G0N a (105 Pa)

Meb (g/mol)

1 2 3 4 5d 6 7d 8f 9f 10f

het-60 het-96 syn-86 iso-71 a-PP s-PP i-PP a-PS s-PS i-PS

0.58 0.52 0.34 0.97 0.42 0.87e 0.43 0.24 0.30 0.16

6900 7700 11800 4100 7050 3400e 6900 17900 14500 27200

Eac (kJ mol−1) 152.7 141.0 130.2 167.9

(0.5) (0.9) (0.6) (0.6)

50.6d 38.7 97(9) 53 (5) 99

a

Calculated from minimum of van Gurp−Palmen plot (Figures S18− S21). bObtained from eqs 1 and 2. cCalculated from the slope of the horizontal shift factor, aT vs 1/T plot (Figure S22). Calculated standard errors given in parentheses. dReference 7g. eReference 23. f Reference 6b. To obtain the G0N and Me the references used for the PS and PP samples are 280 and 190 °C, respectively.

Figure 9 shows a plot of zero-shear viscosity, η0, against weight-averaged molecular weight, Mw, for the various microstructures in the present study. The lines drawn in Figure 9 for the iso- and syndio-PLAs are best fits with a fixed exponent of 3.4 that is well established for the atactic PLAs. The zero-shear viscosity data for heterotactic and atactic PLAs agree well with the relation reported by Othman in 2011 for amorphous PLA (continuous line in Figure 9).10e Isotactically enriched polymers have a higher zero-shear viscosity compared to the syndiotactic counterparts (in agreement with the intrinsic viscosity comparison presented in Figure 4). Horizontal shift factors, aT, determined from the time− temperature superposition were found to agree with those reported in the literature10f and followed an Arrhenius trend, aT = exp{Ea/R(1/T − 1/Tref}, where Ea is the activation energy for flow, R is the universal constant of the ideal gas law, and Tref is G


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Table 4. Average Relaxation Times for Various Polymers Studied Calculated from Linear Viscoelastic Experiments (LVEs) λave (s) entry 1 2 3 4

het-60 het-96 syn-86 iso-71

70 °C

90 °C

110 °C

9.97 7.29 9.46

4.72 4.65 7.26 7.22

2.79 1.48 2.70 2.97

i-PLAs. It seems that these interactions are stronger in the case of isotactic PLAs (enhancing the zero-shear viscosity as discussed above), and thus strain hardening persists even at temperatures of 110 °C at all extensional rates (Figure 11). However, this is not the case for the other types of PLAs where at 110 °C there is no sign of strain hardening (Figures S22− S25).

Figure 11. Tensile stress growth coefficient (measure of elongational viscosity) as a function of time at Hencky strain rates from 0.01 to 10 s−1 for iso-enriched gradient PLA (Table 3, entry 4). It is noted that strain hardening is present at all Hencky strain rates (not the case for all other PLA types), indicating the stronger interactions of L- and Dregions.

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CONCLUSIONS Polymerization Behavior of Reported Indium Complexes Differs between Lactide Monomers. Two series of dinuclear indium coordination complexes were employed to investigate the catalytic ring-opening polymerization of mesolactide. Catalyst 1 showed good molecular weight control for the ring-opening polymerization of meso-LA. Polymerization of meso-LA with (RR/RR)-1 generated PLA with the highest syndiotacticity of all complexes in this study. Catalyst 2 showed less control over molecular weight and displayed lower levels of syndiotactic control. This is in contrast with previous studies of rac-LA where catalyst 2 exerted the greatest control on polymer tacticiy. Stereocontrol mechanisms operating in these systems were probed by studying the kinetics of the polymerization reaction; however, reaction rates alone gave little further evidence of how syndiocontrol arises during the polymerization process. Tacticity Influences Chain Interactions and Consequently Polymer Properties. In an attempt to generate a

Figure 10. Tensile stress growth coefficient (measure of elongational viscosity) as a function of time at Hencky strain rate of 10 s−1 for PLAs of different microstructure (tacticity) at (a) 70, (b) 90, and (c) 110 °C.

deformation.10a,d For a Hencky strain rate of 10 s−1, the characteristic deformation time is 0.1 s. The average relaxation times for each of the polymers studied were calculated from the linear viscoelastic experiments (LVE) experiments and are shown in Table 4. The average relaxation times are high at 70 and 90 °C (5−10 s), which explains the presence of strain hardening at 10 s−1 (with characteristic time of deformation of 0.1 s). However, another contributing factor to strain hardening is the interactions between L- and D-domains with formations of transient stereocomplex aggregates, particularly in the case of H


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Macromolecules

and was dried over CaH2, transferred under vacuum, and degassed before use. CDCl3 was purchased from Cambridge Isotope Laboratories Inc. and dried over CaH2, transferred under vacuum, and degassed through three freeze−pump−thaw cycles before use. 2Propanol was purchased from Fischer Chemicals and was dried over 3 Å molecular sieves before use. Indium(III) trichloride was purchased from Sigma-Aldrich and was used as received. Triethylamine and benzyl alcohol were purchased from Sigma-Aldrich and were dried over CaH2, transferred under vacuum, and degassed before use. rac-LA and meso-LA were a gift from PURAC America Inc. rac-LA was recrystallized twice from hot dried toluene prior to use. Racemic and enantiopure complexes 1 and 2 were prepared according previously reported methods.14b,15a Purification of meso-Lactide. meso-LA (50 g) was fully dissolved in dichloromethane; magnesium sulfate (anhydrous) was added to the solution, which was then stirred for 10 min. The excess solid was removed from white suspension via vacuum filtration, and the solvent was removed from the filtrate under reduced pressure to yield a white crystalline powder. This powder was recrystallized from dry 2propanol five to seven times to gave meso-LA with 97−99% purity (by GCMS, see Figures S10 and S11). Polymerization of meso-LA with Catalyst 1. A glass vial was equipped with a magnetic stir bar and charged with meso-lactide (290 mg, 2.7 mmol) in 2 mL of CH2Cl2. To this, a solution of the catalyst (2.0 mg, 0.002 mmol) in 2 mL of CH2Cl2 was added. The solution was stirred at room temperature for 16 h. Solvent was removed under reduced pressure, and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in minimal CH2Cl2 and precipitated using excess of ice-cold methanol at least three times to remove residual catalyst. The polymer sample was then dried under vacuum for 48 h. No stabilizers were added. Polymerization of meso-LA with Catalyst 2. A glass vial was equipped with a magnetic stir bar and charged with meso-lactide (620 mg, 4.3 mmol) in 2 mL of CH2Cl2. To this, a solution of the catalyst (2.0 mg, 0.001 mmol) in 2 mL of CH2Cl2 was added. The solution was stirred at room temperature for 4 h to minimize transesterification and backbiting reactions. Solvent was removed under reduced pressure, and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in minimal CH2Cl2 and precipitated using excess of ice-cold methanol at least three times to remove residual catalyst. The polymer sample was then dried under vacuum for 48 h. No stabilizers were added. Synthesis of Heterotactic PLA. A vacuum adapted flask was equipped with a magnetic stir bar and was charged with rac-lactide (2.6 g, 0.02 mol) in 5 mL of toluene. InCl3 (2 mg, 0.009 mmol), benzyl alcohol (0.9 μL, 0.009 mmol), and triethylamine (2.5 μL, 0.02 mmol) were then added, and the volume was made up to 15 mL. The reaction was stirred at 80 °C for 3 days under N2. Solvent was removed under reduced pressure, and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in minimal CH2Cl2 and precipitated using excess of ice-cold methanol at least three times to remove residual catalyst. The polymer sample was then dried under vacuum for 48 h. No stabilizers were added. Synthesis of Atactic PLA. A vacuum adapted flask was equipped with a magnetic stir bar and charged with rac-lactide (2.0 g, 0.01 mol) in 5 mL of toluene. Sn(oct)2 (4.5 μL, 0.04 mmol) was then added, and the volume was made up to 15 mL. The reaction was then stirred at 80 °C for 3 days under N2. Solvent was removed under reduced pressure, and the conversion was determined using 1H NMR spectroscopy. The polymer was then dissolved in minimal CH2Cl2 and precipitated using excess of ice-cold methanol at least three times to remove residual catalyst. The polymer sample was then dried under vacuum for 48 h. No stabilizers were added. In Situ Observation ROP of meso-LA. All samples for NMR scale polymerization were prepared in Teflon-sealed NMR tubes under an N2 atmosphere. The NMR tube was charged with a stock solution of catalyst in CDCl3 (0.25 mL, 0.0011 mmol) and frozen in liquid N2. CDCl3 (0.25 mL) was added and again frozen to create a buffer between the catalyst and the lactide. Finally, a stock solution of mesoLA (0.50 mL, 0.45 mmol) and the internal standard 1,3,5-

comprehensive correlation between polymer microstructures determined by in situ NMR methods (Pm/Ps) to bulk polymer properties, we synthesized and carried out rheological studies on families of atactic, syndiotactic, isotactic, and heterotactic polymers with varying molecular weights. To our knowledge this is the first report of the thermal and rheological properties of syndio- or heterotactically enriched PLA. Comparisons of rheological properties showed that to achieve large aggregate domains, chain packing from isotactic and heterotactic domains is superior to those of syndiotactic domains. Our results showed that the measured Me is lowest for isotactic PLA (i-PLA) and highest for syndiotactically enriched PLA (s-PLA). Because of this, s-PLA has the lowest entanglement density for a given Mw (highest Me), which in turn gives rise to the lowest zero-shear viscosity. In addition, heterotactic PLA (h-PLA) and i-PLA have higher Ea,flow than sPLA. This is consistent with h- and i-PLA having larger aggregate regions than s-PLA. These results are in agreement with the solution properties of various PLAs where s-PLA exhibits the lowest intrinsic viscosity and melt viscosity. The linear viscoelastic properties of PLA showed greater dependence on the level of isotactic rather than syndiotactic linkages present in the polymer chain, with isotactically enriched polymers having higher flow activation energy than syndiotactically enriched polymers. Heterotactic PLAs exhibited nascent aggregate domains which were observed by an increase in flow activation energy compared to syndiotactic PLA. Furthermore, it was determined that strain hardening was possible for h- and i-PLA, showing that transient aggregates due to enhanced L- and D-interactions, more common in isotactically and heterotactically inclined polymers compared to syndiotactic polymers, can increase the relaxation times. As a result, strain hardening was present for i-PLA even at temperatures as high as 110 °C.

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EXPERIMENTAL SECTION

General Methods. All the air- and moisture-sensitive manipulations were carried out in an MBraun glovebox or using standard Schlenk line techniques. A Bruker Avance 300 or 400 MHz spectrometer was used to record 1H spectra. A Bruker Avance 600 MHz spectrometer was used to acquire homonuclear decoupled 1 H{1H} spectra of PLA. 1H NMR chemical shifts are given in ppm versus residual protons in deuterated solvents as follows: δ 7.27 CDCl3. Molecular weights, hydrodynamic radii, and intrinsic viscosities were determined by GPC-LLS using a Agilant liquid chromatograph equipped with a Agilant 1200 series pump and autosampler, three Phenogel 5 μm narrow bore columns (4.6 × 300 mm with 500, 103, and 104 Å pore size), a Wyatt Optilab differential refractometer, Wyatt tristar miniDAWN (laser light scattering detector), and a Wyatt ViscoStar viscometer. The column temperature was set at 40 °C. A flow rate of 0.5 mL/min was used, and samples were dissolved in THF (ca. 4 mg/mL). The measurements were carried out at laser wavelength of 690 nm at 25 °C. The data were processed using the Astra software provided by Wyatt Technology Corp. A differential scanning calorimeter (DSC) Q1000 (TA Instruments) was employed to measure the glass transition (Tg) and melting (Tm) temperatures. Thermogravimetric analysis (TGA) traces were collected on a PerkinElmer Pyris 6 TGA with a nitrogen flow rate of 20 mL/min. Shear measurements were performed using a MCR 501 rheometer equipped with 8 mm parallel plates. Uniaxial extensional measurements were performed using the SER-2 extensional fixture attached to an Anton Paar MCR 502 rheometer. Materials. THF was taken from an IT Inc. solvent purification system with activated alumina columns and degassed before use. HPLC grade dichloromethane was purchased from Fischer Chemicals I


Article

Macromolecules

Tang, C. B. Macromolecules 2013, 46 (5), 1689−1712. (d) Chen, G. Q.; Patel, M. K. Chem. Rev. 2012, 112 (4), 2082−2099. (e) Vroman, I.; Tighzert, L. Materials 2009, 2 (2), 307−344. (f) Williams, C. K.; Hillmyer, M. A. Polym. Rev. 2008, 48 (1), 1−10. (g) Gupta, A. P.; Kumar, V. Eur. Polym. J. 2007, 43 (10), 4053−4074. (h) Mecking, S. Angew. Chem., Int. Ed. 2004, 43 (9), 1078−1085. (i) Okada, M. Prog. Polym. Sci. 2002, 27 (1), 87−133. (2) (a) Jamshidian, M.; Tehrany, E. A.; Imran, M.; Jacquot, M.; Desobry, S. Compr. Rev. Food Sci. Food Saf. 2010, 9 (5), 552−571. (b) Peelman, N.; Ragaert, P.; Ragaert, K.; De Meulenaer, B.; Devlieghere, F.; Cardon, L. J. Appl. Polym. Sci. 2015, 132 (48), 42305. (3) (a) Drumright, R. E.; Gruber, P. R.; Henton, D. E. Adv. Mater. 2000, 12 (23), 1841−1846. (b) Kricheldorf, H. R. Chemosphere 2001, 43 (1), 49−54. (4) (a) Dechy-Cabaret, O.; Martin-Vaca, B.; Bourissou, D. Chem. Rev. 2004, 104 (12), 6147−6176. (b) O’Keefe, B. J.; Hillmyer, M. A.; Tolman, W. B. J. Chem. Soc., Dalton Trans. 2001, No. 15, 2215−2224. (c) Chisholm, M. H.; Zhou, Z. P. J. Mater. Chem. 2004, 14 (21), 3081−3092. (d) Bourissou, D.; Martin-Vaca, B.; Dumitrescu, A.; Graullier, M.; Lacombe, F. Macromolecules 2005, 38 (24), 9993−9998. (e) Wu, J. C.; Huang, B. H.; Hsueh, M. L.; Lai, S. L.; Lin, C. C. Polymer 2005, 46 (23), 9784−9792. (f) Wu, J. C.; Yu, T. L.; Chen, C. T.; Lin, C. C. Coord. Chem. Rev. 2006, 250 (5−6), 602−626. (g) Kamber, N. E.; Jeong, W.; Waymouth, R. M.; Pratt, R. C.; Lohmeijer, B. G. G.; Hedrick, J. L. Chem. Rev. 2007, 107 (12), 5813− 5840. (h) Kricheldorf, H. R. Chem. Rev. 2009, 109 (11), 5579−5594. (i) Wheaton, C. A.; Hayes, P. G.; Ireland, B. J. Dalton Trans. 2009, 25, 4832−4846. (j) Ajellal, N.; Carpentier, J.-F.; Guillaume, C.; Guillaume, S. M.; Helou, M.; Poirier, V.; Sarazin, Y.; Trifonov, A. Dalton Trans. 2010, 39 (36), 8363−8376. (k) Vieira, I. D.; Herres-Pawlis, S. Eur. J. Inorg. Chem. 2012, 2012 (5), 765−774. (l) Brown, H. A.; Waymouth, R. M. Acc. Chem. Res. 2013, 46 (11), 2585−2596. (m) Dagorne, S.; Normand, M.; Kirillov, E.; Carpentier, J. F. Coord. Chem. Rev. 2013, 257 (11−12), 1869−1886. (5) Current convention uses Pr and Pm values (probability of racemic and meso linkages) when discussing rac-LA and Pi and Ps values (probability of iso and syndio linkages) when discussing meso-LA. (6) (a) Guenet, J. M.; Picot, C. Macromolecules 1983, 16, 205−10. (b) Huang, C.-L.; Chen, Y.-C.; Hsiao, T.-J.; Tsai, J.-C.; Wang, C. Macromolecules 2011, 44 (15), 6155−6161. (c) Nakaoki, T.; Kobayashi, M. J. Mol. Struct. 2003, 655, 343−349. (d) Wang, C. C.; Lin, C.-C.; Tseng, L.-C. Polymer 2006, 47, 390−402. (e) Woo, E. M.; Sun, Y. S.; Yang, C. P. Prog. Polym. Sci. 2001, 26, 945−983. (f) Chantini, Y.; Shimane, Y.; Ijitsu, T.; Yukinari, T. Polymer 1993, 34 (8), 1625−1629. (g) Chantini, Y.; Shimane, Y.; Inoue, Y.; Inagaki, T.; Ishioka, T.; Ijitsu, T.; Yukinari, T. Polymer 1992, 33 (3), 488−492. (h) Tomotsu, N.; Ishihara, N.; Newman, T. N.; Malanga, M. T. J. Mol. Catal. A: Chem. 1998, 128, 167−190. (i) Cimmino, S.; Di Pace, E.; Martuscelli, E.; Silvestre, C. Polymer 1991, 32, 1080−3. (j) Xu, G.; Chung, T. C. Macromolecules 1999, 32, 8689−8692. (k) Chen, K.; Harris, K.; Vyazovkin, S. Macromol. Chem. Phys. 2007, 208, 2525− 2532. (7) (a) Eckstein, A.; Friedrich, C.; Lobbrecht, A.; Mulhaupt, R. Acta Polym. 1997, 48, 41−46. (b) Eckstein, A.; Suhm, J.; Friedrich, C.; Maier, R. D.; Sassmannshausen, J.; Bochmann, M.; Mülhaupt, R. Macromolecules 1998, 31, 1335. (c) Jones, T. D.; Chaffin, K. A.; Bates, F. S.; Annis, B. K.; Hagaman, E. W.; Kim, M.-H.; Wignall, G. D.; Fan, W.; Waymouth, R. Macromolecules 2002, 35, 5061−5068. (d) Arrighi, V.; Batt-Coutrot, D.; Zhang, C.; Telling, M. T. F.; Triolo, A. J. Chem. Phys. 2003, 119, 1271−1278. (e) Antoniadis, S. J.; Samara, C. T.; Theodorou, D. N. Macromolecules 1999, 32, 8635−8644. (f) Fetters, L. J.; Lee, J. H.; Mathers, R. T.; Hustad, P. D.; Coates, G. W.; Archer, L. A.; Rucker, S. P.; Lohse, D. J. Macromolecules 2005, 38 (24), 10061− 10066. (g) Liu, C.; Yu, J.; He, J.; Liu, W.; Sun, C.; Jing, Z. Macromolecules 2004, 37, 9279−9282. (h) Li, Y.; Yao, Z.; Chen, Z.-H.; Qiu, S.-L.; Zeng, C.; Cao, K. Ind. Eng. Chem. Res. 2013, 52 (23), 7758−7767. (i) Madkour, T. M.; Soldera, A. Eur. Polym. J. 2001, 37, 1105−1113.

trimethoxybenzene (5 mg, 0.03 mmol per 0.50 mL) in CDCl3 was added and frozen. The sealed and evacuated NMR tube was immediately taken to the NMR spectrometer (400 MHz Avance Bruker spectrometer). The sealed tubes were thawed, and the polymerization was monitored at 25 °C. DSC Measurement of Polymers. Approximately 2−3 mg of polymer was weighed and sealed in an aluminum pan. Experiments were carried out under a nitrogen atmosphere. The samples were heated at a rate of 10 °C/min from 25 to 170 °C and held isothermally for 5 min to destroy any residual nuclei before cooling at 5 °C/min. The transition and melting temperatures were obtained from the first heating sequence, performed at 10 °C/min. TGA Measurement of Polymers. Approximately 20 mg of polymer was weighed into a ceramic crucible. Experiments were carried out under a nitrogen atmosphere at a flow rate of 20 mL/min. The samples were heated at a rate of 20 °C/min from 25 to 500 °C Linear Viscosity Measurements. All polymer samples were compression molded at 150 °C into discs with diameters of 16−50 mm and thickness 0.4−0.6 mm. The dynamic linear viscoelastic measurements were carried out within the linear viscoelastic regime at temperatures in the range from 70 to 190 °C. The dynamic measurements were conducted in the range of 0.01−100 rad/s at a strain of 2%. A gap of 0.5 mm was used to minimize edge effects and ensure a reasonable aspect ratio of plate radius and gap. Dynamic time sweep measurements were carried out at an angular frequency of 0.5 rad/s at 180 °C to examine the thermal stability of the samples. The rheological measurements were performed under a nitrogen atmosphere to minimize degradation of the polymer samples during testing. Uniaxial Extensional Measurements. Samples with diameters of 16−50 mm and thickness 0.4−0.6 mm were prepared by the same procedure used for shear samples. Individual polymer specimens were then cut to a width of 1.5−3.5 mm. Measurements were conducted at 70, 90, and 110 °C at Hencky shear rates of 0.01, 0.1, 1.0, and 10 s−1.

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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.5b02568. Molecular weight polymerization data of meso-lactide with catalysts 1 and 2, 1H {1H} NMR spectra of polymers and determination of tacticity, kinetic investigations, and comparison to literature catalysts, GCMS traces, TGA traces, DSC traces, isothermal time sweep experiments, hydrodynamic radius data, Van Gurp− Palmen plots, uniaxial extensional experiment plots, Arrhenius plot, and determination of activation of flow energy (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail [email protected] (P.M.). *E-mail [email protected] (S.G.H.). Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS P.M. and S.G.H. acknowledge support from NSERC in the form of a joint Strategic Grant. REFERENCES

(1) (a) Holmberg, A. L.; Reno, K. H.; Wool, R. P.; Epps, T. H. Soft Matter 2014, 10 (38), 7405−7424. (b) Kataoka, K.; Harada, A.; Nagasaki, Y. Adv. Drug Delivery Rev. 2012, 64, 37−48. (c) Yao, K. J.; J


Article

Macromolecules

11649−11657. (m) Buffet, J.-C.; Okuda, J. Chem. Commun. 2011, 47 (16), 4796−4798. (n) Blake, M. P.; Schwarz, A. D.; Mountford, P. Organometallics 2011, 30 (5), 1202−1214. (o) Bouyahyi, M.; Ajellal, N.; Kirillov, E.; Thomas, C. M.; Carpentier, J. F. Chem. - Eur. J. 2011, 17 (6), 1872−1883. (p) Romain, C.; Heinrich, B.; Laponnaz, S. B.; Dagorne, S. Chem. Commun. 2012, 48 (16), 2213−5. (q) Sauer, A.; Buffet, J. C.; Spaniol, T. P.; Nagae, H.; Mashima, K.; Okuda, J. Inorg. Chem. 2012, 51 (10), 5764−70. (r) Nayab, S.; Jeong, J. H. Polyhedron 2013, 59, 138−143. (s) Klitzke, J. S.; Roisnel, T.; Kirillov, E.; Casagrande, O. L.; Carpentier, J.-F. Organometallics 2014, 33 (1), 309−321. (t) Kong, W. L.; Chai, Z. Y.; Wang, Z. X. Dalton Trans. 2014, 43 (38), 14470−80. (u) Yang, S.; Nie, K.; Zhang, Y.; Xue, M.; Yao, Y.; Shen, Q. Inorg. Chem. 2014, 53, 105. (v) Pilone, A.; De Maio, N.; Press, K.; Venditto, V.; Pappalardo, D.; Mazzeo, M.; Pellecchia, C.; Kol, M.; Lamberti, M. Dalton Trans. 2015, 44 (5), 2157−65. (w) Fuoco, T.; Meduri, A.; Lamberti, M.; Venditto, V.; Pellecchia, C.; Pappalardo, D. Polym. Chem. 2015, 6 (10), 1727−1740. (13) Lyoo, W. S.; Kim, B. C.; Ha, W. S. Polym. Eng. Sci. 1997, 37, 1259. (14) (a) Douglas, A. F.; Patrick, B. O.; Mehrkhodavandi, P. Angew. Chem., Int. Ed. 2008, 47 (12), 2290−2293. (b) Yu, I.; Acosta-Ramirez, A.; Mehrkhodavandi, P. J. Am. Chem. Soc. 2012, 134 (30), 12758− 12773. (c) Acosta-Ramirez, A.; Douglas, A. F.; Yu, I.; Patrick, B. O.; Diaconescu, P. L.; Mehrkhodavandi, P. Inorg. Chem. 2010, 49 (12), 5444−5452. (d) Osten, K. M.; Yu, I.; Duffy, I. R.; Lagaditis, P. O.; Yu, J. C. C.; Wallis, C. J.; Mehrkhodavandi, P. Dalton Trans. 2012, 41 (26), 8123−8134. (e) Aluthge, D. C.; Xu, C. L.; Othman, N.; Noroozi, N.; Hatzikiriakos, S. G.; Mehrkhodavandi, P. Macromolecules 2013, 46 (10), 3965−3974. (f) Mehrkhodavandi, P.; Aluthge, D. C.; Clark, T. J.; Mariampillai, B.; Yan, Y. Salen indium catalysts and methods of manufacture and use thereof. PCT/CA2013/050191, 2013. (g) Osten, K. M.; Aluthge, D. C.; Patrick, B. O.; Mehrkhodavandi, P. Inorg. Chem. 2014, 53 (18), 9897−9906. (h) Osten, K. M.; Aluthge, D. C.; Mehrkhodavandi, P. Dalton Trans. 2015, 44, 6126−6139. (i) Yu, I.; Ebrahimi, T.; Hatzikiriakos, S. G.; Mehrkhodavandi, P. Dalton Trans. 2015, 44, 14248−14254. (15) (a) Aluthge, D. C.; Patrick, B. O.; Mehrkhodavandi, P. Chem. Commun. 2013, 49 (39), 4295−4297. (b) Aluthge, D. C.; Yan, E. X.; Ahn, J. M.; Mehrkhodavandi, P. Inorg. Chem. 2014, 53 (13), 6828− 6836. (c) Aluthge, D. C.; Ahn, J. M.; Mehrkhodavandi, P. Chem. Sci. 2015, 6, 5284−5292. (16) (a) Wisniewski, M.; LeBorgne, A.; Spassky, N. Macromol. Chem. Phys. 1997, 198 (4), 1227−1238. (b) Bernardo, K. D. S.; Robert, A.; Dahan, F.; Meunier, B. New J. Chem. 1995, 19 (2), 129−131. (c) Ovitt, T. M.; Coates, G. W. J. Polym. Sci., Part A: Polym. Chem. 2000, 38, 4686−4692. (d) Ovitt, T. M.; Coates, G. W. J. Am. Chem. Soc. 2002, 124 (7), 1316−1326. (17) (a) Karasz, F. E.; MacKnight, W. J. Macromolecules 1968, 1 (6), 537−540. (b) Grohens, Y.; Brogly, M.; Labbe, C.; David, M. O.; Schultz, J. Langmuir 1998, 14 (11), 2929−2932. (c) Hatada, K. J. Polym. Sci., Part A: Polym. Chem. 1999, 37 (3), 245−260. (18) Garlotta, D. J. Polym. Environ. 2001, 9 (2), 63−84. (19) Burchard, W. Solution Properties of Branched Macromolecules. In Branched Polymers II; Roovers, J., Ed.; Springer: Berlin, 1999; Vol. 143, pp 113−194. (20) Linear viscoelastic measurements last about 20 min, during which time the decrease in molecular weight is estimated to be 6%. This has a minimal impact on the reported LVE properties (see Supporting Information). (21) Cooper-White, J. J.; Mackay, M. E. J. Polym. Sci., Part B: Polym. Phys. 1999, 37 (15), 1803−1814. (22) Trinkle, S.; Friedrich, C. Rheol. Acta 2001, 40 (4), 322−328. (23) Zirkel, A.; Urban, V.; Richter, D.; Fetters, L. J.; Huang, J. S.; Kampmann, R.; Hadjichristidis, N. Macromolecules 1992, 25 (23), 6148−6155. (24) Yamane, H.; Sasai, K.; Takano, M.; Takahashi, M. J. Rheol. 2004, 48 (3), 599.

(8) (a) Guillaume, S. M.; Kirillov, E.; Sarazin, Y.; Carpentier, J. F. Chem. - Eur. J. 2015, 21 (22), 7988−8003. (b) Hillmyer, M. A.; Tolman, W. B. Acc. Chem. Res. 2014, 47 (8), 2390−2396. (c) Liu, H. Z.; Zhang, J. W. J. Polym. Sci., Part B: Polym. Phys. 2011, 49 (15), 1051−1083. (d) Dijkstra, P. J.; Du, H. Z.; Feijen, J. Polym. Chem. 2011, 2 (3), 520−527. (e) Buffet, J. C.; Okuda, J. Polym. Chem. 2011, 2 (12), 2758−2763. (f) Thomas, C. M. Chem. Soc. Rev. 2010, 39 (1), 165− 173. (g) Stanford, M. J.; Dove, A. P. Chem. Soc. Rev. 2010, 39 (2), 486−494. (h) Carpentier, J. F. Macromol. Rapid Commun. 2010, 31 (19), 1696−1705. (9) (a) Stopper, A.; Okuda, J.; Kol, M. Macromolecules 2012, 45 (2), 698−704. (b) Heck, R.; Schulz, E.; Collin, J.; Carpentier, J. F. J. Mol. Catal. A: Chem. 2007, 268 (1−2), 163−168. (c) Arnold, P. L.; Buffet, J. C.; Blaudeck, R.; Sujecki, S.; Wilson, C. Chem. - Eur. J. 2009, 15 (33), 8241−8250. (d) Arnold, P. L.; Buffet, J. C.; Blaudeck, R. P.; Sujecki, S.; Blake, A. J.; Wilson, C. Angew. Chem., Int. Ed. 2008, 47 (32), 6033−6036. (e) Otero, A.; Fernandez-Baeza, J.; Lara-Sanchez, A.; Alonso-Moreno, C.; Marquez-Segovia, I.; Sanchez-Barba, L. F.; Rodriguez, A. M. Angew. Chem., Int. Ed. 2009, 48 (12), 2176−2179. (f) Normand, M.; Kirillov, E.; Roisnel, T.; Carpentier, J. F. Organometallics 2012, 31 (4), 1448−1457. (g) Drouin, F.; Whitehorne, T. J. J.; Schaper, F. Dalton Trans. 2011, 40 (6), 1396− 1400. (h) Buffet, J.-C.; Okuda, J.; Arnold, P. L. Inorg. Chem. 2010, 49, 419−426. (i) Buchard, A.; Carbery, D. R.; Davidson, M. G.; Ivanova, P. K.; Jeffery, B. J.; Kociok-Kohn, G. I.; Lowe, J. P. Angew. Chem., Int. Ed. 2014, 53 (50), 13858−13861. (j) Dove, A. P.; Li, H. B.; Pratt, R. C.; Lohmeijer, B. G. G.; Culkin, D. A.; Waymouth, R. M.; Hedrick, J. L. Chem. Commun. 2006, No. 27, 2881−2883. (k) Zhang, L.; Nederberg, F.; Pratt, R. C.; Waymouth, R. M.; Hedrick, J. L.; Wade, C. G. Macromolecules 2007, 40 (12), 4154−4158. (l) Chen, H. L.; Dutta, S.; Huang, P. Y.; Lin, C. C. Organometallics 2012, 31 (5), 2016−2025. (m) Wang, H. B.; Yang, Y.; Ma, H. Y. Macromolecules 2014, 47 (22), 7750−7764. (n) Isotacticity can be quantified by the probability of meso linkages (Pm) in the polymer. A completely isotactic polymer would have Pm = 1. (10) (a) Palade, L.; Lehermeier, H. J.; Dorgan, J. R. Macromolecules 2001, 34, 1384−1390. (b) Dorgan, J. R.; Janzen, J.; Clayton, M. P.; Hait, S. B.; Knauss, D. M. J. Rheol. 2005, 49 (3), 607−619. (c) Dorgan, J. R.; Janzen, J.; Knauss, D. M.; Hait, S. B.; Limoges, B. R.; Hutchinson, M. H. J. Polym. Sci., Part B: Polym. Phys. 2005, 43, 3100. (d) Hirata, M.; Kimura, Y. Polymer 2008, 49, 2656. (e) Othman, N.; AcostaRamirez, A.; Mehrkhodavandi, P.; Dorgan, J. R.; Hatzikiriakos, S. G. J. Rheol. 2011, 55 (5), 987−1005. (f) Othman, N.; Xu, C.; Mehrkhodavandi, P.; Hatzikiriakos, S. G. Polymer 2012, 53 (12), 2443−2452. (11) (a) Ovitt, T. M.; Coates, G. W. J. Am. Chem. Soc. 1999, 121 (16), 4072−4073. (b) Buffet, J.-C.; Kapelski, A.; Okuda, J. Macromolecules 2010, 43 (24), 10201−10203. (c) Kapelski, A.; Okuda, J. J. Polym. Sci., Part A: Polym. Chem. 2013, 51, 4983−4991. (12) (a) Chamberlain, B. M.; Cheng, M.; Moore, D. R.; Ovitt, T. M.; Lobkovsky, E. B.; Coates, G. W. J. Am. Chem. Soc. 2001, 123 (14), 3229−3238. (b) Chisholm, M. H.; Gallucci, J.; Phomphrai, K. Chem. Commun. 2003, No. 1, 48−49. (c) Chisholm, M. H.; Gallucci, J. C.; Phomphrai, K. Inorg. Chem. 2004, 43 (21), 6717−6725. (d) Zhong, Z. Y.; Dijkstra, P. J.; Feijen, J. J. Am. Chem. Soc. 2003, 125 (37), 11291− 11298. (e) Hormnirun, P.; Marshall, E. L.; Gibson, V. C.; White, A. J. P.; Williams, D. J. J. Am. Chem. Soc. 2004, 126 (9), 2688−2689. (f) Amgoune, A.; Thomas, C. M.; Roisnel, T.; Carpentier, J. F. Chem. Eur. J. 2006, 12 (1), 169−179. (g) Bonnet, F.; Cowley, A. R.; Mountford, P. Inorg. Chem. 2005, 44 (24), 9046−9055. (h) Coulembier, O.; Mespouille, L.; Hedrick, J. L.; Waymouth, R. M.; Dubois, P. Macromolecules 2006, 39, 4001−4008. (i) Chmura, A. J.; Chuck, C. J.; Davidson, M. G.; Jones, M. D.; Lunn, M. D.; Bull, S. D.; Mahon, M. F. Angew. Chem., Int. Ed. 2007, 46 (13), 2280−2283. (j) Liu, X.; Shang, X.; Tang, T.; Hu, N.; Pei, F.; Cui, D.; Chen, X.; Jing, X. Organometallics 2007, 26 (10), 2747−2757. (k) Pietrangelo, A.; Hillmyer, M. A.; Tolman, W. B. Chem. Commun. 2009, No. 19, 2736− 2737. (l) Pietrangelo, A.; Knight, S. C.; Gupta, A. K.; Yao, L. J.; Hillmyer, M. A.; Tolman, W. B. J. Am. Chem. Soc. 2010, 132 (33), K


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