Effect of Self-Poisoning on Crystallization Kinetics of Dimorphic

Feb 9, 2018 - The incident X-ray beam was the Cu Kα line with a wavelength λ = 1.5418 Å. The samples were isothermally crystallized from the melt i...
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Article Cite This: Macromolecules XXXX, XXX, XXX−XXX

Effect of Self-Poisoning on Crystallization Kinetics of Dimorphic Precision Polyethylenes with Bromine Xiaoshi Zhang,† Wei Zhang,† Kenneth B. Wagener,‡ Emine Boz,‡ and Rufina G. Alamo*,† †

Department of Chemical and Biomedical Engineering, FAMU-FSU College of Engineering, 2525 Pottsdamer St., Tallahassee, Florida 32310-6046, United States ‡ The George and Josephine Butler Polymer Research Laboratory, Department of Chemistry, University of Florida, Gainesville, Florida 32611-7200, United States S Supporting Information *

ABSTRACT: High molar mass polyethylenes with bromine atoms placed on each and every 21st, 19th, 15th, or 9th backbone carbon crystallize into two distinctive layered polymorphs by changing undercooling. Crystallization at low temperatures produces Form I, a planar all-trans conformation, while at higher temperatures gauche conformers set for backbone bonds adjacent to the methine due to a close intermolecular staggering of bromines resulting in a herringbone Form II structure. In this work, the sharp range of isothermal crystallization temperatures for the transition between Form I and Form II is first identified via WAXD and melting behaviors for all members of the series. Furthermore, the temperature dependence of the isothermal linear spherulitic growth rates of Form II has been studied for a wide range of crystallization temperatures. The linear growth rates display a discrete minimum with decreasing temperature at a crystallization temperature near the melting point of Form I, a feature which is reminiscent of the minimum found in the crystallization rate of long-chain n-alkanes. Changes in spherulitic morphology and the growth rate minima are analyzed on the basis of self-poisoning at the growth front resulting from frequent but unstable Form I depositions on the growth surface of Form II. The similarity with the behavior observed in the growth of long-chain n-alkanes crystallites supports a polymer crystallization process controlled by events that take place at the crystal growth front.

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INTRODUCTION Polyethylene-like molecules with equidistant placement of a counit along the backbone are interesting models to study the effect of the size and polarity of the counit on chain folding and on the various modes of intermolecular staggering of counits and methylene runs during crystallization.1 In a series of prior works, we reported the synthesis and thermal properties of precision polyethylenes containing halogens that were synthesized via olefin metathesis polycondensation.2−6 In the initial works, the thermal and crystalline properties were analyzed for specimens that were cooled from the melt at a relatively fast rate. By detailed spectroscopic (FTIR, 13C NMR) and crystallographic studies, it was demonstrated that these systems crystallize as homopolymers accommodating 5−9 repeating units in lamellar crystallites with the halogens staggered in layers.1,2,5,6 More recently, we have started to analyze polymorphism and differences in structural and thermal properties between fast and isothermally crystallized samples.7,8 Precision systems with Cl or Br pendant groups are of special interest because they undergo a sharp polymorphic transition with increasing crystallization temperature. Under fast crystallization, Cl- and Br-containing samples adopt Form I, a planar © XXXX American Chemical Society

all-trans conformation of the backbone carbons, while in slowly crystallized samples a close intermolecular staggering of halogens leads to a nonplanar (kinked) or herringbone-type structure (Form II). In the latter, gauche conformers set for bonds of the backbone carbons adjacent to the carbon with the halogen, while the methylene sequence between halogens maintains the all-trans conformation.7−9 In other words, in Form II the all-trans methylene sequences bend back and forth around the carbon with the halogen. Evidence of the crystals transitioning from packing in Form I to Form II with decreasing undercooling was given by large differences in melting points, a change in the X-ray diffraction pattern, and a change in conformation of the backbone bonds vicinal to the halogen substitution point.7 The change from a planar all-trans conformation to the nonplanar herringbone type by changing crystallization kinetics was extracted from quantitative FTIR analyses of the content of crystalline gauche conformers around the halogen substitution and by a detailed analysis of the CH2 Received: December 27, 2017 Revised: February 2, 2018

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DOI: 10.1021/acs.macromol.7b02745 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules rocking progression modes in reference to the spectra of alltrans n-alkanes of equivalent length. WAXD patterns combined with FTIR analysis are powerful structural characterization tools for precision polyethylene-like systems of this nature, as recently shown for series of halogen-containing precision samples.7 Such drastic differences in chain packing between Form I and Form II anticipate equally drastic differences in nucleation barriers. The barrier of the kinetically favored all-trans Form I is expected to be much lower than the corresponding barrier for Form II with the herringbone structure. Some evidence that points to this difference was given in a previous work by polarized optical microscopy that shows much higher nucleation density for Form I than for Form II, thus indicating that Form I is easier to initiate.8 However, up to date isothermal crystallization kinetics have not been studied. In the present work, we analyze the temperature coefficient of the isothermal linear growth rates of the series of precision polyethylenes with Br placed on each 21st, 19th, 15th, or 9th backbone carbons. As will be shown, the gradient of the growth rates of most of these systems displays an unusual deep minimum with decreasing temperature which is reminiscent of the minimum observed in the crystallization rate of long-chain n-alkanes and in a methylterminated low molar mass PEO fraction.10−12 In both systems, the lamellar thickness undergoes a quantized change from extended to an integer folded form (once-folded, twice-folded, etc.) at the temperature where the rate is at a minimum. For long-chain n-alkanes not only is there a sharp increase in the variation of the growth rate with temperatures at the transition between extended and once-folded crystallization mode, but the growth rate of the extended form decreases sharply at temperatures near the melting of the once-folded structure when approached from above.10 This unusual crystallization feature was prevalent in the crystallization of all n-alkanes with a length above the critical length for folding (>150 carbons) crystallized from the melt or from solution. The deep minimum found in the crystallization kinetics of nalkanes raised interest because while the classical Hoffman− Lauritzen (HL) theory could describe the lamellar growth and predicted morphological behaviors of most semicrystalline polymers, it is unable to explain the minimum of the growth rate observed in n-alkanes.13 The extra lateral free energy term added to the HL theory in an attempt to explain the rate minimum raised serious criticism.14−16 The growth rate minimum was explained by Ungar and Keller with the concept of “self-poisoning”.10,11,16,17 In essence, self-poisoning implies that while the n-alkane crystal is growing in the extended-chain form, segments with the folded conformation will attach to the growth front hindering further growth until they detach and growth can proceed in the extended form. The attachment and detachment of the wrong conformation takes some time and causes retardation in the propagation of further growth. The reason the minimum is observed at temperatures above but near the melting of the folded crystals is that at these temperatures once-folded attachment events on the extended growth front are not negligible, but growth is only productive in the extended conformation. How deep the observed minimum of the growth rate is depends on the relative rates of attachment and detachment which are mandated by the difference between the melting points of the extended (TmE) and the once-folded (TmF2) forms. Because the difference TmE − TmF2 is larger in short alkanes, than for larger ones, the lifetime of wrong

attachments for the short molecules is short and the poisoning effect negligible compared to the pronounced G minima observed for long-chain n-alkanes.16,18 From morphological observations, Ungar et al. argued that self-poisoning could be explained by the surface roughening theory of Saddler and Gilmer19−21 and on this basis predicted that self-poisoning should be present in the crystallization of polymers. Selfpoisoning was not found in polymer crystallization; hence, it was also argued that it will be hard to detect in polymers due to the polydispersity in chain length, and because unlike the quantized lamellar thickness of long-chain n-alkanes, the lamellae thickness of polymers changes continuously with Tc.16 Following arguments for surface roughening, the minimum was also predicted via simulations of chain extension of the folded form that attaches to the extended growth front.22 On these grounds, the manifestation of self-poisoning as a deep growth rate minimum with decreasing Tc would not be expected in precision polyethylenes as they are polydisperse and of high molecular weight. A more general treatment of self-poisoning has been recently given by Whitelam et al. for crystal growth of any type of materials.23 They posited that self-poisoning is ubiquitous in any crystalline molecule provided the molecule can attach in at least two different forms to the crystal front. The molecules can attach with the same or with a different crystal conformation. Restrictions also imposed are that the two forms are not energetically equivalent and that the probabilities of binding to the crystal for each form are also significantly different. In other words, these authors modeled the generic case of poisoning at the rough-growth front by a conformation that does not quite match the more stable conformation but that by itself can generate a different stable structure. Although this mean-field simulation emphasizes conformation of the binding segment rather than segment length, is not applicable to the n-alkane situation because it allows both conformations to be part of the growing crystal, one as a defect. It is known that extended nalkanes crystals cannot grow from the folded structure.16 The behavior of n-alkanes may just be a special case of a general polymer lamellar growth driven by events that take place at the growth front. The requirement for monodisperse, relatively short molecules and a mismatch of segment lengths that approach a growing surface, used to explain the observed rate minima of n-alkanes, may not need to be the major factors to observe self-poisoning in crystallization of polymers. Regardless of the length, if the conformation of the segment that binds to a growing surface is somewhat different than at the crystal front, and growth cannot proceed, poisoning will prevail, and the manifestation will be a more or less pronounced minimum of the growth rate. A similar situation should apply to the primary nucleation event. Hence, on the basis of “sensing” the conformation at the growth front, the kinetic trap associated with poisoning could be explained by differences in nucleation and growth kinetics of two possible chain structures when their existence is feasible. The series of precision polyethylenes with halogens are systems that fit quite well with the above postulated premises for the manifestation of self-poisoning of dimorphic high molecular weight polymers. Precision systems containing Cl and Br can develop two different structures with different thermal stabilities and chain conformations. Furthermore, the high nucleation density of Form I indicates a lower energy barrier than to form nuclei of Form II. These features parallel the properties of once-folded and extended long-chain nB

DOI: 10.1021/acs.macromol.7b02745 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules Table 1. Molecular Mass and Thermal Characterization of Precision Br-Containing Polyethylenes sample

mol % (Br)a

Mw × 103 (g/mol)b

Mw/Mnb

Tg (°C)c

Tc,peak (°C)c

Tm,peak (°C)c

ΔHm (J/g)c

PE9Br PE15Br PE19Br PE21Br

11.1 6.67 5.26 4.76

23.6 27.6 38.1 94.1

1.8 1.7 1.7 2.2

−41 ∼−64 ∼−73 ∼−77

−22.0 27.0 49.2 53.1

−13.3 49.0 64.4 70.0

17.5 45.0 67.4 69.4

a

Content of Br, in moles per 100 backbone carbons. bGPC vs PS in THF. cGlass transition temperature (Tg), crystallization peak temperature (Tc,peak), second heating melting peak temperature (Tm,peak), and heat of fusion (ΔHm), obtained by DSC at a rate of 10 °C/min. Estimated Tg are from a linear plot of Tg versus wt % of CHBr using the experimental Tg obtained for PE9Br and Tg of linear polyethylene of −120 °C.24

alkanes.10 With the premise that attachments of Form I may retard the growth of Form II, the requirement for selfpoisoning of binding events occurring with non-negligible probabilities is expected to be met in the temperature range of formation of Form II and at temperatures above but close to the melting temperature of Form I. In this work we first identify the temperature range for the formation of Form I and Form II of a series of precision polyethylenes with bromine by their characteristic X-ray patterns and melting differences. We then analyze the unusual temperature gradient of the isothermal linear growth rates collected in a broad range of temperatures. The implications to explain growth of polymer crystals as a process driven by events that take place at the surface of the lamellae are discussed.

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instrument. The diffractograms were collected at the isothermal crystallization temperature after full transformation. SAXS patterns of selected samples were also collected during melting at 1 °C/min to follow structural changes during fast melt-recrystallization. Polarized optical micrographs were obtained in 40 ± 10 μm films using an Olympus BX51 optical microscope fitted with an Olympus digital camera (Type DP72). Isothermal spherulitic growth was recorded and analyzed using the cellSens software commercialized by Olympus. The sample temperature was controlled using a Linkam hot stage in conjunction with a cryogenic cooling unit (Type FTIR600) with a temperature programmer (Type TMS94), also made by Linkam Scientific Instruments Ltd., UK. Growth rates (G) were obtained from the slope of the linear variation of the radius of the spherulites with crystallization time at a constant crystallization temperature (Tc). At least three different spherulites were measured. The average G and standard deviation were recorded at each Tc. When the morphology changed to axialites, both arms of the axialitic structure were measured independently as a function of crystallization time. The G value from both arms was basically identical. Nucleation rates were measured by counting the number of developing spherulites as a function of crystallization time from frames collected with very short time intervals. The POM images were processed and analyzed using ImageJ (National Institutes of Health).

EXPERIMENTAL SECTION

Materials. Precision polyethylenes with bromine atoms placed on each and every 9th, 15th, 19th, and 21st backbone carbons were synthesized via acyclic diene metathesis (ADMET) followed by exhaustive hydrogenation. Details of the synthesis and chain characterization can be found in previous works.2,3 These precision polyethylenes are labeled PE9Br, PE15Br, PE19Br, and PE21Br, where the number corresponds to the precise location of the bromine atom in the carbon backbone. The repeated structural unit is −[(CH2)x−1− CHBr]n−. The weight-average molecular mass (Mw), molar mass distribution (Mw/Mn), and thermal characterization obtained by DSC are listed in Table 1. Measurements. Isothermal crystallizations were carried out in a TA Q2000 differential scanning calorimeter (DSC) operating under nitrogen flow. The DSC was calibrated for static temperature, thermal lags, and heat of fusion with indium. The TA Q2000 is connected to an intracooler to maximize heat transfer and to allow subambient temperature control. An ∼150 μm thick film from the original powder was first made, and about 4 mg of this film was placed in an aluminum pan without cover. To remove the thermal history, the samples were initially melted at 100 °C for 3 min and cooled at 40 °C/min to the desired Tc. Once the required crystallization times had elapsed, each sample was brought to room temperature and removed from the DSC pan for X-ray testing. After X-ray testing, the isothermally crystallized samples were again encapsulated and melted in the DSC at 10 °C/ min. Wide- and small-angle X-ray diffractograms (WAXD, SAXS) were collected at room temperature or at the crystallization temperature using a Bruker Nanostar diffractometer with IμS microfocus X-ray source and equipped with a HiStar 2D Multiwire SAXS detector and a Fuji Photo Film image plate for WAXD detection. The plate was read with a Fuji FLA-7000 scanner. SAXS and WAXD profiles were calibrated in reference to the patterns of silver behenate and corundum provided by Bruker. The incident X-ray beam was the Cu Kα line with a wavelength λ = 1.5418 Å. The samples were isothermally crystallized from the melt in the DSC at times sufficiently long as to ensure high levels of transformation in each crystalline form. After isothermal crystallization, WAXD patterns of all the precise polyethylenes except for PE9Br were collected at room temperature. PE9Br was isothermally crystallized at different crystallization temperatures in a Peltier device placed inside the sample compartment of the Bruker

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RESULTS AND DISCUSSION

Crystallographic Packing. In prior works it was demonstrated that the crystallization of halogen-containing precision polyethylenes is akin to that of a homopolymer. In forming the crystallites, the chain folds back and forth and includes the halogens in layers that accommodate 5−9 repeating units in the chain axis. However, the kinetic restriction to crystallize by the presence of bromine is obvious by a drastic decrease of crystallization temperatures (Tc) with increasing bromine content. Thus, the crystallization of PE21Br can be followed in a range of temperatures between 40 and 65 °C while PE9Br requires much lower temperatures, between −20 and 10 °C. As shown earlier by data collected at the extreme low and high crystallization temperatures, all Brcontaining precision polyethylenes crystallize in two distinctive crystal forms.7 The crystallization temperature that demarcates the formation of the all-trans planar structure (Form I) and the beginning of the formation of the herringbone structure (Form II) is now identified from X-ray patterns of samples that were crystallized isothermally in a wide range of crystallization temperatures. WAXD patterns of isothermally crystallized PE21Br, PE19Br, PE15Br, and PE9Br are shown in Figure 1. In each panel the red WAXD pattern demarcates the change from crystallization in Form I to Form II. As shown, both forms coexist in a very narrow temperature range. The conformational change between both forms has been previously characterized by FTIR spectroscopy.7 In the four samples studied, Form I develops in the low range of crystallization temperatures. Except for the broad single pattern of PE9Br, the characteristic WAXD patterns of this C

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displacement of the Br in the layers of Form I.9 This displacement is corrected in Form II by intermolecular halogen kinks and the gauche conformation of the C−C bonds adjacent to the methine.7,8 Melting of Form I and Form II. The melting behavior of the same isothermally crystallized samples is given in Figure 2.

Figure 1. WAXD patterns collected at room temperature of isothermally crystallized precision bromine containing polyethylenes. The patterns of PE9Br were collected at the isothermal Tc. The red patterns demarcate crystallization of Form I (low Tc range) and crystallization of Form II (high Tc range). A broad reflection at ∼12° in the patterns of PE19Br is an impurity.

form are two main reflections at 19° and ∼22°. The reflection at 19° is unchanged in the series, while the reflection at ∼22° clearly shifts to lower angles with increasing bromine content. The shift to lower angles is explained as an expansion of the (010) plane proportional to the content of bromine atoms that are incorporated in the triclinic crystalline lattice.2,6 The lattice expansion of Form I is the largest for PE9Br that shows a shift of the 22° reflection to values close to the 19° reflection. Both reflections overlap in a single broad peak as shown by the pattern obtained at Tc = −20 °C in Figure 1d. The WAXD patterns of PE19Br contain an extra broad diffraction at ∼12° from an impurity which does not affect the development of each crystal form. The herringbone structure (Form II) develops at temperatures >52 °C (PE21Br), >45 °C (PE19Br), >25° (PE 15Br), and >−15 °C (PE9Br) and is characterized by additional reflections as seen in the X-ray patterns. The reflection at 2θ = 23° of Form II also shifts appreciably to lower angles with increasing Br, thus indicating an expansion of the crystal lattice proportional to the content of Br included in the crystals, analogous to the expansion effect observed in Form I. In both forms the Br atoms pack in layers, as shown by the prominent low angle layer peak (4° < 2θ < 10°) in the X-ray patterns that corresponds closely to the distance between halogens. A signature of the formation of Form II is also the appearance of second and third orders of the layer peak, reflecting a higher symmetry for this form. From a crystallographic analysis of fiber patterns, it was found that there is some longitudinal

Figure 2. DSC melting thermograms after isothermal crystallization of precision bromine containing polyethylenes at the indicated temperatures. The red thermograms demarcate melting patterns of Form I (low Tc range) and of Form II (high Tc range).

In each panel, the red thermogram also demarcates the transition from Form I crystals in the low Tc range to Form II crystals at the higher crystallization temperatures. As shown, a characteristic behavior of Form I crystals of PE21Br, PE19Br, and PE15Br is their fast melting and recrystallization into Form II on heating. This feature is indicated by a first melting followed by a recrystallization exotherm and further by a melting peak at temperatures where Form II crystals melt. The melt-recrystallization origin of the double melting was further confirmed by melting at different heating rates crystallites formed at the same low temperature. With increasing heating rate the time to recrystallize decreases, and the area of the high melting temperature peak decreases in agreement with a slower recrystallization. Supporting thermograms can be found in Figure SI.1. Crystallites formed in the high range of temperatures where WAXD patterns demonstrate that they belong to Form II melt as single endotherms at higher temperatures. The thermograms of Figure 2 also show that the difference in melting temperatures between Form I (Tm,FI) and Form II (Tm,FI) increases with increasing content of Br. Tm,FII − Tm,FI is only ∼7 °C for PE21Br and increases up to a difference of ∼30 °C for PE9Br. As shown later, this difference in melting D

DOI: 10.1021/acs.macromol.7b02745 Macromolecules XXXX, XXX, XXX−XXX

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not the case as seen by the corresponding thermogram shown in Figure 3 where, for comparison, the melting of crystals

between both forms becomes important to discuss features related to the unusual crystallization kinetics of these systems. The difference in melting reflects drastic differences in thermal stability of both forms as well as differences in their crystallization kinetics. Form I is the kinetically favored form as it develops in the low temperature range, and Form II is thermodynamically more stable crystallizing and melting at higher temperatures. For PE21Br, PE19Br, and PE15Br, Form I melts at temperatures where the crystallization kinetics of Form II are relatively fast; hence, immediately upon melting Form I there is fast recrystallization into Form II, which further melts at temperatures 5−10 °C higher as shown in the thermograms of Figure 2. Because of significantly slower crystallization kinetics, PE9Br is the only member of the series for which Form I does not undergo melting−recrystallization at a heating rate of 10 °C/min. Furthermore, the thermograms of Figure 2 indicate that rapidly formed Form II also undergoes reorganization to crystallites that are of Form II and melt at slightly higher temperature, as shown for PE15Br and PE9Br by double endotherms at ∼50 and ∼15 °C, respectively. These endotherms correspond to Form II for crystallizations at Tc in the transition range between Form I and Form II. Only at the highest crystallization temperatures does the melting of Form II shift at significantly higher values, indicating crystal annealing or the formation of thicker crystallites. It is apparent that the layered packing of these forms undergoes complex reorganizations on heating. Even a highly mobile form with a pseudohexagonal packing has been identified just before melting.9 Notice that should these samples be melted and taken at room temperature to crystallize, their melting temperatures may not be comparable as PE9Br and PE15Br will crystallize in Form II while the other two would do so in Form I. These data make relevant the importance of identifying polymorphism in precision polyethylenes prior to building relationships between their melting temperatures and composition of pendant groups.1 The melting behavior combined with the X-ray patterns of Figure 1 make it obvious that each form has a well-defined temperature range of stability demarcated by the end of the melting peak of Form I. The data also indicate that kinetic factors are relevant and Form I only becomes the dominant form at temperatures well below its melting point (Tm,FI). Let us take as an illustrative example the melting and crystallization range data of PE15Br. For this sample Form I melts at 40 °C (Figure 2c), yet pure Form II is the predominant structure formed above and even 10 °C below Tm,FI as seen in this figure. The formation of Form II well below melting of Form I indicates that although the nucleation barrier of Form II may be significantly higher than for Form I, the drive for crystallization of Form II at high undercooling surpasses slow crystallization kinetics of Form I at low undercooling or at temperatures between 30 and 40 °C for PE15Br. The possibility of a solid−solid transformation of Form I to Form II during heating to explain the double melting was ruled out by the following experiment. PE21Br was first crystallized at 40 °C to develop Form I, and subsequently the temperature was raised to 58 °C and held for 30 min prior to cooling to further register melting of the annealed crystals. As seen in Figure 2a, 58 °C is below the melting of Form I but well above the Tc for the transition from Form I to Form II. Therefore, should metastable Form I be subject to a solid−solid reorganization to Form II, melting of Form I crystals annealed at 58 °C would give only the melting peak of Form II. This is

Figure 3. Melting curves at 10 °C/min of PE21Br crystallized from the melt at Tc = 40 and 58 °C and after crystallization at 40 °C and subsequent annealing for 30 min at 58 °C (top thermogram). After annealing and prior to starting the heating run, the sample was brought to 22 °C at a cooling rate of 40 °C/min.

formed at 40 and 58 °C directly from the melt is also added. The annealing step increased the heat of fusion and sharpness of the endotherm corresponding to Form I, but the characteristic melting−recrystallization of Form I to Form II during heating is unchanged (top thermogram of Figure 3). The annealing results demonstrate that only upon fully melting Form I can Form II recrystallize. The same conclusion about stability of Form I was reached in a study of PE15Cl for which Form I crystals of PE15Cl do not undergo melting− recrystallization.8 The observed melting temperatures for Form I and Form II in the PEXBr series and the crystallization temperature for the transition are plotted as a function of content of bromine in Figure 4. Clearly, Tm FII − Tm FI decreases with increasing distance between bromines atoms or as the content of bromine in the chain decreases. Both melting temperatures merge at ∼2 mol % Br or for a chain with bromines spaced by ∼50 methylenes, indicating that in such long-spaced precision

Figure 4. Observed melting temperatures of Form I (diamonds) and Form II (squares). Observed transition crystallization temperature between Form I and Form II (open circles) and observed minimum of growth rate (open triangles) as a function of content of bromine for precision polyethylenes. Melting of both forms merge at a composition corresponding to a distance between bromines of ∼50 methylenes. E

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Figure 5. Representative polarized optical micrographs for the growth of Form II in the series of precision polyethylenes with bromine at increasing isothermal crystallization temperatures. The isothermal temperature and crystallization times are indicated. Scale bars are 100 μm.

crystallization temperatures as shown in this figure. In general, with increasing crystallization temperature, the number of spherulites appears to decrease consistent with the expected nucleation driven process. Moreover, a closer look at the spherulites of PE21Br obtained at 64 and at 65 °C reflects an unexpected feature. The spherulites that develop at 65 °C after 147 min are larger than those formed at 64 °C even when they were held for a longer time (171 min). In other words, the spherulites grow faster at 65 °C than at 64 °C while the opposite was expected for a nucleation driven crystallization process. This morphological feature is reproducible and already points to an inversion of the temperature gradient of the growth rate. The inversion takes place in a narrow temperature range and is corrected at increasing temperatures, as seen in the micrographs at 67 °C that required much longer times (270 min) for the spherulites to develop about the same diameter. The inversion of spherulitic growth with increasing temperature is more difficult to identify in the optical micrographs of PE15Br, but there is a clear morphological change between Tc of 44.5 and 46 °C. As shown, at 46 °C the spherulites become more open, transitioning into axialites at even higher crystallization temperatures. The morphological change of PE9Br with increasing crystallization temperature is more striking. A change from Tc −8 °C to −4 °C causes a decrease in nucleation density as expected, with spherulites developing higher symmetry as Tc increases. At temperatures in the range of 0−4 °C the nucleation density remains basically constant, but the spherulitic pattern changes to ribbon-like aggregates

systems only the kinetically favored Form I will develop. The implication of the melting-halogen content relation of Figure 4 is that Form II will be increasingly more difficult to develop as the length of the methylene spacer increases beyond 20 methylenes and will not develop in systems with bromine spaced by >50 backbone carbons. Morphological Changes. With the range of isothermal crystallization temperatures for the formation of Form I and Form II demarcated by the data of Figures 1 and 2, we now analyze morphological differences between both forms and the isothermal linear spherulitic growth rates. At any crystallization temperature in the range of formation of Form I, profuse nucleation invariably leads to a fine morphology of very small spherulites or aggregates that pervade over the complete view area, making it impossible to follow growth rates of Form I (examples are given in Figure SI.2). Even at crystallization temperatures in the transition between Forms I and II, the crystallization rates and nucleation density are too high to follow growth by microscopy. Kinetic data in this temperature region are available via DSC or FTIR and will be reported in future works. Measurements of the linear growth are feasible at temperatures below and above Tm,FI where pure Form II develops. Representative polarized optical micrographs for three of the samples studied spanning crystallization temperatures for fast and relatively slow growth are given in Figure 5. Some interesting morphological features become apparent comparing images obtained from the same view area but at different F

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negative temperature gradient of G in a narrow range of Tc. Following the morphological evidence shown in Figure 5, with increasing undercooling (decreasing crystallization temperature) the growth rate first increases, passes through a maximum, which is sharply followed by a minimum, and then increases rapidly with further decreasing temperature. The decrease of the growth rate of PE21Br at 64 °C by more than half of the value obtained at 65 °C is a remarkable retardation just by one degree change in undercooling. In PE19Br the minimum of the growth rate is observed at 60 °C and is significantly shallower than in PE21Br. A similar shallow but reproducible minimum is seen at 44.5 °C for the G data of PE15Br. No clear G minimum but a discontinuity at about 2 °C is observed for the variation of G with temperature for PE9Br. The shallow minimum in the G data of PE15Br and the discontinuity at a Tc ∼ 2 °C for PE9Br coincide with the morphological change observed by optical microscopy. The signature of the G minimum or discontinuity with increasing undercooling is clearer in plots of ln G as a function of temperature (shown in Figure SI.3). The maximum of G at a temperature of −12 °C for PE9Br reflects the dominance of the transport term at very low crystallization temperatures or those that approach the glass transition temperature of PE9Br (Tg = −41 °C). The inversion of the temperature gradient of the growth observed in Figure 6 is reminiscent of the minimum of the crystallization rate of long-chain n-alkanes first observed by Ungar and Keller10,11 and later extensively studied by Ungar and co-workers for pure components and blends of long-chain n-alkanes of different lengths crystallized from the melt and from solution.25−29 The rate minimum was interpreted as the manifestation of self-poisoning. In a framework analogous to the formation of Forms I and II by changing undercooling in precision Br-containing polyethylenes, n-alkanes longer than 150 carbons can crystallize in extended chain or quantized folded structures (once, twice, etc.) with decreasing crystallization temperature. The melting temperature of each of these structures is sharp and increases proportionally to the quantized crystal thickness. In the range of Tc where once-folded and extended chain crystallites are formed, Ungar and co-workers found the occurrence of a deep crystallization rate minimum at a Tc extremely close to the melting of once-folded crystallites and explained the anomalous rate behavior as a feature of selfpoisoning at the growing surface. The rationale is the following. At temperatures above the melting point of once-folded crystals (TmF2) only extended crystallites can develop. However, at crystallization temperatures approaching TmF2 from above, although only extended-chain depositions on the growth front are stable, the lifetime of a folded chain deposition becomes significant. During this short lifetime extended chains cannot grow on the folded chain substrate; hence, lamellar growth is temporarily blocked until the folded chain overgrowth detaches and growth of the extended structure can again be productive.16 The observed retardation of the rate was interpreted as the temporary blocking of the extended chain surface by shorter segments with the wrong structure. Discontinuities of the growth rates found by Kovacs and coworkers in low molar mass fractions of poly(ethylene oxide) at the transition between folded and extended structures were reinterpreted also as the manifestation of self-poisoning.16,30 From the n-alkane studies, it was found that the strength of the retardation effect, or depth of the minimum, increases with length of the n-alkane. Conversely, the minimum becomes

denoting again a change in morphology. Interestingly, at higher crystallization temperatures (Tc > 6 °C), the nucleation density does not decrease as expected. Instead, at Tc = 6 °C the number of nuclei increases drastically compared to the number of spherulites that develop at Tc between 4 and −4 °C for the same crystallization time. At even higher temperatures (Tc > 8 °C), the induction time increases exponentially, but the nucleation density is still very high. The unusual change in nucleation density and macromolecular morphology of PE9Br with increasing crystallization temperature is not due to the formation of a different crystal form at Tc > 4 °C that may have a lower nucleation barrier than for Form II. As demonstrated in Figure 1d, all WAXD patterns collected Tc > −10 °C are identical and of Form II, thus ruling out the formation of additional polymorphic forms. It will be shown below that this unusual slowdown of the primary nucleation rate of PE9Br with increasing undercooling when Tc is changed from 8 to 4 °C is associated with the same events that cause the inversion of the spherulitic growth rate for PE21Br at temperatures between 64 and 65 °C as well as those responsible for the drastic morphological change in PE15Br between Tc of 44 and 46 °C. Linear Growth Rates and Self-Poisoning. The variation of the linear growth rates (G) with increasing crystallization temperature is given in Figures 6a−d. It is first important to

Figure 6. Linear growth rates of Form II as a function of crystallization temperature. The red dashed lines indicate the melting temperature of Form I. Experimental error bars are added.

emphasize that G data could only be obtained for Form II; in other words, the x-axis for each panel of Figure 6 corresponds to crystallization temperatures above the transition and where only Form II develops. The only exceptions are G data for the two lowest crystallization temperatures of PE9Br. The data of PE21Br, PE19Br, and PE15Br show a clear inversion of the G

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methylenes between bromines and only a weak discontinuity is observed at Tc ∼ 2 °C for PE9Br. A notable difference between monodisperse n-alkanes and polydisperse high molecular weight polymers is in the extent of undercooling for crystallization. n-Alkanes barely undercool, meaning that their crystallization kinetics can be followed at temperatures up to 1 °C or less from their sharp melting point. Polymers need to be undercooled to a much larger extent to observe their crystallization. Crystallization at temperatures close to the polymer melting point is not viable, and only at much lower temperatures is the kinetics within a reasonable time frame for observation. The difference in undercooling explains why for n-alkanes the folded structure crystallizes just below its melting point and the extended structure is not observed at Tc below the rate minimum, while for the precision polymers of Figure 6 Form II is still the dominant form below the minimum. For example, PE21Br crystallizes in pure Form II at Tc between 59 and 64 °C which are below Tm FI. As shown in Figures 1 and 2, only after reaching crystallization temperatures of 52−54 °C both forms develop in a time frame within the limits of observation, but growth cannot be measured below 59 °C due to high nucleation density. The kinetically favored Form I is the only structure that develops for PE21Br at temperatures below 52 °C. It is remarkable that only for PE21Br is the rate minimum (or highest effect on the retardation of growth) observed at the melting point of Form I. The G minimum is observed at temperatures that are 2, 4, and ∼10 °C higher than Tm,FI for PE19Br, PE15Br, and PE9Br, respectively. In other words, not only the minimum becomes shallower as mentioned earlier, but the temperature at which self-poisoning is manifested increases with respect to melting of Form I with decreasing methylene run length between bromines. This feature is explained by recalling the increasing difference between the melting points of Form I and Form II shown in Figure 4 with increasing content of bromine in the series. Since Tm,FII is more remote from Tm,FI with increasing bromine, the lifetime of the obstructing deposition of Form I is shorter, explaining the shallower minimum as Br increases in the series. The temperature at which the retardation has the maximum effect depends on the relative ratio between the growth rate of Form II and the rate of deposition of Form I. Above the minimum, deposition of Form I is basically zero. Below the minimum and above the Form I to Form II transition, the growth rate of Form II is much faster than depositions of Form I and the poisoning effect less pronounced. At intermediate temperatures, poisoning may even lead to an inversion of the rate when GFII decreases to levels commensurate with the attachment rate of Form I. As mentioned, above the transition but below Tm,FI, both crystals may coexist; however, the drive for crystallization of Form I (undercooling) is too low, more so with increasing Br, for effective growth and only Form II develops. For PE9Br, the range of temperatures available for self-poisoning is the widest, and the growth at any temperature is so slow that the major manifestation of self-poisoning is in the primary nucleation rather than in growth. This feature is quite prominent in the optical micrographs of Figure 5. The experimental data of Figure 6 point to events at the surface of the growing polymeric lamellae that are general features for the crystallization of polymers and are not specific to the crystallization of monodisperse short chains. The series of precision polyethylenes studied provide among the first example of a deep minimum of the growth rate for high

shallower, spanning a wider temperature range, when the difference in melting temperatures between the two n-alkane crystal structures becomes wider.16 It was then hypothesized that self-poisoning must prevail in the crystallization of polyethylene at any Tc as stems of multiple lengths will approach the growing surface. It turns out that the characteristics of the n-alkane rate minima with changing length or melting of the n-alkane structure are also observed in the growth rate data of the precision polyethylenes analyzed here. The dashed vertical line added in each panel of Figure 6 corresponds to the melting temperature of Form I, taken at the end of the first melting peak in the thermograms of Figure 2. The melting temperature of Form I of PE21Br (Tm,FI = 64 °C) corresponds closely with the temperature at which the minimum of the growth rate is observed; therefore, a similar interpretation based on self-poisoning can be posited to explain this unusual rate inversion. At crystallization temperatures above Tm,FI only Form II is stable as shown, crystallization proceeds in the herringbone Form II structure, and the growth rate follows the usual negative temperature coefficient. This is the range 65 °C < Tc < 71 °C for PE21Br. However, at Tc approaching 64 °C from above, and due to a low nucleation barrier for Form I (the kinetically favored form), the probability of deposition of segments of Form I on the surface of Form II is significant. Since growth of the herringbone structure cannot proceed from the all-trans Form I surface, the growth of Form II is temporarily blocked until the segment with the wrong conformation rearranges to Form II or detaches. Hence, the deep retardation of the rate in these precision systems is explained by a temporal blockage of the growing surface of Form II due to the attachment of the all-trans conformation, as shown schematically in Figure 7. In other words, the role of

Figure 7. Schematics of self-poisoning. During growth of Form II, chains with the all-trans conformation of Form I attach momentarily to the growth front. The temporal blockage retards the growth rate.

folded and extended chain structures in n-alkane self-poisoning is taken by Forms I and II, respectively, in precision Brcontaining polyethylenes. Given the fact that all brominecontaining polyethylenes studied undergo the same transition from Form I to Form II with increasing crystallization temperature (Figures 1 and 2), a similar effect of self-poisoning or retardation of the crystallization rate approaching the melting point of Form I from above is expected. The inversion is indeed observed in the growth rate data of PE19Br and PE15Br and is accompanied by a change in morphology as shown in Figure 5. Moreover, similar to the n-alkane behavior, the minimum becomes shallower with decreasing number of H

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spherulites are associated with primary nuclei. For PE9Br it was impossible to count developing objects with any accuracy in an extended range of crystallization temperatures, and instead we registered the time for the first events to appear in the polarized optical microscope. These data are plotted in Figures 8a,b.

molecular weight polymers, thus lending support for a mechanism for growth of polymeric lamellar crystallites led by events that take place at the growing surface. The need for monodisperse relatively short molecules (such as the long-chain n-alkanes) that crystallize as extended or as integer folded molecules to observe a retardation in the crystallization rate is replaced in precision polyethylenes by the regular placement of a moiety (the Br atom in the systems studied here) along the polyethylene backbone. Provided the moiety is assembled in layers inside the crystal, the possibility of layer packing with more or less drastic differences in conformation will exist, and energetically unequivalent crystal forms will develop with differences in crystallization kinetics and melting. Poisoning at the growing surface of the high melting form at temperatures approaching the low melting form will prevail and may be experimentally observed as in the case studied here. The series of precision polyethylenes with bromines is just one example of a more general feature, as poisoning should be prevalent in most precision polyethylenes that display double melting behavior and develop layered crystallites.1 The peculiarity of the precision polyethylenes with halogens studied here is that they can assemble in two crystalline forms with drastic differences in chain conformation. The change of the conformation adjacent to the Br substitution by changing Tc produces large and easily detectable effects during crystal growth. In addition to long-chain n-alkanes and low molecular weight fractions of poly(ethylene oxide), minima in the crystallization rate with increasing undercooling have been found in other high molecular weight polymers that can also develop two different crystal forms by changing undercooling. These systems include polyesters,31−35 polyamides,36 polyketones,37 syndiotactic polystyrene,38 and isotactic polypropylene.39,40 However, in these systems the rate minimum is observed at the transition temperature between both polymorphs and at temperatures well below the melting points of any of the two polymorphs. Hence, although the effect on the crystallization kinetics is analogous to the inversion found in long-chain nalkanes and in precision Br-containing polyethylenes, the source of the inversion must be different because in the later the G inversion takes place at or above the lowest melting form (see Figure 4). For the polyesters, aromatic polyketone, sPS, and iPP where a rate minimum was also found, instead of the conventional single bell-shape behavior with a maximum in the rate at approximately half way between Tg and Tm, two maxima associated with the bell shape of each polymorph are distinctive, each in a given range of crystallization temperatures. The rate minimum corresponds with the temperature where the two bell-shaped kinetics overlap. For these systems, the minimum is explained by differences in nucleation regimes of each crystal form.31−40 Effect of Self-Poisoning on Nucleation Rate. The analysis of nucleation events per unit area as a function of crystallization temperature in the long-chain n-alkane C246 H494 crystallized from the melt led to the same retardation effect as observed for the growth rate.41 It was then concluded that selfpoisoning was retarding both primary crystal nucleation and crystal growth. To test if the retardation is also observed in the primary nucleation rate of precision polyethylenes, we computed the number of spherulites observed at the beginning of the crystallization as a function of time at a fixed isothermal crystallization temperature, and obtained the nucleation rate from the slope of the initial data. The early developing

Figure 8. (a) Nucleation rate at increasing crystallization temperatures for PE21Br. (b) Time for incipient crystallization at increasing crystallization temperatures for PE9Br.

Figure 8a displays nucleation rate data for PE21Br where a minimum is evident at Tc = 64 °C in full agreement with the temperature gradient of the growth rate. A discontinuity is also appreciable in the time to start crystallization data of PE9Br for temperatures in the range of 2 and −2 °C (Figure 8b), which is basically the same range where a discontinuity in the growth rate was found. Therefore, as in the case of long-chain nalkanes, we conclude that the observed retardation rate effect which is consistent with frequent but inefficient deposition events of Form I on the surface of Form II is operative in both the nucleation and growth of the crystallization process. Relevance to Crystallization Rate Theories. The observation of a minimum in the nucleation and growth rate of long-chain n-alkanes at temperatures where folded chains begin to form (TmF2) presented a unique case for interpretation of crystallization theories, and although it was predicted to be applicable to the general crystallization of high molar mass polydisperse polymers, direct evidence in the latter has been elusive.16 To the best of our knowledge, the series of precision bromine-containing polyethylenes constitute the first clear evidence of a polymer system with retardation of crystallization rate consistent with self-poisoning. They make the case for reviewing the theories that describe polymer crystallization, especially kinetic theories that incorporate events at the developing surface nucleus and during crystal growth. The drastic retardation of growth rate observed in long-chain n-alkanes was first used by Ungar and co-workers16,18 to contrast kinetic theories of growth of a polymer lamellar crystal by events that take place at the crystal surface. The most relevant are the secondary nucleation theory of Lauritzen and Hoffman (LH)13 and the surface roughness theory of Sadler and Gilmer.19−21 In the first, growth is controlled by the deposition of a new chain stem with exactly the same length and spreading via regular and adjacent folds on a suitable lateral surface.13 In the second, shorter and longer stems are continuously absorbing and desorbing creating a rough surface where segments with the wrong length may spend a long time in a metastable configuration before they end in the stable form.19 In other words, in the HL theory, the surface is smooth, and the length of the stem that is attaching to the surface is fixed while it is variable in SG theory. I

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Macromolecules In its original form, the LH theory cannot describe the observed retardation; hence, the role of different stem lengths in the second appeared initially to be the appropriate theoretical path to explain the blocking effect of a half-length segment on the propagation of the extended one in n-alkane crystal growth.16 On these grounds, Ungar and co-workers dedicated extensive efforts to document temperature gradients of the growth rates and associated morphological changes that reflected the expected flatness (low Tc) or curvature (high Tc) of the crystal surfaces.22,28,29 Adherence of the SG morphological predictions with the experimental observations was not general. It was shown for long-chain n-alkanes that growth is in fact nucleation controlled to a significant degree while the SG model uses a non-nucleated roughness-pinning approach.29 The morphologies of the precision bromine series studied here (Figure 5) do not provide support for the SG approach either. With increasing temperature, the morphology of PE15Br changes before and after the rate minimum; however, similar changes are not observed in PE21Br. Hence, neither the LH not the SG theory can describe adequately the minimum of the growth rate or the crystal morphology. Both theories have been criticized almost since their inception on different grounds.15,42,43 The observed sharp minima of the growth rates of long-chain n-alkanes and for the polymer studied here would be even more difficultly explained by more recent polymer crystallization theories that are based on multistep evolution of the original melt in phases that acquire increasing order by continuous accretion and reorganization within the phases.44−46 While none of the present theories can adequately explain the observed rate retardation, testing the possible controlling role in crystal growth of the length of the stem that is attaching to the surface is of interest in view of nucleation theories that base productive stem attachment on entropic barriers at the interface between the melt and the developing nucleus.19,43,47 In the series of precision polyethylenes investigated here, the change in conformation of the lateral packing stem between Form I and Form II is so drastic that conformation rather than stem length appears as the most plausible cause for selfpoisoning. Moreover, if the crystals of Form I are considerably thinner than those of Form II, the stem length may be a controlling factor in the retardation of the rate in a similar manner to the behavior of n-alkanes. To analyze if there is significant increase in crystal thickness at the transition between Form I and Form II, SAXS data were recorded on heating at a rate of 1 °C/min samples of PEI5Br and PE21Br that were rapidly crystallized from the melt into Form I. During slow heating, Form I melts and quickly recrystallizes into Form II as shown in Figure 2 and Figure SI.1. Therefore, if there are major differences in crystal thickness between both forms, a strong increase of the long period will be observed after recrystallization. A composite is shown in Figure 9 of the DSC melting thermograms (Figure 9a), the variation of long spacing (L) during heating (Figure 9b), and the projection of Lorentz corrected SAXS intensity versus scattering vector (Figure 9c). As shown, for both samples the long spacing (L) increases slowly up to the first melting due to temperature expansion and removal of thin imperfect crystallites. In PE15Br the meltrecrystallization is prominent at temperatures between 40 and 45 °C by a deep decrease in the intensity of the SAXS peak (melting) and a decrease in L (recrystallization), in full agreement with the DSC data. Moreover, the L value of the

Figure 9. Structural change of Form I for PE15Br and PE21Br during heating at 1 °C/min: (a) DSC melting thermograms, (b) long spacing, and (c) Lorentz-corrected SAXS patterns with increasing temperature.

initial crystallites (Form I) and the value after recrystallization at 45 °C (Form II) are basically the same (∼170 Å). After recrystallization, at T > 45 °C, further melting is again indicated by the increase in long period. The region of second melting is demarcated in Figure 9 by the dashed line. Data of the core crystal thickness were extracted from one-dimensional correlation function analysis of the SAXS data using AFM images as reference to distinguish between the thickness of crystalline and amorphous regions. The core crystal thickness at temperatures prior and after recrystallization is 120 ± 10 Å (data shown in Figure SI.4). Because melting and recrystallization of Form I of PE21Br take place in a narrower range of temperatures during heating, the effect on L is less prominent, but it is also evident in Figure 9 by the decrease of SAXS intensity at 65 °C and recovery after recrystallization at higher temperatures. More important is the fact that the long spacing of Form I before (T < 65 °C) and after recrystallization to Form II (T = 65 °C) is basically unchanged at ∼230 ± 10 Å. Only during melting of Form II (T > 67 °C) does the long spacing increase again as expected. The J

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SAXS patterns collected at the isothermal crystallization temperatures for temperatures below and above the transition from Form I to Form II of PE15Cl also demonstrated that L is basically unchanged.8 In view of the small change in crystal thickness at temperatures above and below the transition between both forms, we conclude that it is the conformation rather than the length of the stem approaching the growing surface what drives self-poisoning. More important is the observation of a rate minimum in high molecular weight polymers that confirms early predictions based on the initial observations in n-alkanes. As previously advocated,48,49 the present data on high molar mass precision polyethylenes lend additional support of a polymer crystallization process controlled by events that take place at the growth front rather than by other events that may occur at earlier stages in the polymer melt. We anticipate that similar rate minima should prevail in other precision polyethylenes that develop layered polymorphic crystalline structures controlled by a difference in crystallization kinetics.

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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.7b02745. Melting thermograms of Form I at different heating rates; polarized optical micrographs of Form I; natural logarithm of linear growth rates of Form II as a function of crystallization temperature; linear correlation function, long spacing, and crystal thicknesses on melting at 1 °C/ min for PE15Br and PE21Br (Form I); AFM images of PE21Br (Form II) (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected] (R.G.A). ORCID

Rufina G. Alamo: 0000-0002-3061-499X Notes

The authors declare no competing financial interest.

CONCLUSIONS

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Precision polyethylenes with bromines atoms placed on every 21, 19, 15, or 9 backbone carbons develop two types of layered crystalline forms depending on undercooling. The all-trans, Form I structure is formed in the low crystallization temperature (Tc) range while a herringbone type of structure, Form II, develops in the high Tc range. The Tc at the transition from Form I to Form II has been identified by WAXD patterns and by differences in melting points. For all samples, Form I melts at 10−30 °C lower temperatures than Form II. The difference Tm,FII − Tm,FI decreases with increasing distance between bromines and merges at ∼95 °C for a chain with Br spaced by ∼50 carbons. It is predicted that such long spaced polyethylenes with bromine will not develop the herringbone structure. The temperature gradient of the growth rate displays an anomalous inversion with increasing undercooling analogous to the crystallization rate minima observed in long-chain n-alkanes. The rate minimum in precision Br-containing polyethylenes is observed at or very near the melting temperature of Form I and is consistent with “self-poisoning”. At Tc near Tm,FI approached from above, frequent but unstable depositions of Form I on the growing surface of Form II block momentarily growth until the wrong segment conformation detaches. The depth of the minimum depends on differences of melting points between each form. It is deep for PE21Br and becomes shallower for PE15Br and PE9Br. Although it was predicted much earlier, the present growth rate data on precision polyethylenes provide the first direct experimental evidence in support of self-poisoning in high molar mass polymers, and while the classical crystallization theories are unable to explain the observed growth minima, the present data are in support of polymer crystallization mechanisms led by events that take place at the growth front. The peculiarity of precision polyethylenes with Cl or Br is that they can assemble in two crystalline forms with drastic differences in chain conformation and nucleation barriers, such that approaching the transition between both forms, large and easily detectable effects can be observed during growth.

ACKNOWLEDGMENTS Funding of this work by the National Science Foundation, Polymer Program DMR 1607786, is gratefully acknowledged. We thank Dr. Carolina Ruiz-Orta, Dr. Juan M. Majada, Dr. Masafumi Tasaki, and Ms. Amanda Stefin (REU student) for help with X-ray, DSC, and POM experiments. We are also indebted to the High Performance Materials Institute of Florida State University for access to X-ray instrumentation.

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REFERENCES

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DOI: 10.1021/acs.macromol.7b02745 Macromolecules XXXX, XXX, XXX−XXX