Microstructure of Crystallizable α-Olefin Molecular ... - ACS Publications

Article Cite This: Macromolecules 2018, 51, 872−883

Microstructure of Crystallizable α‑Olefin Molecular Bottlebrushes: Isotactic and Atactic Poly(1-octadecene) Carlos R. López-Barrón,* Andy H. Tsou, Jarod M. Younker, Alexander I. Norman, Jonathan J. Schaefer, John R. Hagadorn, and Joseph A. Throckmorton ExxonMobil Chemical Company, 5200 Bayway Drive, Baytown, Texas 77520, United States S Supporting Information *

ABSTRACT: Isotactic and atactic poly(1-octadecene) (iPOD and aPOD) have been synthesized by organometallic coordinative insertion polymerization of 1-octadecene. Analyzing X-ray and neutron scattering data of POD melts identifies their bottlebrush structures as flexible rods where the rod length is the extended backbone length and rod radius is the side chain coil dimension. Upon cooling, both iPOD and aPOD melts crystallize by fully extending their coiled side chains to form orthorhombic alkane crystals in iPOD and nematically ordered rotator alkane crystals in aPOD, as determined by X-ray scattering and Raman spectroscopy. Molecular dynamics simulations of isotactic and atactic 48-mers of 1-octadecene were applied to define and verify melt and crystalline structures and scattering peak assignments, respectively. Modeling suggests that side chains of both crystallized isotactic and atactic PODs align at 70° and 160° to the 4/1 spiral backbone of equal probability, at an average of 115°, and POD chains pack in an antiparallel pattern. Large wheat-sheaf structural assembly of fibril bundles can be observed in aPOD, which render high opacity to these samples. Each of those fibrils is made of several bottlebrush molecules packed into a hexagonal lattice. Faster crystallization observed in iPODs hinders the formation of large crystallites, which results in translucent samples.

■

INTRODUCTION Recent advances in polymerization allow the synthesis of precisely spaced comb-branched polymers with relatively narrowly dispersed comb lengths.1−7 By reducing backbone spacer length between the comb branch points, a crossover from combs to bottlebrush architecture occurs.8,9 This crossover starts when the coil dimension of the comb arm reaches the unperturbed spacer size. Once the comb arm’s coil size is greater than the backbone spacer size, these side chains would force the backbone to be in a fully extended state8 and hinder the interpenetration of neighboring molecules.9 A corresponding X-ray diffraction peak arises from neighboring backbone correlations,10 indicating the development of a bottlebrush structure. Most bottlebrush polymers reported in the literature have been synthesized by “grafting onto” or “grafting from” synthetic approaches, based on controlled radical polymerization or ring-opening metathesis polymerization (ROMP).1−5 Previous reports of homopolymerization of macromers (via the so-called graft-through method) utilized anionic polymerization11 and ROMP2 to synthesize bottlebrush polymers. Recently, it was demonstrated that a bottlebrush hydrocarbon polymer can be prepared simply by organometallic coordinative insertion polymerization of vinyl-terminated polyolefin macromers.10 The lowest molecular weight macromer used in ref 10 was vinyl-terminated atactic polypropylene with Mn = 120 g/mol. A review of various © 2018 American Chemical Society

synthetic approaches to prepare molecular bottlebrushes can be found elsewhere.12 Instead of using vinyl-terminated polyolefin macromers, one can simply use higher linear α-olefin monomers to prepare bottlebrush polymers, provided that the alkyl chain length is greater than six (i.e., 1-heptene and above). With insertion polymerization, the backbone spacer will be two carbons, and poly(1-hexene) would have four-carbon pendant side chains. The square root of 4 (coil dimension of the side chain) is 2, and poly(1-hexene) is expected to be a crossover polyolefin from combs to bottlebrushes. It was reported that the melting temperature decreases from isotactic polypropylene to isotactic poly(1-hexene) and then increases again and keep on increasing from isotactic poly(1-heptene) to isotactic poly(1octadecene).13−16 In addition to the melting point minimum of isotactic polyolefins being at poly(1-hexene), the specific volume maximum and thermal expansion and theta temperature minima are also at this composition.15 Since the melting point minimum at poly(1-hexene) corresponds to the changeover from helical backbone crystallization to parallel side chain crystallization, it was suggested16 that a critical side chain length is reached preventing backbone from participation in Received: November 29, 2017 Revised: January 9, 2018 Published: January 24, 2018 872

DOI: 10.1021/acs.macromol.7b02524 Macromolecules 2018, 51, 872−883

Article

Macromolecules Table 1. Reaction Temperature and Molecular Characteristics of Atactic and Isotactic Polyoctadecene sample

max reaction temp, °C

Mn, kg/mol

Mw, kg/mol

Mw/Mn

mm, %

mr, %

rr, %

Tm, °C

Tc, °C

ΔHm, J/g

aPOD1 aPOD2 aPOD3 iPOD1 iPOD2

117 115 71 45 45

17 52 149 245 260

54 113 352 691 732

3.17 2.17 2.36 2.82 2.81

27.2 25.9 21.7 82.1 91.0

24.7 35.7 38.6 10.5 3.7

48.1 38.4 39.7 7.4 5.3

41.9 41.9 41.9 62.0 60.8

38.3 38.6 37.9 54.4 53.5

87.3 65.6 66.7 58.3 62.1

rhombic alkane crystals in iPODs and rotator alkane crystals in aPODs. Molecular dynamics simulation was applied to analyze the SAXS/WAXD scattering patterns of iPOD and aPOD and to determine their crystalline structures. Further arrangements of these crystallites into higher level structures can only be seen in atactic PODs with “wheat-sheaf” assemblies found by AFM (atomic force microscopy). This study demonstrates a very important distinction between crystalline structures of low αolefin polymers (such as polyethylene, polypropylene, and poly(1-butene)) and α-olefin polymer bottlebrushes. The former crystallize by backbone folding to produce lamellar structures, and the latter crystallize by arrangement of side chains from adjacent bottlebrush chains.

crystallization. In other words, poly(1-hexene) corresponds to the crossover of a polyolefin from a comb to a bottlebrush, from interpenetrating chains to noninterpenetrating chains,9 and from flexible backbones to extended backbones. Crystalline structures of isotactic poly(α-olefins) have been examined earlier on polyolefins synthesized using Ziegler− Natta (ZN) titanium chloride catalysts.16 These polyolefins contained significant amounts of atactic components (40− 50%17). Extraction in boiling heptane was employed to separate atactic polyoctadecene from isotactic polyoctadecene.17,18 Thus, extracted atactic polyoctadecene can crystallize by selfassembly of their side chains into ordered rotator phases.17 More recently, C2 symmetric metallocene catalysts activated with methylaluminoxane have been employed to prepare isotactic polyolefins from poly(1-butene) to poly(1-octadecene).14,19 Only thermal properties of these metallocenecatalyzed isotactic polyolefins were reported; however, the examination of their melt and crystalline structures is lacking. The crystalline structure of ZN-catalyzed isotactic polyoctadecene (iPOD), in the presence of crystallizable atactic polyoctadecene (aPOD), has been thoroughly examined by Turner-Jones16 using wide-angle X-ray diffraction (WAXD), and two different crystalline forms, type I and II, were reported in nonpurified iPOD depending on thermal history. It was proposed that both crystalline types have backbones aligned in plane with fully extended side chains which are at right angles to the backbone in type I and tilted by 130.6° with respect to the backbone in type II. Both crystalline types have orthorhombic unit cells. Crystalline iPOD of type II is denser and thermally more stable than those of type I and potentially can have two arrangements which cannot be differentiated by X-ray scattering. One is parallel type II where side chains of neighboring backbones are tilting in the same direction but with 4/1 spiral backbone organizing in up and down fashion. The other is antiparallel type II with side chains of neighboring backbones tilting up and down but with neighboring backbones spiraling in the same direction. Unfortunately, as recognized by VanderHart and co-workers,17 the high content of aPOD in these iPOD samples prepared by ZN catalysis greatly complicates their structural analysis. It follows that synthesis of pure iPODs and aPODs is necessary in elucidating microstructures present in these bottlebrush macromolecules. In this study, we synthesized iPOD and aPOD by organometallic coordinative insertion polymerization of 1octadecene. Their bottlebrush dimensions, as flexible rods, and scattering patterns in the melt were measured by X-ray and neutron scattering20 and compared with computed bottlebrush structures. Light and X-ray scattering, along with Raman spectroscopy and rheological measurements, were utilized to follow melting and crystallization of iPOD and aPOD, and their crystalline structures were established by X-ray scattering and Raman spectroscopy. It was found that coiled side chains in iPODs and aPODs were converted to fully extended side chains prior to crystallizing, which allowed them to form ortho-

■

EXPERIMENTAL SECTION

Synthesis. Atactic polyoctadecene (aPOD) polymers were prepared by the batch polymerization of 1-octadecene in toluene solution using the catalyst formed by a combination of Cs symmetric metallocene (diphenylmethylene)bis((1,2,3,3a,7a-η)-1H-inden-1ylidene))dimethylhafnium with 1 mol equiv of N,N-dimethylanilinium tetrakis(pentafluorophenyl) borate activator. The reactor was first cleaned with bis(diisobutylaluminum) oxide scavenger, and the polymerization was run adiabatically by controlling only the initial reactor temperature. All aPOD polymerizations were run for 2 h with their maximum reactor temperatures tabulated in Table 1. Afterward, polymers were recovered, dried, and characterized by GPC-IR (gel permeation chromatography−infrared detector, Polymer Char) for molecular weight, by 13C NMR for tacticity, and by differential scanning calorimetry (DSC) for crystallization and melting temperatures (Tc, and Tm, respectively). These characterization results are listed in Table 1. Isotactic polyoctadecene (iPOD) polymers were prepared by the batch polymerization of 1-octadecene in hexane solution using the catalyst formed by the combination of the C 1 symmetric pyridyldiamide complex (N-(2,6-bis(1-methylethyl)phenyl)-6-(2((cyclopentylamino-κN)methyl)-1-naphthalenyl)-α-(2-(1methylethyl)phenyl)-2-pyridinemethanaminato(2-)-κN1,κN2)dimethylhafnium with 1 mol equiv of N,N-dimethylanilinium tetrakis(pentafluorophenyl) borate activator. The reactor was first cleaned with bis(diisobutylaluminum) oxide scavenger, and the polymerization was run adiabatically with a starting temperature of 23 °C. The polymerization for iPOD1 ran for 3.2 h with half the octadecene and catalyst loadings (relative to iPOD2), whereas the polymerization for iPOD2 ran for 2 h with a full loading of octadecene and catalyst. Afterward, polymers were recovered, dried, and characterized by GPCIR for molecular weight and by 13C NMR for tacticity. Their reaction temperatures, molecular weight, tacticity values, and melting temperatures are listed in Table 1. In-Situ X-ray Scattering. A Linkam CSS-450 optical stage with a Kapton window was used to heat up and cool down iPODs and aPODs, at 1 °C/min, for X-ray scattering measurements. X-ray measurements were carried out using the 12-BM-B beamline at Advanced Photon Source (Argonne National Laboratory). The energy of the X-ray is 12 keV with a 0.103 nm X-ray wavelength and a 500 μm × 500 μm beam size. Two-dimensional (2D) SAXS and WAXD patterns were collected using Mar 225 and Mar 165 charged coupled device (CCD) detectors, respectively, with an exposure time of 0.5 s and read-out time of approximately 2 s per frame. SAXS and WAXS 873

DOI: 10.1021/acs.macromol.7b02524 Macromolecules 2018, 51, 872−883

Article

Macromolecules patterns were calibrated for q space by using standards silver behenate and alumina, respectively. The 2D images to 1D plots were performed over a 0−360° integration, and corrections were made for detector background and air scattering backgrounds. Crystallinity was initially determined using a series of peak fitting routines to compute the fraction of (hkl) reflections relative to the total area under the WAXS trace. As discussed below, at higher temperatures this proved problematic, and the iPOD bulk crystallinity from the 1D SAXS correlation function was performed. For the aPOD sample, WAXS fits of the rotator phase peak (to a Lorentzian function) were carried out, and the area under this curve was compared to the area of the amorphous halo (fitted to a Pearson VII function). This computes a value of the % of rotator phase “crystal”. It is not a crystal in the true physical sense but has a degree of order akin to a nematic liquid crystal phase. For the purpose of clarity throughout this paper we simply refer to this as an ordered packed rotator phase. Raman and Light Transmission Measurements. Raman measurements were performed using a confocal Raman microscope CRM alpha 300 (Witec) equipped with a RayShield coupler. Raman spectra were produced using a 785 nm excitation wavelength and 60 s integration time on a CCD camera (ANDOR). Temperaturedependent Raman spectra of both iPOD and aPOD samples were generated on a Linkam Testing Stage model TST350 (Linkam), using a heating/cooling rate of 1 °C/min. The same Linkam stage was used to collect light transmittance measurements, during heating and cooling at 1 °C/min, using a homemade light scattering instrument. Dynamic Mechanical Thermal Analysis (DMTA). Dynamic temperature ramp measurements were performed using a straincontrolled rheometer ARES-G2 (TA Instruments) using 8 mm serrated plates over a temperature range from 25 to 50 °C for the aPODs and from 25 to 80 °C for the iPODs. All measurements were perfomed at a heating/cooling rate of 1 °C/min, using a frequency ω = 1 rad/s and 1% strain. Atomic Force Microscopy (AFM). Morphologies of aPODs and iPODs were examined by a bimodal AFM (Cypher, Asylum Research) after cryo-facing using a cryo-microtome (Leica) at −120 °C. Bimodal AFM, where the cantilever-tip ensemble is simultaneously excited at two eigenmodes, can deliver enhanced contrast especially when the second-mode phase images are used.21,22 In applying bimodal AFM to polymers, repulsive tip−sample interaction is necessary, which requires optimization of oscillation amplitude and set point for each mode. Based on numerical simulation and experimental findings, appropriate parameters for operating bimodal AFM were identified to minimize the contrast inversion artifacts.23 These optimized bimodal AFM operating conditions were employed to acquire AFM phase morphologies in this study. Molecular Dynamics Simulation. Simulation cells are composed of varying numbers of POD chains where each polymer chain is composed of 48 monomers (either isotactic or atactic). Molecular dynamic simulations utilized the united atom Trappe force field24 with Large-scale Atomic/Molecular Massively Parallel Simulator (LAMMPS)25 as the engine (using constant N/P/T, isobaric and isothermal, time step 1 fs). Three different atom types of CH (tertiary backbone carbon), CH2 (secondary backbone or side chain carbon), and CH3 (primary side chain carbon) were designated. For melt simulations, four chains were packed into an amorphous cell using Scienomics MAPS 4.0. Following an initial geometry optimization, these cells were heated at 600 K and 1 atm for 10 ns, then linearly cooled to 347 K (isotactic) or 323 K (atactic) for 10 ns, and equilibrated at 347 or 323 K for 20 ns. The production run, from which statistics and trajectory frames were gathered, continued for an additional 10 ns. For crystalline simulations, 8 initial POD chains were constructed (4 isotactic and 4 atactic) with 4 different backbone torsions (30°, 45°, 60°, and 90°). Each chain was placed in a simulation cell with the head and tail carbons connected across the periodic cell (z direction), in essence creating a chain of infinite MW. The cell was replicated along the other two dimensions to create cells composed of either 4 (2 × 2 cell) or 9 (3 × 3 cell) chains. Chains were manually translated to remove steric clashes and to minimize the free volume with cell dimensions changed accordingly. Following an initial

geometry optimization, the triclinic cell was compressed at 100 K and 10 MPa (x, y) or 0 MPa (z) for 1 ns. The cell was then heated at 600 K and 1 atm (x, y, z) for 10 ns, linearly cooled to 295 K for 10 ns, and equilibrated at 295 K for 20 ns. The production run, from which statistics and trajectory frames were gathered, continued for an additional 10 ns. For the 2 × 2 chain models, we observed at times >∼ 30 ns a significant energy drop, with a decrease in density by ∼3%. This correlated with the increase in cell length of one of the dimensions (with an equal decrease of length along one or both of the remaining dimensions). We initially assumed this was an artifact of the violation of the minimum energy convention (i.e., a box length had decreased to less than 2 times of the pairwise cutoff value of 12 Å). Therefore, the frame at 26 ns was replicated along the smallest cell dimension to create an eight-chain model, which was equilibrated at 295 K and 1 atm (x, y, z) for 15 ns. The production run, from which statistics and trajectory frames were gathered, followed for an additional 10 ns. However, unlike most other molecular dynamics engines, LAMMPS does not use a minimum energy convention; instead, ghost atoms with box lengths added/subtracted are used to calculate the pair energies. Thus, cutoffs can be greater than the dimensions of the simulation cell. Torsion distributions, intermolecular radial distribution functions, and molecular visualizations were performed with Scienomics MAPS 4.0. X-ray scattering predictions were done using BioVia Materials Studio 2017R2 (cutoff 100 Å) by averaging six equally spaced frames from the production runs. A total of 12 MD models were examined, and they are listed in Table 2.

Table 2. MD Models Evaluated in This Study model

initial backbone

no. of chains

model

initial backbone

no. of chains

1a 2a 3a 4a 5 6

30 45 60 90 30 45

4 4 4 4 9 9

1b 2b 3b 4b 7 8

30 45 60 90 60 90

8 8 8 8 9 9

■

RESULTS AND DISCUSSION Bottlebrush Structures of APOD and IPOD Melts. X-ray scattering profiles of iPOD2 and aPOD3 melts, above their

Figure 1. X-ray scattering patterns of aPOD3 and iPOD2 melts.

melting temperatures, are shown in Figure 1. The same scattering patterns can be found in iPOD1, aPOD1, and aPOD2. The broad peak at q value around 1.33−1.36 Å−1, corresponding to a characteristic distance of 0.46−0.47 nm, is 874

DOI: 10.1021/acs.macromol.7b02524 Macromolecules 2018, 51, 872−883

Article

Macromolecules

that PODs synthesized in this study can be described as rods with low flexibility.20 In a recent study,20 we showed that upon crystallization SANS data in the high q region remain unaffected, indicating no changes in cross-sectional area and stiffness of these bottlebrush PODs. This is in agreement with the previous observation that the radius of gyration of polymer chains does not change during crystallization.26,27 However, an upturn in the low q SANS profiles appears after crystallization, suggesting formation of large scattering objects bigger than radii of gyration of individual chains. Fittings of the SANS data of PODs in the crystalline state using Beaucage model28,29 provided the radius of gyration of crystalline aggregates, Rg,c, which is proportional to the MW of aPOD, as shown in Table 3. Frames and predicted X-ray scattering patterns from molecular dynamics simulations of the amorphous melts composed of 48-monomer iPOD and aPOD chains are shown in Figure 2. The high q peak at 1.5 Å−1 reflects a slighter shorter van der Waals contact between carbon atoms of side chain segments (0.42 nm versus 0.46 nm measured/shown in Figure 1). This is the result of alkane side chain (CH2) sigma pair coefficient being equal to 3.95 Å in the TraPPE force field. The distinct low q peak at 0.25 Å−1, observed in the X-ray scattering patterns (Figure 1), is not well resolved from the MD simulations; however, it can be seen for the iPOD with low intensity. The distribution of distances between neighboring backbones is shown in Figure S1 of the Supporting Information, which shows that the distribution of backboneto-backbone distances (CH−CH) is broad at >10 Å and why no distinct low q feature can be observed. Crystallization of iPOD and aPOD. Cooling of iPOD and aPOD melts leads to crystallization of both. SAXS/WAXS patterns as a function of temperature during crystallization are shown in Figure 3. The intensity of the backbone spacing peak at q = 0.25 Å−1 (d ∼ 2.5 nm) of iPOD starts to decrease upon crystallization at 71 °C and disappears completely at 68 °C. A new set of scattering peaks at q = 0.17 Å−1 (d ∼ 3.7 nm) and q ∼ 0.025 (d ∼ 250 nm) appear at 71 °C. As for the aPOD sample, upon crystallization at 44 °C, the backbone-tobackbone peak at q = 0.27 Å−1 remains at the same position and a new peak emerges at q = 0.085 Å−1 (d ∼ 7.3 nm). The low q peaks appearing in both iPOD and aPOD upon crystallization indicate the formation of new and rich

Table 3. Chain Dimensions of aPODs and iPOD1 melt

solid

sample

R, nm

L, nm

b, nm

Rg,POD, nm

Rg,c, nm

aPOD1 aPOD2 aPOD3 iPOD1

1.77 1.71 1.75 1.83

5.56 34.7 155 328

4.51 4.55 4.48 4.81

4.98 6.77 13.3 22.3

75.5 77.1 90.2

the so-called amorphous halo found in alkanes, long chain paraffins, and polyethylene and can be attributed to van der Waals interactions between carbon atoms of different chain segments in random coil configuration. In molecular bottlebrush PODs, this high q scattering peak can only arise from side chains since neighboring backbones are far apart.16 The low q scattering peak, centered at q values around 0.25−0.27 Å−1, corresponds to the distance between the neighboring POD backbones (2.4−2.5 nm). It has been reported16 that this peak does not appear in polyolefins with olefin carbon number lower than six (i.e., it starts to appear in poly(1-hexene)), and its relative intensity to the high q peak (at 1.36 Å−1) decreases with increasing olefin carbon number, reflecting the decreasing backbone-to-backbone correlations as the side chain length increases. Additionally, it was reported16 that this low q scattering peak position shifts to lower q positions in polyolefins from poly(1-hexene) to POD due to further separation of bottlebrush backbones. We have found that this backbone spacing is proportional to the square root of the side chain length, or the side chain coil dimension, from polyoctene to POD as expected for bottlebrush polymers.8,9 Bottlebrush melts of aPOD and iPOD can be considered as flexible cylinders whose lengths, L, correspond to their backbone contour length, and their constant radii, R, correspond to their side chain coil dimension. SANS measurements of the aPOD and iPOD melts, reported elsewhere,20 were utilized to measure their chain dimensions as a confirmation of their bottlebrush structures. Table 3 lists the chain dimensions, including L, R, Kuhn length, b, and radius of gyration, Rg, obtained by SANS measurements.20 As expected, the three aPODs have the same bottlebrush radius. The slightly higher radius shown for iPOD1 may be due to the fact that this sample was measured at higher temperature (80 °C) than the aPOD samples (60 °C). A linear dependence of the radius of gyration on the molecular weight can be found, which suggests

Figure 2. Final frames from the amorphous iPOD (left) and aPOD (middle) molecular dynamics simulations with atoms colored by chain. Right: predicted X-ray scattering. 875

DOI: 10.1021/acs.macromol.7b02524 Macromolecules 2018, 51, 872−883

Article

Macromolecules

Figure 3. X-ray scattering patterns of (a) aPOD3 and (b) iPOD2 during cooling at 1 °C/min.

Figure 4. Normalized Raman spectra of iPOD1 showing changes in structure when cooling over the temperature range of 80−28.1 °C. (a) Raman spectra with spectral regions identified and arrows depicting changes in band intensity with decreasing temperature. (b) Raman spectra of CH2 bending region showing orthorhombic crystalline band at 1420 cm−1.

Figure 5. (a) Normalized Raman spectra of iPOD2 and aPOD1 in the CH2 Bending region between 1400 and 1500 cm−1. (b) iPOD orthorhombic crystallinity quantified from the CH2 bending peak at 1420 cm−1.

are transformed from coils to extended chains upon crystallization and that crystallization involves only side chains. Additionally, orthorhombic alkane crystal peaks of (110) and

mesophases. Results from temperature-dependent X-ray scattering profiles and Raman spectra during crystallization of iPOD and aPOD suggest that side chains of iPOD and aPOD 876

DOI: 10.1021/acs.macromol.7b02524 Macromolecules 2018, 51, 872−883

Article

Macromolecules Table 4. All-Trans Mass Conformers Fraction and Crystallinities for POD Samples, Measured at Room Temperature crystallinity (mass fraction) iPOD2 aPOD1 aPOD2 aPOD3 a

all-trans (mass fraction)

Raman

SAXS

0.38 0.45 0.41 0.43

0.29a N/A N/A N/A

0.29b 0.57c

Orthorhombic crystallinity. bBulk crystallinity. cRotator crystallinity.

(200) at q positions corresponding to distances of 0.43 and 0.375 nm, respectively, appear on top of the amorphous side chain scattering peak at 0.46 nm of iPOD (see Figure 3) upon crystallization, whereas a rotator alkane crystal peak at 0.41 nm emerges on top of amorphous scattering peak of aPOD. This suggests that the side chains of iPOD crystallize into orthorhombic crystals and side chains of aPOD arrange into rotator crystals, as first reported by VanderHart et al.17 Temperature-dependent Raman spectra were collected from iPOD2, aPOD1, aPOD2, and aPOD3 using a Linkam stage that was heated at 1 °C/min. An overlaid temperature series of Raman spectra from iPOD1 is presented in Figure 4, which was

Figure 7. SAXS correlation function curves of iPOD2 over a heating cycle.

cooled from 80 to 28.1 °C. Only prominent regions in the temperature-dependent Raman spectra are shown in Figure 4a, including the terminal C−C stretch, C−C stretch, CH2 twist, CH2 bend, and CH2 stretch regions. Crystalline and amorphous bands are superimposed on each other as can be seen in Figure

Figure 6. Melting and crystallization aPOD3 (left panel) and iPOD2 (right panel) during heating/cooling cycles measured at 1 °C/min using five different probes (from top to bottom): Raman spectroscopy, SAXS/WAXS, light transmittance, rheology, and DSC. Photographs show aPOD1 and iPOD2 at the indicated temperatures. 877

DOI: 10.1021/acs.macromol.7b02524 Macromolecules 2018, 51, 872−883

Article

Macromolecules

Figure 10. X-ray scattering profile of crystalline aPOD3 measured at 22 °C.

sample presented in Figure 5b was quantified over the temperature range of 26.2−80 °C using the method reported by Strobl and Hagedorn.31 They proposed a three-phase model for polyethylene, which consists an orthorhombic crystalline, isotropic amorphous, and an anisotropic disordered phase. Temperature-dependent Raman spectra were normalized the area of the CH2 twist vibration at 1295 cm−1a band with an area that is temperature independent. After normalization to the CH2 twist, the area of the orthorhombic peak at ∼1420 cm−1 was divided by 0.46, which is the area of an extended chain polyethylene reported by Strobl.31 Temperature-dependent orthorhombic crystallinity in iPOD1 ranged from 29% crystallinity at 26.2 °C to zero crystallinity at 71.9 °C in the melt. Orthorhombic crystallinity decreased gradually over the temperature range from 26.2 to 71.9 °C, suggesting limited long-range cooperativity of the iPOD1 chains during melting. Upon cooling from the melt there is a sharp jump in orthorhombic crystallization at 55.1 °C, suggesting some longer range cooperativity between side chains during crystallization. In other words, cooperative interactions between more ordered trans chains influence the conformation of neighboring chains as seen at 55.1 °C in Figure 5b with a sharp increase in orthorhombic crystallinity during crystallization. Below 53.2 °C, the crystallization process slows with temperature, and there is less cooperative motions between the chains. All-trans and amorphous conformer fractions were calculated from temperature-dependent Raman spectra of iPODs and aPODs. POD Raman spectra were normalized to the CH2 twist at 1295 cm−1, and a weighted normalized molten spectrum of polyethylene (160 °C) was subtracted to remove the amorphous phase contribution. The all-trans phase contribution from the PODs was quantified by fitting the crystalline CH2 twist peak at 1295 cm−1 after the weighted subtraction of the molten polyethylene spectrum. A weighted subtraction of molten polyethylene was used due to the all-trans C−C stretching bands of the POD samples showing intensity at 1063 and 1130 cm−1 at temperatures above melting, which suggests that the backbone may not be completely disordered. As indicated in Table 4, all aPODs contain an all-trans conformer mass fraction ranging from 0.41 to 0.45 regardless of their molecular weights. This suggests that the all-trans conformer mass fraction content in an aPOD may be independent of its molecular weight. The all-trans conformer

Figure 8. AFM micrographs of aPOD1 (upper panel) and iPOD1 (lower panel).

Figure 9. X-ray scattering profile of crystalline iPOD2 measured at 22 °C.

4a with arrows indicating changes in normalized Raman scattering intensity as the molten iPOD1 is cooled from a temperature of 80 °C. Temperature-dependent Raman spectra of both iPOD and aPOD were normalized to the area of the CH2 twist vibration with a peak maximum at 1295 cm−1 using methods developed for polyethylene analysis.30,31 The CH2 twist area has been found to be invariant with temperature for polyethylene and was used as an internal standard to quantify the all-trans and amorphous phase conformers in the PODs. As shown in Figure 4b, Raman spectroscopy confirms the presence of an orthorhombic crystalline phase in the iPOD1. The Raman peak located at 1420 cm−1 in the iPOD1 spectrum of Figure 4b is associated with orthorhombic unit cell and caused by crystal field splitting of the fundamental vibration at 1440 cm−1. Furthermore, the orthorhombic crystalline phase was not detected in any of the aPODs examined by Raman as depicted in Figure 5a. Orthorhombic crystallinity of the iPOD1 878

DOI: 10.1021/acs.macromol.7b02524 Macromolecules 2018, 51, 872−883

Article

Macromolecules

Figure 11. Hierarchical microstructure of wheat-sheaf crystallites in aPOD samples.

Figure 12. Along the chain (left) and down the chain (right) views of the final frames from crystalline iPOD molecular dynamics simulations with atoms colored by chain. Top: iPOD model 3a. Middle: iPOD model 3b. Bottom: iPOD model 6.

mass fractions of iPOD2 and aPOD1 as a function of temperature are plotted in Figure 6. Examination of the alltrans conformer fraction in iPOD2 shows a gradual decline in the all-trans conformers with increasing temperature, which is similar to the evolution of orthorhombic crystallinity for the iPOD2 (also shown in Figure 6). A more surprising result is the melting behavior of aPOD2 given in Figure 6. The aPOD1 melted over a much smaller temperature range and had less of a pretransition to melting than that of iPOD2, which indicates

greater long-range cooperation between the aPOD chains that between iPOD chains. The other aPODs showed similar trends in melting behavior with the only small differences in melting temperature. The crystallinity of the POD samples was also quantified by SAXS and WAXS. The all-trans conformer mass fraction and crystallinities obtained by Raman and SAXS are given in Table 4. Upon crystallization of aPOD3, a single Bragg peak appears on top of an amorphous halo (see Figure 3a), which resembles 879

DOI: 10.1021/acs.macromol.7b02524 Macromolecules 2018, 51, 872−883

Article

Macromolecules

Figure 13. Along the chain (left) and down the chain (right) views of the final frames from crystalline aPOD molecular dynamics simulations with atoms colored by chain. Top: aPOD model 2a. Top middle: aPOD model 4a. Bottom middle: aPOD model 4b. Bottom: aPOD model 5.

the diffraction peaks observed in rotator phases in alkanes.17,18,32 The relative area of this Bragg peak with respect to the area of the amorphous halo is used to compute the fraction of ordered rotator phase, which is plotted as a function of temperature in Figure 6. This is illustrated in Figure S2, which shows representative results at four temperatures. WAXS scattering profiles of iPOD2 at several temperatures were fit to a series of crystalline and amorphous peaks with representative fitting results shown in Figure S3. At 22 °C, the degree of crystallinity from WAXS fits is 29.3%, consistent with the finding from Raman spectroscopy. However, as the temperature increases, the broadening of these weak crystalline reflections, while still present, become so broad (see Figure S3) that the fits become unreliable. The uncertainty in mathematical peak fitting is thus raised in analyzing WAXS profiles at higher temperatures. Hence, instead of using WAXS scattering profiles, bulk crystallinity values of iPOD1 were calculated from the Fourier-transformed Lorentz-corrected SAXS data. These thus-obtained correlation function curves, which are essentially damped Patterson function curves, are

plotted in Figure 7. Gamma represents the real space correlation function, given by Gamma(R ) =

1 Qs

∫0

∞

(I(q)q2 cos(qR )) dq

where I(q) is the scattering intensity and Qs is the experimental invariant obtained from the Lorentz-corrected (I(q)·q2) 1D SAXS profile scattering between the experimental limits of q1 (the first real data point) and q2 (where I(q) becomes constant). The correlation function assumes an ideal two-phase lamellae morphology and can be used to extract long period, crystalline block thickness, amorphous block thickness, and estimated crystallinity. The second maximum in a given correlation function curve (Figure 7) is used to extract its long spacing as the curve crosses the Gamma = 0 line with an assumption that this broad SAXS peak is an electron density contrast between a two-phase system consisting of crystalline and amorphous regions with a sigmoidal interface in between. Subsequent analysis for crystalline and amorphous block 880

DOI: 10.1021/acs.macromol.7b02524 Macromolecules 2018, 51, 872−883

Article

Macromolecules

Figure 14. Predicted X-ray scattering pattern for (a) iPOD and (b) aPOD models. Inset in (a) shows measured SAXS/WAXS scattering profile of iPOD2 plotted against computed scattering profile from iPOD simulation model 6. Inset in (b) shows measured SAXS/WAXS profile of aPOD3 plotted against computed scattering profile of aPOD simulation model 5.

chains allow them to fold into lamellar structures, whereas the high rigidity in POD bottlebrushes forbid them from bending enough to acquire the chain folded lamellae structure. Rather, the crystallization in both iPOD and aPOD is via cooperative ordering processes of the side chains, similar to alkane crystallization.33 The X-ray scattering profile of iPOD2, shown in Figure 9, has all the scattering peaks expected from the so-called type II crystalline form first proposed by Turner-Jones,16 including the two orthorhombic crystal peaks of (110) and (200). Characteristic distances computed as d = 2π/q*, where q* is the position of the diffraction peaks, are shown in Figure 9, along with a table listing all the Miller indices of Turner-Jones type II diffraction peaks. In the proposed type II form, side chains are tilted by 130.6° to the backbone 4/1 helical axis and can be up and down among neighboring backbones or in the same direction side-by-side. We did not see type I crystalline form and suspect that its appearance in Turner-Jones’ iPOD may come from the aPOD contamination. The POD examined by Turner-Jones, which was synthesized using a ZN catalyst without purification, contained a large amount of atactic components. An additional very broad peak centered at q ∼ 0.026 Å−1 (corresponding to d ∼ 240 Å) is observed in the iPOD2 sample but not reported by Turner-Jones.16 This peaks arises from correlation between small crystallites in iPOD, which are evident in Figure 8, where distances between hard crystallites (dark domains) are in the range of 10−60 nm. The X-ray scattering profile of aPOD3 is illustrated in Figure 10. The narrow sharp peak at q*R = 1.51 Å−1, which resembles the rotator phase observed in alkane crystals,32 describes the local arrangement between aPOD side chains. Unlike the side chains in iPOD bottlebrushes, which crystallize into orthorhombic unit lattices, the side chains in aPOD do not have crystalline lattice registration. Rather, these side chain trans conformers pack tightly as nematic liquid crystals where side chain axes align in the director direction, but they lack lateral order, analogous to the rotator phase in alkanes. The characteristic side chain distance decreases from 0.47 nm (in the melt) to 0.41 nm upon crystallization. We postulate that the crystallization of the aPOD samples is driven by the formation of this nematically ordered “rotator” phase of the side chains, which in turn produces rearrangements of the whole

thicknesses can then also be used to determine local and bulk crystallinity values. Analyzing SAXS data takes away the uncertainty in WAXS peak fitting which is problematic in iPOD due to the lower q broader peaks, as first reported by Tuner-Jones.16 However, this SAXS analysis is dependent on the two-phase crystal−amorphous stack model with sigmoidal interfaces between the two phases, whereas WAXS analysis is not dependent on this assumption. Orthorhombic crystallinity values of iPOD2 as a function of temperature during heating and cooling computed from SAXS profiles and Raman spectra as well as rotator crystallinity values of aPOD1 obtained from WAXS profiles are shown in Figure 6. This figure also shows temperature-dependent all-trans conformer, light transmittance, dynamic mechanical moduli, and DSC heat flux values. The melting temperature difference between iPOD and aPOD (see Figure 6 and Table 1) can be attributed to their different alkane crystalline unit cell. It should be noted that despite the lower crystallization temperature, aPOD is opaque upon crystallization, while iPOD remains transparent (see photographs in Figure 6). Additionally, aPOD has a higher solid modulus that can be established right after crystallization. Both of these characteristics of aPOD can be explained by the tighter crystalline structural assembly observed in aPOD compared to iPOD. This is evidenced by the higher mass density (determined picnometrically) measured at 25 °C in aPOD3 (0.895 kg/m3) compared to that measured in iPOD2 (0.852 kg/m3). Note that the melt densities, measured at 82 °C, in both samples are practically identical (0.762 kg/m3). Representative AFM micrographs of aPOD1 and iPOD1 are shown in Figure 8. A “wheat-sheaf” type of crystallite assembly can be seen in aPOD whereas only tiny crystallites can be found in iPOD. The opacity in aPOD is due to the large crystallite sizes (>1 μm), which produces significant light scattering. Contrarily, the crystalline domains in iPOD are less than 100 nm in size, and they are well distributed throughout (i.e., not clustered). No crystallite assembly at sizes greater than 100 nm can be found in iPOD, which explain its macroscopic translucency. Crystalline Structures of aPOD and iPOD. It is worth noting that due to strong steric effects between side chains, the POD chains do not crystallize in the way polyethylene (PE) or polypropylene (PP) does. The high flexibility of PE and PP 881

DOI: 10.1021/acs.macromol.7b02524 Macromolecules 2018, 51, 872−883

Article

Macromolecules

illustrate the packing with the embedded cartoons in the Figures 13 and 14. For the 4-, 8-, and 9-chain models, different arrangements result. In the 8- and 9-chain models we observe packing such that the side chains from different chains lie on top of one another, and the maximum distance between backbones is the length of the side chain. A closer look at the intermolecular RDFs between CH3 and CH to quantify this distance can be found in Figure S7. The same is true for iPOD model 3a and aPOD model 2a, except we observe two flattened chains stacking on top of each other prior to side chain stacking. In iPOD model 4a, we observe linear stacking of the flattened chains. The predicted Xray scattering patterns obtained from the iPOD and aPOD models are shown in Figure 14. For the iPOD models, a high q peak from 1.48 to 1.62 Å−1 and a low q peak at 0.11−0.13 Å−1 can be found. iPOD model 3a features a split high q peak. As for the predicted aPOD scattering patterns, a high q peak at 1.57 Å−1 which is broader and less intense for the 4-chain models is indicated. The predicted low q peak for the 8and 9-chain models is located at q ∼ 0.12 Å−1, while the 4-chain models show two low q peaks at 0.09 and 0.17 Å−1 for aPOD model 2a and at 0.10 and 0.21 Å−1 for aPOD model 4a. The inset in Figure 14a shows a comparison between the computed scattering pattern of an iPOD (model 6) and that of experimentally measured. Although they are similar, differences in peak positions are found, more specfically at lower q scattering. Considering that 100% isotacticity, infinite molecular weight, and monodispersed molecular weight were assumed during iPOD MD modeling, these reductions in peak positions, or increases in spacing between chains, are expected from model predictions. Comparison of the computed diffraction pattrern for aPOD using model 5 with the measured X-ray diffraction (inset of Figure 14b) shows shifting of the low q peak to higher q values, when the shifting toward lower q is expected. Given that this peak arises from interfiber−bundle correlations, as discussed above, which is not a feature built in the MD models, it was not expected that the MD simulation would predict this peak.

bottlebrush molecules. These rearrangements manifest as additional diffraction peaks at lower q values, which are discussed next. The amorphous backbone scattering peak located at a q value of 0.27 Å−1 (labeled as q*Am in Figure 10, where “Am” stands for amorphous) in aPOD3 remains in the solid state, although with lower intensity. This indicates that significant amorphous fraction remains when aPOD is cooled below Tc. An additional peak centered at a q value of ∼0.85 Å−1 is observed in Figure 10. We postulate that this arises from the distance between the fibrils forming the wheat-sheaf crystallites and is labeled as q*Fib, where “Fib” stands for fibrils. We hypothesize that these fibrils, which are made of individual bottlebrush chains bundled together, are separated by an average distance of dFib = 2π/q*Fib ∼ 7.4 nm. The hierarchical structure of the aPOD wheat-sheaf crystallites is illustrated in Figure 11, which shows that the bottlebrush molecules within each of the fibrils arrange into a hexagonal (HEX) cylinder phase. This configuration is supported by the diffraction peaks centered at q values of 0.45, 0.78, and ∼0.90 Å−1, which correspond to the HEX cylinder space group with relative spacings q/qH* = 1, √3, √4, and √7, where q*H represents the position of the principal reflection.34 Note that the last reflection peak at √7q*H is not well-defined due to overlapping with the peak centered at qR*, but nevertheless, it is resolvable. The lattice parameter of the hexagonal structure, a = 1.61 nm, which represents the distance between the centers of adjacent cylinders, is obtained from the position of the first diffraction peak, qH*, as a = 4π/√3qH*. This distance a is smaller than the average distance between bottlebrush backbones in the melt (dBB_melt = 2.33 nm). Molecular dynamics simulations were applied to understand crystalline structures of iPOD and aPOD solids. Total energy distributions of the final 10 ns of the molecular dynamics simulations for the different models of cyrstalline iPOD and aPOD are shown in Figure S4. Energies can only be compared if the simulations have an identical number and types of atoms. Therefore, at least three models will be discussed. The lowest energy iPOD models are 3a, 3b, and 6 (see Table 2). The lowest energy aPOD models are 2a, 4a, 4b, and 5 (models 5 and 8 converged to similar end points). Along the chain and down the chain orientations of the final frames of these simulations are illustrated in Figures 12 and 13 for iPOD and aPOD, respectively. A visual analysis of the trajectories for the lowest-energy models finds a consistent motif for the polymer chain with minimal interpentatration of the side chains, which is in agreement with recent results from theoretical analyses and simulation results of bottlebrush moleucles.9 Instead, each bottlebrsuh molecule is a flattened rod with on average three chains per 360° turn of the backbone or four monomers per one turn (4/1 helical). An analysis of the distribution of backbone torsion angles (CH−CH2−CH−CH2 in Figures S5 and S6) reveals two different distributions of equal probability: ∼70° and ∼160°, whose average is 115°, or ∼120° per turn (4/ 1 helical). Excluding the first and second side chain carbons, the side chains are in a trans configuration with torsions of ∼180° as shown in Figures S5 and S6. The second side chain carbon is primarily at ∼180°, but with a small distribution at ∼75°. However, the first side chain carbon, which dictates the alignment of the side chain to the backbone, has two equally probable torsion angles: ∼70° and ∼160° with an average of 115°. The packing of the flattened rods is governed by the periodic boundary conditions of the simulations. We endevoured to

■

CONCLUSIONS Isotactic and atactic poly(octadecene) bottlebrush polymers were synthesized by organometallic coordinated insertion polymerization of 1-octadecene. These bottlebrush iPODs and aPODs are flexible rods with their rod lengths corresponding to their backbone extended lengths while having a constant rod diameter dictated by their side chain coil dimension. Upon cooling, both iPOD and aPOD melts crystallize by extending side chain through transformation to all-trans conformers followed by side chain assembly either into orthorhombic alkane crystals in iPODs or nematic-like rotator crystals in aPODs. Molecular dynamics simulation suggests that side chains of both crystallized isotactic and atactic PODs align at an average of 115° to the 4/1 spiral backbone and pack in an antiparallel pattern. The aPOD bottlebrush molecules assembly into a hierarchical structure consisting of bundles of bottlebrush rods packed into a hexagonal lattice, which form fibrils that further assemble into large wheat-sheaf crystallites. These wheat-sheaf crystallites scatter the light and render aPOD opaque. No crystallite assembly at sizes greater than 100 nm can be found in iPOD, which explain its macroscopic translucency. 882

DOI: 10.1021/acs.macromol.7b02524 Macromolecules 2018, 51, 872−883

Article

Macromolecules

■

(15) Wang, J.; Porter, R. S.; Knox, J. R. Physical Properties of the Poly(1-olefin)s. Thermal Behavior and Dilute Solution Properties. Polym. J. 1978, 10, 619−628. (16) Turner-Jones, A. Crystallinity in isotactic polyolefins with unbranched side chains. Makromol. Chem. 1964, 71, 1−32. (17) VanderHart, D. L.; Pérez, E.; Bello, A.; Vasquez, J.; Quijada, R. Effect of Tacticity on the Structure of Poly(1-octadecene). Macromol. Chem. Phys. 2004, 205, 1877−1885. (18) Aubrey, D. W.; Barnatt, A. Transition temperatures and crystallinity in polyoctadecene-1. Journal of Polymer Science Part A-2: Polymer Physics 1968, 6, 241−248. (19) Brüll, R.; Pasch, H.; Raubenheimer, H. G.; Sanderson, R.; Wahner, U. M. Polymerization of higher linear α-olefins with (CH3)2Si(2-methylbenz[e]indenyl)2ZrCl2. J. Polym. Sci., Part A: Polym. Chem. 2000, 38, 2333−2339. (20) López-Barrón, C. R.; Tsou, A. H.; Hagadorn, J. R.; Zeng, Y.; Kang, S.; Mattler, S. J. Chain Dimensions of α-Olefin Molecular Bottlebrushes in the Melt and Solution. In preparation, 2017. (21) Martinez, N. F.; Patil, S.; Lozano, J. R.; Garcia, R. Enhanced compositional sensitivity in atomic force microscopy by the excitation of the first two flexural modes. Appl. Phys. Lett. 2006, 89, 153115. (22) Garcia, R.; Proksch, R. Nanomechanical mapping of soft matter by bimodal force microscopy. Eur. Polym. J. 2013, 49, 1897−1906. (23) Chakraborty, I.; Yablon, D. G. Cantilever energy effects on bimodal AFM: phase and amplitude contrast of multicomponent samples. Nanotechnology 2013, 24, 475706. (24) Chen, B.; Potoff, J. J.; Siepmann, J. I. Monte Carlo Calculations for Alcohols and Their Mixtures with Alkanes. Transferable Potentials for Phase Equilibria. 5. United-Atom Description of Primary, Secondary, and Tertiary Alcohols. J. Phys. Chem. B 2001, 105, 3093−3104. (25) Plimpton, S. Fast Parallel Algorithms for Short-Range Molecular Dynamics. J. Comput. Phys. 1995, 117, 1−19. (26) Fischer, E. W. Neutron Scattering Studies on the Crystallization of Polymers. Polym. J. 1985, 17, 307−320. (27) Fischer, E. W. Studies of structure and dynamics of solid polymers by elastic and inelastic neutron scattering. Pure Appl. Chem. 2009, 50, 1319. (28) Beaucage, G. Approximations Leading to a Unified Exponential/ Power-Law Approach to Small-Angle Scattering. J. Appl. Crystallogr. 1995, 28, 717−728. (29) Beaucage, G. Small-Angle Scattering from Polymeric Mass Fractals of Arbitrary Mass-Fractal Dimension. J. Appl. Crystallogr. 1996, 29, 134−146. (30) Lagaron, J. M.; Dixon, N. M.; Reed, W.; Pastor, J. M.; Kip, B. J. Morphological characterisation of the crystalline structure of colddrawn HDPE used as a model material for the environmental stress cracking (ESC) phenomenon. Polymer 1999, 40, 2569−2586. (31) Strobl, G. R.; Hagedorn, W. Raman spectroscopic method for determining the crystallinity of polyethylene. J. Polym. Sci., Polym. Phys. Ed. 1978, 16, 1181−1193. (32) Sirota, E. B.; King, H. E. J.; Shao, H. H.; Singer, D. M. Rotator Phases in Mixtures of n-Alkanes. J. Phys. Chem. 1995, 99, 798−804. (33) Anwar, M.; Turci, F.; Schilling, T. Crystallization mechanism in melts of short n-alkane chains. J. Chem. Phys. 2013, 139, 214904. (34) Alexandridis, P.; Olsson, U.; Lindman, B. A Record Nine Different Phases (Four Cubic, Two Hexagonal, and One Lamellar Lyotropic Liquid Crystalline and Two Micellar Solutions) in a Ternary Isothermal System of an Amphiphilic Block Copolymer and Selective Solvents (Water and Oil). Langmuir 1998, 14, 2627−2638.

ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.7b02524. Figures S1−S7 (PDF)

■

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] (C.R.L.-B.). ORCID

Carlos R. López-Barrón: 0000-0002-9620-0298 Notes

The authors declare no competing financial interest.

■

ACKNOWLEDGMENTS Experimental support provided by Beth Welke, Hillary Passino, Sarah Mattler, and Shuhui Kang is greatly acknowledged. We thank Eric Sirota of ExxonMobil Research and Engineering Company for his valuable insights.

■

REFERENCES

(1) Beers, K. L.; Gaynor, S. G.; Matyjaszewski, K.; Sheiko, S. S.; Möller, M. The Synthesis of Densely Grafted Copolymers by Atom Transfer Radical Polymerization. Macromolecules 1998, 31, 9413− 9415. (2) Li, Z.; Zhang, K.; Ma, J.; Cheng, C.; Wooley, K. L. Facile syntheses of cylindrical molecular brushes by a sequential RAFT and ROMP “grafting-through” methodology. J. Polym. Sci., Part A: Polym. Chem. 2009, 47, 5557−5563. (3) Jha, S.; Dutta, S.; Bowden, N. B. Synthesis of Ultralarge Molecular Weight Bottlebrush Polymers Using Grubbs’ Catalysts. Macromolecules 2004, 37, 4365−4374. (4) Xia, Y.; Olsen, B. D.; Kornfield, J. A.; Grubbs, R. H. Efficient Synthesis of Narrowly Dispersed Brush Copolymers and Study of Their Assemblies: The Importance of Side Chain Arrangement. J. Am. Chem. Soc. 2009, 131, 18525−18532. (5) Zhang, M.; Müller, A. H. E. Cylindrical polymer brushes. J. Polym. Sci., Part A: Polym. Chem. 2005, 43, 3461−3481. (6) Pesek, S. L.; Li, X.; Hammouda, B.; Hong, K.; Verduzco, R. Small-Angle Neutron Scattering Analysis of Bottlebrush Polymers Prepared via Grafting-Through Polymerization. Macromolecules 2013, 46, 6998−7005. (7) Dalsin, S. J.; Hillmyer, M. A.; Bates, F. S. Molecular Weight Dependence of Zero-Shear Viscosity in Atactic Polypropylene Bottlebrush Polymers. ACS Macro Lett. 2014, 3, 423−427. (8) Paturej, J.; Sheiko, S. S.; Panyukov, S.; Rubinstein, M. Molecular structure of bottlebrush polymers in melts. Sci. Adv. 2016, 2, e1601478. (9) Liang, H.; Cao, Z.; Wang, Z.; Sheiko, S. S.; Dobrynin, A. V. Combs and Bottlebrushes in a Melt. Macromolecules 2017, 50, 3430− 3437. (10) López-Barrón, C. R.; Brant, P.; Eberle, A. P. R.; Crowther, D. J. Linear rheology and structure of molecular bottlebrushes with short side chains. J. Rheol. 2015, 59, 865−883. (11) Pantazis, D.; Chalari, I.; Hadjichristidis, N. Anionic Polymerization of Styrenic Macromonomers. Macromolecules 2003, 36, 3783− 3785. (12) Yuan, J.; Müller, A. H. E.; Matyjaszewski, K.; Sheiko, S. S. In Polymer Science: A Comprehensive Reference; Matyjaszewski, K., Möller, M., Eds.; Elsevier: Amsterdam, 2012. (13) Turner-Jones, A.; Aizlewood, J. M. Crystalline forms of polypentene-1. J. Polym. Sci., Part B: Polym. Lett. 1963, 1, 471−476. (14) Kaminsky, W. In Handbook of Polymer Synthesis; Kricheldorf, J. S., Porter, R. S., Knox, J. R., Eds.; Marcel Dekker: New York, 2004. 883

DOI: 10.1021/acs.macromol.7b02524 Macromolecules 2018, 51, 872−883