Effect of Hydroxyl-Functionalization on the Structure and Properties of

3 Jul 2013 - Synthesis of Functional Polypropylene Containing Hindered Phenol Stabilizers and Applications in Metallized Polymer Film Capacitors...
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Effect of Hydroxyl-Functionalization on the Structure and Properties of Polypropylene Sahil Gupta,† Xuepei Yuan,‡ T. C. Mike Chung,‡ S. Kumar,§ M. Cakmak,† and R. A. Weiss†,* †

Department of Polymer Engineering, The University of Akron, 250 South Forge Street, Akron, Ohio 44325, United States Department of Materials Science and Engineering, The Pennsylvania State University, University Park, Pennsylvania 16802, United States § Department of Chemical Engineering, Columbia University, New York, New York 10027, United States ‡

ABSTRACT: Hydroxyl-modified polypropylenes (PPOH) with side chains containing OH groups were synthesized by copolymerization of the propylene and undecenyloxytrimethylsilane monomers. Copolymers with hydroxyl concentration up to ∼4 mol % were produced by hydrolysis of the propylene−undecenyloxytrimethylsilane copolymers and their structural, thermal and rheological properties were compared with that of unmodified polypropylene (PP). The presence of H-bonding between the hydroxyl groups led to a densification of the amorphous domains and an increase in the glasstransition temperature with increasing hydroxyl concentration. The incorporation of comonomer in the PP backbone led to a reduction in the melting point, as predicted by Flory’s theory of melting point depression in copolymers. During crystallization, the OH containing polymer segments are excluded from the crystalline domains and their presence disrupt the growth behavior leading to reduced lamellar thickness, leaving the α-monoclinic crystal structure unaffected. Because of the steric hindrance by the comonomer side-chains to the ordered packing of PP backbone chains into the lamellar structure, a reduction in crystallinity by as much as 15% and a significant reduction in the overall rate of crystallization was observed for the PPOH copolymers. Both the storage and the loss moduli decreased for the PPOH copolymers below the melting point, owing to a lower crystallinity than PP, but increased above the melting point due to the dominating effect of H-bonding interactions in the melt state.

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INTRODUCTION

kinetics and the thermal and rheological properties of isotactic polypropylene (PP).

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In the past decade, research in the area of polymeric capacitors has been aimed at increasing the dielectric constant, energy density and the breakdown strength while decreasing the energy loss of the dielectric material.1−4 The present state-ofthe-art material for high voltage capacitors is metallized biaxially oriented polypropylene (BOPP).2,5,6 BOPP, however, has a low dielectric constant (∼2.2) and provides a low energy density of 3−5 J/cm3 at breakdown, which occurs at 600−750 V/μm for ∼10 μm thick films.4,7 One method for improving the dielectric properties of polymers is to incorporate functional side groups that are highly polarizable and that can align under the influence of an electric field. Yuan et al.4 recently reported the synthesis of hydroxylfunctionalized isotactic polypropylene (PPOH) copolymers, which exhibited energy densities 2−3 times that of BOPP without a major increase in the energy loss. While the incorporation of hydroxyl functionality improves the dielectric properties, it also affects the polymer structure and the thermal and rheological properties, which are important with regard to the processing of film and the capacitor application. This paper describes the effects of the incorporation of hydroxyl groups (−OH) on the crystalline microstructure, crystallization © 2013 American Chemical Society

EXPERIMENTAL SECTION

Materials. The synthesis and structure of the PPOH copolymers are shown in Figure 1. The detailed synthesis procedure is described elsewhere.4 A heterogeneous Ziegler−Natta copolymerization of propylene and undecenyloxytrimethylsilane was used to prepare the PPOH copolymers. The intermediate silane was the product of the polymerization, and it was hydrolyzed under acidic conditions to form the corresponding PPOH copolymer. The reactivity ratios for this copolymerization were r1 = 68 for propylene and r2 = 0.032 for the comonomer. This means that both species are far more reactive with propylene than with the silane comonomer. Thus, during the early stages of the batch reaction, propylene rich chains are formed, but as the polymerization proceeds to higher conversion, more comonomer is incorporated due to the depletion of propylene monomer. During the later stages of polymerization when the comonomer concentration is high, the comonomer is incorporated at a higher rate than at the earlier stages, which produces a OH-richer copolymer structure.8 The concentration distribution of the copolymers, however, was not determined. For comparison purposes, a PP sample was also prepared Received: April 26, 2013 Revised: June 24, 2013 Published: July 3, 2013 5455

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Figure 1. Synthesis of PPOH copolymers.

Table 1. Molecular and Thermal Characteristics of PP and PPOH Polymers ϕm (%) polymer PP PPOH1.3 PPOH1.4 PPOH1.7 PPOH3.9

[OH]a (mol %)

Mnb (kDa)

Mwb (kDa)

Mzb (kDa)

PDI

0 1.3

38 16

385 243

1479 1405

10.2 14.8

1.4

27

488

2235

17.9

1.7

−

−

−

−

3.9

−

−

−

−

Mvc (kDa)

ρi (g/cm3)

ρaj (g/cm3)

43 35

0.902 0.895

0.856 0.864

34

33

0.895

0.865

35

32

33

0.899

0.874

37

34

34

0.902

0.877

Tgd (°C)

Tα*e (°C)

Tmf (°C)

Tdegh (°C)

DSCf

− −

−4.2 −2.0

98 68

163 158

343 324

50 38

49 35

−

−1.7

67

157

352

37

388

0

65

157

324

301

0.3

65

157

325

DSCg WAXD

From 1H NMR spectra. bFrom HTGPC with PS standards; PPOH-1.7 and PPOH-3.9 could not be dissolved. cCalculated from [η] = K.Mvα, where K = 1.05 × 10−5 and α = 0.89, [η] was measured in decalin at 135 °C. dFrom DSC, 3rd heat at 50 °C/min. eFrom DMA, E″ at f = 100 Hz. fFrom DSC, 1st heat at 10 °C/min. gFrom DSC, 2nd heat at 10 °C/min after quenching from the melt. hFrom TGA at 20 °C/min under N2 flow, temperature corresponds to 1% degradation. iFrom gas-pycnometer at 26 °C. jUsing ρa = (ρ − ρcϕm)/(1 − ϕm). a

samples. Tg was defined as the inflection point in the change of the heat capacity associated with the glass transition. A heating rate of 10 °C/min was used to determine the melting point and the enthalpy of fusion (ΔHf). The melting point (Tm) was defined as the peak of the melting endotherm, and the mass fraction crystallinity of the polymer (ϕm) was calculated from the ratio of ΔHf/ΔHPP, where ΔHf was the area under the DSC melting endotherm and ΔHPP is the enthalpy of fusion for a 100% crystalline PP. Reported values of ΔHPP vary from 138 to 250 J/g.9 According to Wunderlich,10 the most reliable values of ΔHPP are direct calorimetric measurements and the average value from five different research groups is 165 ± 18 J/g. In this paper, the value of ΔHPP was used. Any error in that value, however, will change the absolute values of the heats of fusion reported, but the relative comparisons between the different PPOHs and PP will effectively be the same. DSC was also used to study the isothermal crystallization kinetics of the polymers at 124, 126, 128, and 130 °C. The polymer sample was first heated and equilibrated to 200 °C (above the equilibrium melting temperature of iPP, Tmo = 185 °C9), held isothermally for 3 min at 200 °C to remove any thermal history, cooled at 40 °C/min to the chosen crystallization temperature (Tc), and then maintained isothermally until crystallization was complete as determined by a zero differential heat flux. The data were analyzed using the Avrami equation11 to obtain the Avrami constant (K), which is associated with the overall rate of crystallization, and the Avrami exponent (n) that is associated with the mode of crystallization. The thermal stability of the polymers was measured by TGA using a heating rate of 20 °C/min and a nitrogen atmosphere. DMA was used to measure the temperature and frequency-dependent dynamic modulus, E′(T,f), and to identify the secondary-transition temperature associated with the slippage of crystal planes, (Tα*).12 Tα* was defined as the temperature of the peak in the loss modulus, E″(T,f), associated with that transition. The DMA measurements were performed in tension using a multiple-frequency strain mode and frequencies of f = 100 and 1 Hz from −100 to +150 °C, a heating rate of 2 °C/min and a strain of 0.2%. The crystal structures were characterized with wide-angle X-ray diffraction (WAXD) and small-angle X-ray scattering (SAXS). WAXD measurements were performed using a Bruker AXS D8 Goniometer and SAXS measurements were performed using a Rigaku Micro-

using the same synthetic procedure, but without the comonomer. The samples used in this study were denoted as PPOH-y, where y is the molar concentration of the OH groups (i.e., the comonomer concentration), and their characteristics are summarized in Table 1. Characterization. Molecular weight averages were measured by high temperature gel permeation chromatography (HTGPC) and intrinsic viscosity measurements. For HTGPC measurements, 4 mg of polymer was dissolved in 1 mL trichlorobenzene (containing 0.1 wt % each of Irganox 1010 and Irgafos 168 antioxidants) at 145 °C and the solution was analyzed in the GPC at 135 °C. The instrument was calibrated using polystyrene standards. The PPOH copolymers with 1.7 and 3.9 mol % [OH] were not completely soluble in trichlorobenzene, which was the carrier solvent used in the HTGPC. For those two samples, intrinsic viscosity measurements were made with a Cannon-Ubbelohde viscometer using decahydronaphthalene (Decalin) solutions at 135 °C. The viscosity-average molecular weight was estimated by the Mark−Houwink equation using the constants for polypropylene: [η] = KMvα where K = 1.05 × 10−5 and α = 0.80.9 Mw and Mv were in the range of ∼250−500 kDa for all the polymers used, though the polydispersity (PDI) appeared to increase with increasing OH concentration. Overall, the relatively high PDI values are associated with high monomer conversion during the batch copolymerization reactions. Polymer films were compression molded at 200 °C, quenched to room temperature at ∼40 °C/min and annealed under vacuum at 80 °C for 24 h. Hydrogen bonding interactions were characterized by Fourier transform infrared spectroscopy (FTIR), using a Thermo Scientific Nicolet 380 spectrometer. Absorption spectra of 20 μm thick films were measured using 64 scans with a resolution of 4 cm−1. Thermal properties were measured with a TA Instruments Q-200 differential scanning calorimetry (DSC), a TA Instruments Q-50 thermogravimetric analysis (TGA) and TA Instruments Q-800 dynamic mechanical analysis (DMA). DSC was used to measure the glass transition temperature (Tg), the melting point (Tm) and the heat of fusion (ΔHf) of the polymers. Before the measurements, temperature calibration was performed using an Indium standard (Tm = 156.6 °C). Tg was measured from heating thermograms using a heating rate of 50 °C/min. The unusually fast heating rate was needed to resolve the relatively low change in the heat capacity at the glass transition due to the high crystallinity of the 5456

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Figure 2. (a) FTIR spectra of PP and PPOH-y polymers at room temperature for the spectral region of 3100 to 3700 cm−1; (b) FTIR spectrum of PPOH-1.7 at 40, 100, 160, and 200 °C. Max002+ diffractometer. Both instruments used CuKα radiation (λ = 0.1542 nm). The WAXD data were used to compute the crystal lattice spacings (dhkl) and the degree of crystallinity, and SAXS was used to determine the long-spacing (LP) of the crystals. Bulk density (ρ) measurements of the polymers were made with a AccuPyc II 1340 gaspycnometer at room temperature (∼26 °C). Linear viscoelastic (LVE) properties at 165−250 °C, f = 0.1−100 rad/s and a strain amplitude of 5−8% were measured with a TA Instruments ARES G2 rheometer using small amplitude oscillatory shear (SAOS) and a parallel-plate fixture. Prior to the SAOS measurements, a strain-sweep was performed at 200 °C and 1 Hz to determine the linear response limit. LVE mastercurves were developed using time−temperature superposition with a reference temperature of To = 200 °C.

interchain OH dimers. The presence of H-bonded, bound water produced clusters that enhanced the dielectric constant of the polymer. The bulk density (ρ) of the polymers was measured with a gas pycnometer and the values are summarized in Table 1. The bulk density was assumed to be a weighted average of the densities of the crystalline and amorphous phases, ρ = ϕmρc + (1 − ϕm)ρa

(1)

where ρc is the density of a 100% PP crystal (0.946 g/cm3),9,10 ρa is the density of the amorphous phase, and ϕm is the mass fraction crystallinity of the polymer. It is assumed that the density of the crystalline phase was unaffected by the introduction of OH in the polymer, i.e., the OH groups were located exclusively in the amorphous phase. That assumption is supported by the molecular dynamic simulations of Agarwal et al.16 Eq 1 was used to calculate the density of the amorphous phase. Those values are included in Table 1, and they indicate that the density of the amorphous phase increased with increasing [OH]. One might intuitively expect that incorporation of the OH and the alkyl chain to which it is attached might decrease the density of the amorphous phase as a consequence of steric hindrance of the alky chains. No evidence of side-chain crystallinity was observed. It is also very unlikely that the side chains would be crystalline, since the melting point of a C9 alkane crystal (undecane) is −25 °C and the melting point for any shorter segments of the side chain would be even lower. The only phenomenon that would explain the increase in the packing of the amorphous phase by the introduction of the OH substituent is a densification of the amorphous phase due to H-bonding interactions, which again is consistent with the results described in ref 15. The glass-transition temperature (Tg) increased with increasing OH concentration, see Table 1, which is probably also a consequence of the hydrogen bonding and the densification of the amorphous phase. The melting points (Tm) of the copolymers were 7−8 °C lower than that of the PP, though increasing the concentration of OH from 1.3 to 3.9 mol % had little effect on Tm. For random copolymers, one expects that the introduction of a noncrystallizable comonomer into a crystallizable polymer will lower the melting point17,18 and Tm will decrease monotonically with increasing comonomer concentration, in this case the −OH-containing monomer.

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RESULTS AND DISCUSSION Physical Properties. The thermal properties of the polymers are summarized in Table 1. The thermal degradation temperature (Tdeg), which was defined as the temperature corresponding to a 1% mass loss, was greater than 300 °C for PP and the PPOH polymers. The introduction of the OH groups increased the polymer’s affinity for water, but had an insignificant effect on its thermal stability in the temperature range used for processing or application. H-bonding of the OH groups in the PPOH copolymers was evident from the room temperature FTIR absorption spectra for PP and the PPOH copolymers in the spectral region 3700− 3100 cm−1 where the O−H stretching vibration occurs, Figure 2a. A low intensity peak at 3637 cm−1 corresponds to the stretching vibration of free (non-hydrogen bonded) OH groups and a higher intensity, broad absorption at 3332 cm −1 corresponds to the stretching vibration of hydrogen bonded OH groups.13,14 Both absorptions are absent in the PP sample. The intensity of the 3332 cm−1 peak increased relative to the 3637 cm−1 peak as the concentration of OH in the polymer increased. Although the strength of the hydrogen bond decreased with increasing temperature, as evident from a shift of the 3332 cm−1 peak position to a lower higher-frequency, a significant fraction of the hydrogen bonded OH groups persisted at melt temperatures, Figure 2b. These results indicate that most of the OH groups are associated by H-bonding in the PPOH-y, and that observation is consistent with molecular simulations by Wang et al.15 that showed that the OH groups in a hydroxyl modified polyethylene formed a network structure from H-bonded 5457

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An example of peak separation and curve fitting for PP is shown in Figure 4. The lattice spacings (dhkl) were calculated

However, as seen in Table 1, the Tm of the polypropylene crystals decreased from 163 to 157−158 °C upon the addition of ∼1.3 mol % OH, but the Tm remained constant for higher hydroxyl concentrations. That result can be explained by microphase separation of the OH-containing species due to the H-bonding of the OH groups. The simulations of Wang et al.15 show the formation of hydrogen bonded OH groups for relatively low concentration of hydroxyl groups attached directly to a polyethylene chain, and one might reasonably expect the same for polypropylene. In addition, there might be a tendency to form OH nanoclusters above a critical hydroxyl concentration. The decoupling of the OH groups in the PPOHy copolymers by a long alkyl chain adds additional mobility to the OH species and the evidence for hydrogen bonding discussed above suggest that microphase separation may have occurred in these copolymers. Unfortunately, the electron density contrast between O−H and C−H phases was insufficient to observe microphase separation with small-angle X-ray scattering. Crystalline Structure. Figure 3 shows 1-D WAXD patterns after background correction and the assignments of the Miller

Figure 4. Peak separation and curve fitting from WAXD data of PP. The experimental data in the middle graph is represented by the solid curve and the fit is given by the dotted curve.

using Bragg’s law (nλ = 2d[sin θ], where n is an integer, λ is the wavelength of the radiation (λCuKα = 0.1542 nm), and θ is half of the diffraction angle) and are plotted as a function of hydroxyl concentration in Figure 5. The [OH] had no effect on

Figure 3. WAXD of compression-molded PP and PPOH copolymers. The diffractograms are shifted vertically for clarity. The data are consistent with an α-monoclinic structure. The peak at 2θ = 15.95° for lattice index (300) corresponds to a β-hexagonal structure. Note that this peak is seen only for PP and PPOH-1.4 and the intensity is negligibly low as observed during the peak resolution procedure. The inset is an example of the 2-D WAXD data for PP.

Figure 5. Lattice spacings as a function of [OH] for the different (hkl) crystalline planes.

the d-spacings, which indicated that the size of the unit cell of the PP crystals was unaffected by the incorporation of the comonomer in polypropylene. The fwhm values for the major lattice planes, which are summarized in Table 2, were also unaffected by the hydroxyl concentration. Note that the contribution to peak broadening from instrumental broadening was not decoupled from that of crystal distortion, but it was assumed to be the same for all the polymers. The WAXD data further support the conclusion that the OH groups are located solely in the amorphous phase. The data in Figure 3 also show that the amount of crystallinity of the copolymers decreased with increasing [OH], which is consistent with the DSC results, see Table 1. The mass fraction of crystalline phase, ϕm, was calculated from the WAXD data using eq 2,

indices (hkl) for the crystal planes of the PP and PPOH copolymers. Each material shows the presence of an αmonoclinic crystal structure,19−21 which is typical for PP. The insensitivity of the crystal structure to the incorporation of the OH-containing monomer is consistent with the conclusion that the branch and OH groups were excluded from the crystals. The van der Waal (VDW) molar volume of a methyl group (−CH3) is much larger than that of the hydroxyl group, 13.67 cm3/mol vs 8.04 cm3/mol.22 For H-bonded OH groups, the VDW volume of the hydroxyl may decrease by as much as 1.8 cm3/mol depending on the angle and bond length of the Hbond.22 As a consequence, if the OH were incorporated into the PP crystal, one would expect to detect a change in the lattice spacing (dhkl) and/or the crystal distortions that should manifest as peak broadening in WAXD.23 dhkl and the full width at half-maximum (fwhm) were calculated from the data in Figure 3 by deconvoluting the amorphous and crystalline peaks. The baseline was fitted with a cubic function, a Gaussian function was used to fit the crystalline peaks and a Voigt function was used to fit the amorphous peak.

ϕm = Ac /(Ac + A a )

(2)

where Ac is the area under the crystalline peaks and Aa is the area under the amorphous peak in Figure 4. The values of ϕm calculated from the WAXD data were comparable, maybe a little lower than those calculated from the DSC data, see Table 5458

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Table 2. FWHM and average crystal dimensions for PP and PPOH polymers fwhm (degree) polymer

[OH] (mol %)

(110)

(040)

(130)

(111)

(041)

LP (nm)

lc (nm)

la (nm)

PP PPOH-1.3 PPOH-1.4 PPOH-1.7 PPOH-3.9

0 1.3 1.4 1.7 3.9

0.7 0.8 0.8 0.8 0.8

0.6 0.6 0.5 0.6 0.5

0.7 0.7 0.6 0.7 0.7

0.7 0.7 0.7 0.7 0.8

0.8 0.9 0.9 0.9 0.8

14.0 12.5 12.4 12.0 12.1

7.4 5.0 4.9 4.4 4.7

6.6 7.5 7.6 7.6 7.4

Figure 6. (a) SAXS patterns for PP and PPOH polymers. The intensity data were smoothed and corrected for parasitic scattering, Lorentz effect and the geometry. The inset is an example of the 2-D SAXS pattern for PP. (b) 1D correlation functions.

Figure 7. Isothermal crystallization exotherms at different crystallization temperatures for (a) PP and (b) PPOH-1.4. Note the different scales for the time axis in the two plots.

More details of the structure of the polymers can be obtained from the autocorrelation function, Γ(r) which provides the spatial correlation of the electron density fluctuations,24,25

1, but the general effect of [OH] on crystallinity determined by both techniques was essentially the same. The SAXS patterns (Iq2 vs q) are shown in Figure 6a for all the polymers. The data were smoothed using a Loess function and corrected for parasitic scattering, the Lorentz effect and geometry. The Lorentz and geometric corrections involved multiplying the measured intensity by the scattering wave vector (q).23 The SAXS patterns reveal a well-defined interference peak that arises from the periodic arrangement of the crystalline lamellae. For the PP homopolymer a shoulder of lower intensity at higher q is also observed, which is the second order reflection of the large peak. The ratio of the Bragg angles of the second and first maxima is 1.95 (q1 = 0.43 nm−1, q2 = 0.84 nm−1), which is close to the value of 2.0 that is expected for a lamellar structure. The second order reflection was not observed for the PPOH samples, which indicates that the lamellar periodicity is not as well developed in those materials.

Γ(r ) =

1 Q

∫0

∞

I(q)q2 cos(qr ) dq

(3)

where r is the distance between any two points that are in the same phase, Q is the scattering invariant Q=

∫0

∞

I(q)q2 dq=(2π )2 ϕm(1 − ϕm)(Δρ)2

(4)

and Δρ is the electron density contrast between the two phasesin this case, the crystalline and amorphous regions. The one-dimensional autocorrelation function used here assumes that the electronic density in a crystal changes only along the chain axis, i.e., perpendicular to the lamellar plane, or that the lamellae spacings occur along one fixed axis. The SAXS data in Figure 6a were extrapolated to q = 0 and q → ∞ using a Guinier and a sigmoid function, respectively, and numerical 5459

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Figure 8. Relative crystallinity as a function of time for the isothermal crystallization of (a) PP and (b) PPOH-1.4.

Figure 9. Avrami plot for the isothermal crystallization of (a) PP and (b) PPOH-1.4. The solid lines are fits of eq 8 to the experimental data. Longtime experimental data corresponding to secondary crystallization are not included.

methods were used to integrate eqs 3 and 4. The Γ(r) functions for the PP and the PPOH polymers are shown in Figure 6b. The supramolecular crystalline phase consists of stacks of alternating crystalline and amorphous lamellae. It was assumed that the dimensions of those stacks were larger than what was resolved in the SAXS data shown in Figure 6,25 such that the peak in the autocorrelation function (Figure 6b) corresponds to the long spacing of the crystal, LP = lc + la, where lc is the thickness of the crystal lamella and la is the thickness of the amorphous folds. The 2-D SAXS pattern shown in the inset in Figure 6a indicates that the morphology of the polymers was isotropic. For the PPOH copolymers, Lp decreased about 10−15% upon the introduction of the OH-containing comonomer, but the long-spacing was relatively insensitive to the extent of the modification, see Table 2. The thicknesses of the crystal lamella (lc) and the amorphous fold-region (la) for each polymer were calculated from the crystal volume fraction (ϕv), ϕv =

accommodate the long alkyl chains containing the OH groups, and the insensitivity of the lc and la values to the concentration of OH support the conclusion that the microstructure of those materials was somewhat blocky. Isothermal Crystallization Kinetics. Figure 7 shows isothermal crystallization exotherms for PP and PPOH-1.4 for different crystallization temperatures (Tc). For all the polymers, the time required for complete crystallization decreased with increasing undercooling, ΔT = Tmo − Tc). The time-dependent relative crystallinity X(t) was calculated from the data in Figure 7 using the equation X (t ) =

Q∞

=

t

( ddHt ) dt ∞ ∫0 ( ddHt ) dt ∫0

(6)

where Qt is the total exothermic heat flux measured by DSC at time t during crystallization and Q∞ is the total heat flux for the complete crystallization process. Figure 8 shows the isothermal kinetics of the crystallization of PP and PPOH-1.4. These results were fit to the Avrami model,11

ϕmρm ϕmρm + (1 − ϕm)ρa

Qt

(5)

where ϕm is the crystal mass fraction given in Table 1 and ρc and ρa are the mass densities of the crystalline and the amorphous phase, respectively. The values of lc = ϕvLP and la = (LP − lc) are summarized in Table 2. The thicker crystal lamellae for PP is consistent with its higher crystallinity, cf. Tables 1 and 2. The smaller lc and lower crystallinity of the PPOH copolymers is due to the inability of the crystals to

X(t ) = 1 − exp( −Kt n)

(7)

where K is the Avrami constant, which is associated with the overall rate of crystallization, and n is the Avrami exponent, which is associated with the mode of nucleation, growth geometry and growth velocity. Equation 7 may be rewritten as, 5460

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log[− ln(1 − X (t ))] = log K + n log t

that strong intermolecular interactions, such as H-bonding, suppress crystal growth rate and, therefore, the overall rate of crystallization.26 The increase of K for the highest hydroxyl concentration may be a consequence of an increase in the heterogeneous (primary) nuclei density. The simulations of Agarwal et al.16 suggest that for the PPOH-3.9 sample microphase separation of nanoclusters of the hydroxyl species may occur, and those nanostructures may provide sufficient order for accelerating crystallization27 or nucleating crystallization in a manner similar to what has been proposed for semicrystalline ionomers.28 Table 3 also includes the crystallization half-times for the PP and PPOH polymers calculated from eq 9. t1/2 increased by an order of magnitude as the [OH] increased from zero (PP) to 1.7 mol % (PPOH-1.7), and that was followed by a decrease when [OH] increased to 3.9 mol %. Those results are consistent with the trends observed for K. Figure 10 compares the polarized optical micrographs of isothermally grown crystals of PP, PPOH-1.7 and PPOH-3.9 at 130 °C. The spherulitic texture in PP was more organized and space-filling compared to the coarser and less definite texture in the PPOH copolymers. As discussed before, that was a result of steric hindrance from the comonomer side chains that disrupts the packing of the PP backbone chains in the crystalline phase. Rheology. Dynamic mechanical results for PP and the PPOH copolymers are shown in Figure 11. The data below Tm were obtained from tensile measurements and the data above Tm were calculated from shear measurements using E′ = 3G′ and E″ = 3G″. The dynamic tensile modulus, E′, below Tm decreased when the OH-containing monomer was incorporated, which is due to the decrease in crystallinity. Above Tm, E′ increased with increasing [OH], which indicates that the copolymer melts were more elastic than that of PP. The dynamic shear viscosity, η′ = G″/ω also increased with increasing [OH]. Those results are consistent with the idea that hydrogen bonding between hydroxyl groups occurs in the PPOH melts, which increases the viscosity and the elasticity as a consequence of strong intermolecular interactions. The increase of the melt elasticity suggests that the PPOH copolymers may have higher melt strength than PP, but this property was not assessed in this study. Two transitions in E′ and two maxima in E″ are seen in the data in Figure 11. For E′, the two transitions are associated with Tg and Tm, the latter producing a much larger change in the modulus as the material transforms from a semicrystalline solid to a viscous liquid melt. The lower temperature peak in E″ corresponds to the glass transition, which is usually referred to as the α-transition. The higher temperature peak in E″ is the

(8)

The half-time for crystallization, (t1/2), is calculated by setting X(t) = 0.5 in eq 7,

t1/2 =

⎛ ln 2 ⎞1/ n ⎜ ⎟ ⎝ K ⎠

(9)

Figure 9 shows the Avrami plot for PP and PPOH-1.4. The solid lines are linear fits of eq 8 to the experimental data. The values of n and K, calculated from the slope and intercept of the Avrami plot, respectively, are summarized in Table 3. For all the Table 3. Isothermal Crystallization Constants for PP and PPOH Polymers polymer

PPOH-1.3

PPOH-1.4

PPOH-1.7

PPOH-3.9

a

Tc (°C)

N

K × 102 (min−n)

t1/2 (min)a

124 126 128 130 124 126 128 130 124 126 128 130 124 126 128 130 124 126 128 130

1.9 2.0 2.1 2.1 2.0 1.9 1.9 1.8 2.1 1.9 1.9 1.9 2.1 2.0 1.9 1.9 2.0 1.9 1.8 1.8

11.75 5.89 3.09 1.29 0.63 0.30 0.19 0.09 0.43 0.30 0.12 0.06 0.16 0.09 0.03 0.01 0.56 0.27 0.14 0.06

2.6 3.5 4.3 6.9 10.7 17.0 24.5 38.8 11.7 18.2 28.0 43.3 17.4 27.0 51.3 79.9 11.5 18.6 29.0 46.9

Calculated from eq 9

polymers and Tcs used, n ∼ 2, which suggests a heterogeneous mode of nucleation with a two-dimensional growth and a constant growth velocity. Hence, incorporation of the comonomer in PP did not affect the mode of crystallization. The increase of K with decreasing Tc, see Table 3, is consistent with the observed faster crystallization kinetics with increasing supercooling. K decreased by about an order of magnitude when [OH] increased from 0 to 1.7 mol %, but then increased when the hydroxyl concentration was 3.9 mol %. The decrease in K with increasing [OH] is consistent with the idea

Figure 10. Polarized optical micrographs for (left to right) PP, PPOH-1.7 and PPOH-3.9 during isothermal crystallization. Scale bar corresponds to 100 μm. 5461

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The hydrogen bonding suppressed the crystallization rate, but at sufficient high hydroxyl concentration, in this case ∼4 mol %, the crystallization rate increased, which may be due to additional nucleation by microphase separation of OH-rich nanodomains that were predicted by molecular simulations. Hydrogen-bonding was also responsible for increasing the melt elasticity and viscosity of the copolymer compared with polypropylene. One potential application of these copolymers is capacitor films, and the improved elasticity and the lower Tα* of the hydroxyl copolymers is thought to be beneficial with regard to biaxial stretching of the film.

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AUTHOR INFORMATION

Corresponding Author

*(R.A.W.) E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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Figure 11. Storage modulus (E′) and loss modulus (E″) for PP and PPOH polymers at f = 100 Hz. Thermal transitions are observed at Tg, segmental motion; Tα*, crystal−crystal slip; and Tm, melting. Data above the melting point were calculated from small amplitude oscillatory shear (SAOS) measurements (E′ = 3G′, E″ = 3G″).

ACKNOWLEDGMENTS Financial support for this work through a Multidisciplinary University Research Initiative (MURI) grant from the Office of Naval Research (Contract No. N00014-10-1-0944) is greatly appreciated. The authors are also grateful to Prof. E. Bryan Coughlin and Brian Cromer at the University of Massachusetts, Amherst, for providing the high temperature GPC data.

α*-transition, which is typically observed between Tg and Tm in semicrystalline polymers. That thermo-mechanical relaxation is associated with slippage of the crystal planes in the direction of deformation.12 Tα* decreased substantially with the incorporation of the OH-containing monomer, but it was relatively insensitive to the concentration of comonomer used in this study. The decrease in Tα* for the PPOH copolymers may be explained by the requirement of less thermal energy for crystal−crystal slip for thinner crystals, which would be consistent with the smaller crystal size (lc) observed when the OH-containing comonomer was incorporated into PP. The activation energy of α* process estimated from the frequency dependence of Tα* indicated that activation energy for the α* process was lower in the PPOH polymers compared to that in PP. Quantitative values for the activation energy were not possible, because Tα* relaxation overlapped with the glass transition for the PPOH materials at frequencies 1 Hz or less. The α*-transition has a lower thermal activation energy than does the glass transition.

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REFERENCES

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CONCLUSIONS Copolymers of propylene and 10-hydroxy-1-undecenyl with 0− 4 mol % of the latter were synthesized by Ziegler−Natta polymerization. The copolymers were semicrystalline for hydroxyl concentrations up to at least 4 mol %, however, the introduction of the OH and the alkyl side chain lowers the melting point and crystallinity and increases the glass transition temperature. During crystallization, the hydroxylundecenyl groups were excluded from the crystals, and steric hindrance of the side chain disrupted the lamella thickening, which led to a decrease in the melting point. The crystal unit cell was unaffected by the modification. Although the melting point decreased with the introduction of the comonomer, it plateaued for copolymers with less than 2 mol % hydroxyl-undecenyl groups, which suggests that the microstructure of copolymers was somewhat blocky. Intermolecular hydrogen bonding of the hydroxyl groups in the copolymers increased the density of the amorphous phase, which produced an increase in the glass transition temperature. 5462

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