Investigation of Micromechanical Behavior and Voiding of

Article Cite This: Ind. Eng. Chem. Res. XXXX, XXX, XXX−XXX

pubs.acs.org/IECR

Investigation of Micromechanical Behavior and Voiding of Polyethylene Terephthalate/Polyethylene-stat-methyl Acrylate Blends during Tensile Deformation Bongjoon Lee,† Sebla Onbulak,‡ Yuewen Xu,∥ Vasily Topolkaraev,§ Ryan McEneany,∥ Frank Bates,*,† and Marc Hillmyer*,‡

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†

Department of Chemical Engineering and Materials Science, University of Minnesota, Minneapolis, Minnesota 55455-0431, United States ‡ Department of Chemistry, University of Minnesota, Minneapolis, Minnesota 55455-0431, United States § Corporate Research & Engineering, Kimberly-Clark Corporation, Neenah, Wisconsin 54957, United States ∥ Void Technologies Inc., Neenah, Wisconsin 54956, United States S Supporting Information *

ABSTRACT: Micromechanical deformation of polyethylene terephthalate (PET)/ethylene-stat-methyl acrylate copolymer [p(E-sMA)] blends was investigated for various MA contents and molar masses of p(E-s-MA). The copolymers were synthesized by ringopening metathesis polymerization and subsequent hydrogenation. Varying the MA content and molar mass of the copolymer alters the interfacial adhesion between the PET and the copolymer and the mechanical properties of the copolymer significantly. Transmission electron microscopy images of the blends obtained after tensile deformation reveal that the composition and the molar mass of the copolymer determine whether debonding, cavitation, both, or neither occurs during stretching. The extent of void formation associated with tensile testing was characterized by density measurements.

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INTRODUCTION

weight and enhanced breathability and in the automotive industry due to improved impact resistance.4−6 Understanding the fundamentals of voiding behavior in polymers blends is critical for controlling the microcellar structure for many targeted applications. Cavitation of semicrystalline homopolymers has been studied extensively, including in polyethylene (PE),7,8 polypropylene (PP),9 poly(L-lactide) (PLA),10 poly(ε-caprolactone) (PCL),11 polyvinylidene fluoride (PVDF),12 polyoxymethylene (POM),13 and polyamide (PA).14 During uniaxial extension, at temperatures above the glass transition temperature but below the melting point, cavities are generated in the amorphous domains between the crystalline lamellae in equatorial regions of spherulites (i.e., where lamellae are stacked parallel to the stress) and between lamellae in the polar region of the spherulite (i.e., lamellae that are stacked perpendicular to the stress).14 Cavitation, which occurs near the yield point, is affected by morphological factors such as

Dispersing modifier particles in glassy brittle polymers is a simple way of improving the mechanical property of the matrix material. Via incorporation of modifier particles with suitable physical properties, the impact strength and elongation at break can be greatly enhanced.1 Understanding toughness in multicomponent polymers is complicated by sensitivity to temperature, particle size, interparticle distance, and other morphological and micromechanical parameters.2,3 A generally accepted view is that large deformation leads to either particle cavitation or interfacial debonding, which alters the state of stress around the particles from a triaxial to biaxial state thus facilitating matrix shear yielding. This overall sequence of mechanisms enhances energy absorption resulting in improved material toughness.3 Cavitation and debonding also offer opportunities to develop nanocellular structures that may lead to improved mechanical properties and lower material density. Recently, Xu et al. reported a 14-fold increase in impact strength in colddrawn and voided polylactic acid (PLA)-based blends compared to that of the undeformed material.4 Void formation in polymer blends reduces the density leading to applications in the personal care and sportswear industry due to their light © XXXX American Chemical Society

Received: December 23, 2018 Revised: March 15, 2019 Accepted: March 17, 2019

A

DOI: 10.1021/acs.iecr.8b06362 Ind. Eng. Chem. Res. XXXX, XXX, XXX−XXX

Article

Industrial & Engineering Chemistry Research

micromechanical deformation behavior of PET/copolymer blends. We first present results that reveal how the p(E-s-MA) composition and molar mass influence the structure and linear mechanical properties of the blends, followed by an assessment of micromechanical deformation and its effect on voiding behavior.

crystal thickness, degree of crystallinity, arrangement of crystalline elements, density of tie molecules, morphology of the amorphous phase (e.g., continuity and degree of entanglement), and experimental factors such as temperature and strain rate.6 A complex interplay between these factors determines the extent of cavitation; for example, for thick and defect-free crystalline lamellae, the cavitation precedes the high shear yield stress. Numerous studies have been carried out to understand the mechanism of plastic deformation, voiding, lamellar fragmentation, crystallographic slip, and deformation of the amorphous phase in semicrystalline polymers.6 Voiding in phase-separated polymer blends is more complicated than in neat semicrystalline homopolymers due to a host of additional variables, for example, the interfacial tension between the polymer domains. Bucknall et al. have proposed models for explaining rubber particle cavitation, multiple crazing, and shear yielding for glassy and semicrystalline blends containing rubbery particles. According to their model, cavitation occurs when the strain energy released by cavitation exceeds the excess surface energy that would result from voiding; hence, the extent of cavitation strongly depends on the particle size, surface energy, and stiffness of the rubber.15−17 Declet-Perez et al. have demonstrated cavitation upon production of glassy epoxy containing in block copolymer micelles with rubbery cores (∼30 nm diameter); block copolymer micelles with a glassy core block do not cavitate.18 Kim et al. have reported various void morphologies associated with drawn particle-filled semicrystalline polymer blends prepared with different phase morphologies and variable particle−matrix interfacial adhesion.3,19 Most of the studies carried out on cavitation of polymer blends have aimed to understand the mechanism of rubber toughening rather than addressing the fundamental mechanisms of voiding or the impact of achieving lightweight materials. It is commonly known that debonding between the matrix and the particle occurs if the interfacial adhesion is poor. However, factors that alter adhesion, for example, the polarity and molar mass of the particle polymer, have not been studied carefully. To achieve targeted morphologies and desired mechanical properties from drawn polymer blends, it is crucial to systematically study the micromechanical deformation behavior that controls the initiation and growth of voids. Here polyethylene terephthalate (PET) was used as the semicrystalline polymer matrix, and a set of model statistical copolymers of ethylene (E) and methyl acrylate (MA), denoted p(E-s-MA), was employed to generate particle inclusions. PET is a commodity thermoplastic used in a multitude of applications, including food packaging and clothing. Understanding the voiding behavior of PET-based blends will have considerable impact on many commercial products. Copolymers of ethylene and methyl acrylate are often used to enhance the mechanical properties of PET such as impact strength at low temperature.20 Other ethylene copolymer rubbers have been also reported to toughen PET through cavitation and shear yielding.21−23 Due to the relatively high polarity of MA, copolymers of ethylene and methyl acrylate wet PET resulting in a tunable interfacial adhesion depending on the composition. We have systematically varied the composition and molar mass of model p(E-sMA) to investigate the effect of the interfacial adhesion, the density of tie molecules in the amorphous domains, interfacial entanglement, and the crystalline lamellae thickness on

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EXPERIMENTAL SECTION Materials and Methods. Polyethylene terephthalate (PET 7200) with an intrinsic viscosity of 0.68 dL/g was purchased from Auriga polymers and used as received. All solvents were purchased from Fisher Scientific and used as received unless noted otherwise. Anhydrous CHCl3 was purified by distillation over P2O5 prior to polymerization. All commercially available reactants and reagents were purchased from Sigma-Aldrich and used without further purification. Cyclooctene (COE) was purchased from Acros Organics. The chain transfer agent (CTA) cis-4-octene was purchased from GFS chemicals and purified by vacuum distillation over CaH2. The methyl (Z)cyclooct-4-ene-1-carboxylate (MEC) monomer was synthesized according to literature procedures.24,25 1 H nuclear magnetic resonance (NMR) spectra were recorded on Bruker AX-400 and Bruker HD-500 spectrometers using residual solvent peaks as internal standards; CDCl3 and toluene-d8 were used as the solvents. Number-average molar mass (Mn) values for the polymers were determined on a Hewlett-Packard 1100 series liquid chromatograph fitted with a Hewlett-Packard 1047A refractive index detector and three PLgel columns (Polymer Laboratories columns with 500, 103, and 104 Å pore sizes) calibrated using polystyrene standards with chloroform as the eluent at a flow rate of 1 mL/min at 35 °C. A Polymer Laboratories GPC-220 liquid chromatograph fitted with three PlGel 10 μm Mixed-B columns and equipped with a refractometer also was employed for samples with polyethylene components and operated at 135 °C with 1,2,4trichlorobenzene as the eluent at a flow rate of 1.0 mL/min. Synthesis of PMEC Polymers. Step 1: Ring-Opening Metathesis Polymerization. To a 50 mL round bottom flask with a magnetic stir bar were added MEC (and COE), cis-4octene, and anhydrous chloroform in a glovebox at a concentration of 3 M. To this solution was added Grubbs’ second-generation (G2) catalyst (0.02 mol %) as a solution in anhydrous chloroform. The vial was capped and immersed in an oil bath at 50 °C. After 20 h, the reaction mixture was cooled to room temperature, the reaction quenched with 0.1 mL of ethyl vinyl ether, and the mixture stirred for an additional 30 min. The polymer was precipitated from chloroform into methanol (twice). Methanol was decanted, and the resulting polymer was dried under reduced pressure at 30 °C for 1 day: 1H NMR (500 MHz, CDCl3) δ 5.45−5.27 (m, 4H), 3.66 (s, 3H), 2.41−2.28 (m, 1H), 2.06−1.87 (m, 8H), 1.72−1.62 (m, 2H), 1.62−1.53 (m, 2H), 1.52−1.40 (m, 2H), 1.39−1.19 (m, 8H). Step 2: Hydrogenation. The unsaturated polymers, ptoluenesulfonyl hydrazide (1.6 equiv to the double bond), tributylamine (1.75 equiv to the double bond), and BHT (20− 30 mg/g of polymer) were dissolved in xylenes (15 mL/g of polymer). The reaction mixture was refluxed at 135 °C for 8 h, allowed to cool to room temperature, and subsequently precipitated in methanol. The polymer was isolated by decantation and purified by repeating the precipitation using a chloroform/methanol/isopropanol mixture and then dried under high vacuum at 60 °C overnight to afford hydrogenated B


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Industrial & Engineering Chemistry Research polymers: 1H NMR (500 MHz, CDCl3) δ 3.66 (s, 3H), 2.37− 2.27 (m, 1H), 1.65−1.49 (m, 2H), 1.48−1.37 (m, 2H), 1.30− 1.18 (m, 24H). Preparation of Blends and Sample Preparation. Blends were prepared on an Xplore MC 5 microcompounder. Prior to blending, PET pellets were dried under vacuum at 130 °C overnight. The polymer ingredients, in all cases 93% PET and 7% by weight copolymer, were compounded at 200 rpm for 5 min at 270 °C, then extruded and quenched in water, and dried at 130 °C under reduced pressure prior to further processing. Compression-molded tensile specimens (22 mm × 5 mm × 0.2−0.5 mm) conforming to the ASTM D1708 standard were prepared on a hydraulic hot press with a plate temperature of 280 °C. Samples were kept on the bottom plate for 3 min and then compressed for 2 min. The mold was subsequently cooled on a cold water-jacketed press where the mold was held for 5 min. A PET/p(E-s-MA)/PET trilayer was prepared for T-peel testing. Two rectangular PET sheets (0.6 cm × 3 cm, 300 μm thick) were laminated on either side of a 100 μm thick layer of p(E-s-MA) using a hot press operated at 280 °C for 1 min. Subsequently, the trilayer was cooled on a cold water-jacketed press and held for 5 min. Mechanical Testing. A tensile and T-peel test was conducted after the specimens had been aged at room temperature (20 °C) for 2 days. The tensile and T-peel test was performed by using a Shimadzu Autograph AGS-X Tensile Tester. Mechanical data were averaged over 3−10 runs. Thermal Characterization. Thermogravimetric analysis (TGA) was conducted on a TA Instruments TGA Q500 apparatus. Samples were heated to 550 °C at a ramping rate of 10 °C/min under air. Differential scanning calorimetry (DSC) traces were collected on a TA Instruments Discovery DSC apparatus using a temperature ramp rate of 10 °C/min. The samples were heated to 150 °C, cooled to −85 °C, and then heated to 150 °C. In the case of PET, samples were heated to 275 °C. The melting transition temperatures (Tm) were measured by analysis of the second heating ramp. Extensional dynamic mechanical thermal analysis (DMTA) was performed using an RSA-G2 apparatus from TA Instruments. Sample were cut into rectangular shapes with thicknesses of 0.2 mm for DMTA measurement. DMTA was conducted at a ramping rate of 5 °C/min with an oscillating strain of 0.03−0.1% and an angular frequency of 1 Hz. X-ray Scattering. Small-angle X-ray scattering (SAXS) and wide-angle X-ray scattering (WAXS) were conducted with a SAXSLab Ganesha instrument outfitted with an X-ray microsource with a Cu target and position sensitive Eiger 1 M (Dectris) detector, and located at the Characterization Facility at the University of Minnesota. Data were acquired at room temperature under vacuum. Two-dimensional SAXS and WAXS patterns were azimuthally integrated to obtain plots of scattered intensity vs momentum transfer vector, q=(4π/λ) sin(θ/2), where θ is the scattering angle and λ = 1.54 Å (CuKα). Contact-Angle Measurement. Thin sheets of copolymer samples were hot pressed at 180 °C and subsequently quenched by cooling jackets with running water. Subsequently, contact-angle measurements were conducted conforming to ASTM D7490-13 using a KRUSS ADVANCE drop shape analyzer. Contact angles of small drops of distilled water and diiodomethane on the copolymer sample were measured five times at room temperature, and the surface tension was

calculated on the basis of the Owens−Wendt−Kaelble equation.26 Microscopy. Scanning electron microscopy (SEM) was conducted with a Hitachi S4700 microscope. PET/copolymer blend samples were placed in a liquid N2 bath for approximately 5 min. The samples were removed, quickly fractured with a razor blade, and then mounted on doublesided carbon tape. All blends were sputter coated with 5 nm of Ir before SEM analysis. The diameter of the dispersed p(E-sMA) phase was measured by ImageJ software and averaged over 200 particles. Transmission electron microscopy (TEM) images were acquired using a FEI Tecnai G2 Spirit BioTWIN apparatus with a LaB6 gun and an accelerating voltage of 120 kV in bright-field mode. Prior to imaging, the bulk samples were cryo-microtomed at −120 °C to 70 nm thickness sections using LEICA UC6 microtome and placed on the copper grid. Measurement of Density with a DI/NaBr Solution. A series of solutions of NaBr in deionized water (DI/NaBr) with densities ranging from 1 to 1.37 g/mL in increments of 0.013 g/mL were prepared. After the tensile bars were strained, two to three pieces from the neck were cut using a razor blade and immersed in the DI/NaBr solution. The density of the polymer was measured as the range between the density of two DI/ NaBr solutions where the sample sinks in one and floats in the other. To verify the validity of this method, polymer samples with various densities were cross calibrated using a Mettler Toledo XPE205 density measurement kit. The comparative results shown in Figure S13 confirm the good agreement.

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RESULTS AND DISCUSSION The methyl (Z)-cyclooct-4-ene-1-carboxylate (MEC) monomer was prepared by Pd-catalyzed carboxymethylation of 1,5cyclooctadiene (1,5-COD) following previously reported procedures.24,25 The homopolymerization of MEC was performed in anhydrous chloroform using 0.02 mol % G2 at 50 °C (Scheme 1). cis-4-Octene was used as a chain transfer Scheme 1. Synthesis of p(E-s-MA) Polymers

agent (CTA) to control the Mn of the resulting polymers. In the 1H NMR spectrum of the polymer PMEC, the chemical shift at 3.7 ppm corresponds to −OCH3 groups of the methyl acrylate (MA) units. The olefinic protons of the monomer between 5.7 and 5.6 ppm are shifted upfield and merged into a broader peak with a shoulder at 5.3 ppm, indicating that the ring was opened and both E and Z double bonds were present in the polymer backbone.25 ROMP of the 5-substituted C


Article

Industrial & Engineering Chemistry Research Table 1. Characterization Data for the Synthesized Copolymer of Ethylene and Methyl Acrylateg name

FMA (wt %)

Mn (kDa)a

Đ

Tg (°C)b

Tm (°C)c

Td,5% (°C)d

X (%)e

p(E-s-MA(0))-20 p(E-s-MA(0))-90 p(E-s-MA(10))-23 p(E-s-MA(10))-55 p(E-s-MA(10))-93 p(E-s-MA(20))-19 p(E-s-MA(20))-33 p(E-s-MA(20))-79 p(E-s-MA(50))-59

0 0 10 10 10 20 20 20 50

20 90 23 55 93 19 33 79 59

3.7 2.0 3.9 3.0 2.3 2.5 2.2 2.8 2.3

−120 −119 −36 −34 −26 −33 −30 −65 −43

116 131 95 103 104 91 89 83 −

302 331 381 385 386 306 329 351 340

51 39 22 23 24 22 25 13 −

E (MPa)f 352 255 73.7 61.3 85.8 60.2 58.6 13.3 1.3

± ± ± ± ± ± ± ± ±

19 19 6.6 1.3 4.1 2.9 2.6 0.1 0.1

UTS (MPa)f 10.5 8.3 11.5 18.4 31.1 6.5 13.7 6.7 0.10

± ± ± ± ± ± ± ± ±

0.9 0.6 0.3 0.9 4.0 1.2 0.7 1.6 0.05

εb (%)f 130 265 544 630 685 331 624 512 555

± ± ± ± ± ± ± ± ±

64 225 6 26 102 115 39 103 143

Number-average (Mn) and weight-average (Mw) molecular weights determined by 1,2,4-trichlorobenzene SEC at 135 °C using universal calibration and the Mark−Houwink coefficients for high-density polyethylene. bGlass transition temperature was determined by DSC during 2nd heating at a ramping rate of 10 °C/min and DMTA at a ramping rate of 5 °C/min. Complete lists of DMTA curves, DSC curves are shown in Figure S20 and S21. Comparison of Tg values from DMTA and DSC are tabulated in Table S3. cPeak position of the melting peak during the second heating used for determining Tm, where the ramping rate was 10 °C/min. dA 5% mass loss determined by TGA at a rate of 10 °C/min under air. eThe crystallinity of the copolymer was determined using the enthalpy of fusion of HDPE with 100% crystallinity to be 293 J/g.29 fThe elastic modulus, ultimate tensile strength (UTS), and strain at break reflect the average from at least three experiments in which the variation represents the range of the data. gComplete lists of tensile curves for each samples are shown in Figure S16. a

cyclooctene monomers gave regio- and stereoirregular polymers consistent with the literature.27,28 Unsaturated polymers were then hydrogenated using diimide generated in situ by the thermolysis of p-toluenesulfonylhydrazide. The hydrogenation was quantitative according to the 1H NMR spectrum, where the peaks in the olefinic region between 5.4 and 5.3 ppm were absent, and the allylic methylene protons of the end groups originally at 2.0−1.9 ppm shifted upfield (Figure S1). We also prepared statistical copolymers with COE (Scheme 1). Varying the feed ratio between the COE and MEC was used to tune the properties of the final unsaturated copolymers. Long reaction times (20 h) in the ROMP step were utilized to promote an even sequence distribution along the backbone by cross-metathesis. The crude 1H NMR spectra of the copolymers indicated that MEC and COE were fully incorporated after 20 h. As a result, control over the monomer composition and the final molar mass of the polymers was achieved. The copolymers were hydrogenated in the same manner as described above for PMEC homopolymers. The copolymer backbone was saturated quantitatively on the basis of the 1H NMR spectra (Figure S2). Hydrogenated polymers were analyzed by high-temperature SEC in 1,2,4-trichlorobenzene at 135 °C. The characterization data of copolymers synthesized are summarized in Table 1. The differential scanning calorimetry (DSC) curves for the p(E-s-MA) copolymers and PET are given in Figure 1. All of the curves were measured upon a second heating cycle at a ramping rate of 10 °C/min. The thermal properties of copolymers, including Tm, Tg, and Td,5%, are listed in Table 1. As the number of MA units in the copolymer increases, the melting temperature (Tm) decreases and the peak shape becomes much broader due to the distribution of the ethylene sequence length in the chain. This suggests that the crystalline morphology of the copolymer underwent a transition from large lamellae with a uniform size distribution to small lamellae with a broad size distribution.30 As the MA content reaches 50 wt %, the copolymer can no longer form crystalline domains and shows only a glass transition at −46 °C. In addition to the MA content, the molar mass of the copolymer also plays a role in the melting transition. For 0 and 10 wt % MA content copolymers, the Tm of the highest-molar mass samples is 15 and 10 °C higher than that of the lowest-molar mass samples,

Figure 1. Differential scanning calorimetry curves of p(E-s-MA) series and PET upon a second heating cycle at 10 °C/min. As the content of the noncrystallizing MA co-monomer, FMA (wt %), increases, the melting transition (Tm) of the copolymer resulting from PE crystallinity becomes broader due to the broad distribution of sequence lengths. Also, the Tm of the copolymer shifts to lower temperatures, suggesting that the average crystallite thickness decreases. The Tm of PET is 254 °C. The curves are shifted vertically for the sake of clarity.

respectively. In contrast, a 20 wt % MA content copolymer with a higher molar mass shows a lower Tm. This suggests that for low MA content, the longer chains can fold into larger lamellae because they have a lower number of chain ends; however, for high MA content, the mobility of chains is more important for the formation of larger lamellae because of the noncrystallizing MA co-unit. The Tg was determined by DSC and DMTA by the peak position of the loss modulus. Due to the relatively small change in heat capacity, Tg was not detected by DSC for some of the copolymers; however, DMTA was able to detect Tg for all samples. Comparisons between Tg values measured by DMTA and DSC are listed in Table S3, and they are in good agreement. The complete sets of DMTA and DSC curves for the copolymers are shown in Figures S20 and S21, respectively. The Tg of p(E-s-MA) ranged from −120 to −26 °C. PET showed a melting point at 254 °C by DSC. D


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the amorphous domain stays almost constant with an increase in FMA. The wide-angle X-ray scattering (WAXS) data for the series of quenched p(E-s-MA) are shown in Figure 3. Two strong

The small-angle X-ray scattering (SAXS) traces obtained from quenched p(E-s-MA) are shown in Figure 2. The samples

Figure 2. Small-angle X-ray scattering (SAXS) data for the series of quenched p(E-s-MA) copolymers. The SAXS traces show either one or two maxima corresponding to the long spacing of lamellar stacking. As the content of the noncrystallizing co-unit, FMA (wt %), of copolymer increases, the long spacing of lamellar stacking decreases and finally disappears at FMA = 50 wt %. The curves are shifted vertically for the sake of clarity.

Figure 3. Wide-angle X-ray scattering (WAXS) data for a series of quenched p(E-s-MA) copolymers. Two strong scattering peaks corresponding to the (110) and (200) planes of the orthorhombic polyethylene unit cell are observed for FMA = 0 wt % at q values of 1.525 and 1.686 Å−1, respectively. These reflections broaden and decrease in intensity with an increase in MA content in the copolymer. The (110) and (200) peaks shift to lower q values with an increase in MA concentration. This shift results from expansion of the a-axis of the unit cell due to crystal strain at the interphase of the crystallites by accumulation of noncrystallizing MA units.34 The copolymer with FMA = 50 wt % shows only an amorphous halo. The curves are vertically shifted for the sake of clarity.

were hot pressed at 180 °C, and the mold was subsequently cooled on a cold water-jacketed press where the mold was held for 5 min. Crystalline and amorphous domains have different densities because polymer chains are more closely packed in the crystalline form, and therefore, the periodic stacking of crystalline and amorphous domains leads to Bragg scattering. Hence, SAXS can characterize the long spacing of the lamellar stacking composed of crystalline and amorphous domains from the primary peak position. Copolymers with FMA values of 0, 10, and 20 wt % show either one or two maxima that correspond to the long spacing of the lamellar stacking. The copolymer with an FMA of 50 wt % does not show any maximum as it is unable to form crystalline lamellae, consistent with the DSC data. Two maxima are observed for the FMA = 0 wt % copolymer. We interpret these features as the smaller long period (L1) corresponding to well-defined lamellae and amorphous layers and the larger long period (L2) being due to thin defective lamellae formed at the later stage of cooling from the melt. These thin lamellae are produced between the thick lamellae formed at the earlier stage of cooling and contain defects such as chain ends and branching. Therefore, the electron density difference between these thin defective lamellae and the amorphous material is strongly reduced and may not contribute to the scattering in some regions. As a result, larger long periods are detected by SAXS if one neglects these structures that are contained between well-defined thick lamellae.31,32 Only one maximum was observed for the FMA = 10 and 20 wt % copolymers. The long spacing of the p(E-sMA) copolymers, plotted in Figure S4, decreases with an increase in FMA from 0 to 20 wt %, suggesting smaller lamellae are formed. The thickness of the crystalline lamellae and amorphous domains can be calculated on the basis of the long spacing from the SAXS analyses and crystallinity from the DSC measurements. The procedure and the results, described in the Supporting Information and plotted in Figure S7, show that the thickness of the crystalline lamellae decreases while that of

scattering peaks that correspond to the (110) and (200) planes of the orthorhombic polyethylene cell are observed at q values of 1.525 and 1.686 Å−1, respectively, for FMA = 0 wt %.33 The scattering peaks broaden and decrease in intensity with an increase in FMA, suggesting that the copolymer becomes less crystalline and the sizes of the crystalline domains decrease. This is consistent with the DSC and SAXS data. Shifting of the (110) and (200) peaks to lower q values with an increase in MA content (see Figure S5) is attributed to expansion of the aaxis of the orthorhombic polyethylene unit cell, presumably due to accumulation of the co-monomer (MA) on the surface of the thin crystals as reported for random copolymers of ethylene and 1-decene, 1-hexane, 4-methylpentene, and dicyclopentadiene.34 We note that the slightly dilated PE crystals may be more susceptible to lattice fragmentation, thus serving as potential cavitation sites. We conclude from the DSC and X-ray scattering measurements that the composition and molar mass of the copolymer significantly influence the crystalline morphology on multiple length scales, most notably the distribution of crystalline lamellar thickness and the packing of PE repeat units within the crystalline domains. These morphological factors affect the mechanical properties of the blends as shown below. Stress−strain curves obtained from the p(E-s-MA) copolymers at a strain rate of 5 mm/min are plotted in Figure 4. The samples were hot pressed at 180 °C, and the mold was subsequently cooled on a cold water-jacketed press where the mold was held for 5 min. Dog-bone samples were fabricated on the basis of the ASTM D1708 standard. Likely due to the lower crystallinity of the higher-FMA copolymers, the stress− E


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S6, the fraction of tie molecules increases from 5% to 26% as the molar mass increases from 19 to 79 kg/mol for FMA = 20 wt %. Increasing the tie molecule content enhances the tensile toughness at a constant crystallinity; hence, the tensile toughness improves in going from 19 to 33 kg/mol. However, the crystallinity decreases as the molar mass is further increased to 79 kg/mol due to smaller crystallites. Small crystallites are more susceptible to either fragmentation or shear yielding during reorientation of the chains when they are subjected to a tensile strain, reducing their effectiveness as cross-links between tie molecules. A similar argument applies to explain the higher tensile toughness of the FMA = 10 wt % sample compared to that of the FMA = 20 wt % systems with a comparable crystallinity. The smaller crystallite sizes at FMA = 20 wt %, evidenced by the lower melting temperatures in Figure 1, imply mechanically weaker crystals that cannot withstand high stress. Further increasing FMA to 50% results in amorphous material that does not strain harden and fails at low stress levels; Figure S8 shows the data in Figure 4 plotted on a log scale to discern the FMA = 50 wt % result. The composition of the p(E-s-MA) copolymers also influences the surface properties of the materials. Thin sheets of p(E-s-MA) were hot pressed at 180 °C, and the mold was subsequently cooled on a cold water-jacketed press where the mold was held for 5 min. The surface tension was estimated by contact-angle measurements conforming to ASTM D7490-13. The polar component, dispersion component, and total surface tension are plotted in Figure S9. The polar component of the surface tension increases, and the dispersion component decreases with an increase in MA content. Therefore, improved adhesion between PET and p(E-s-MA) is expected for copolymers with higher MA content as PET also has polar functional groups along the backbone. The total surface tension, which is the sum of the polar and dispersion contributions, stays almost constant between 32 and 37 mJ/ m2. The total surface energy is relevant to particle cavitation in blends because cavitation is accompanied by the formation of new surfaces.15,16 As the total surface tension is almost constant, we can assume that the surface energy contribution of the voids is negligible. We have employed Wu’s equation to calculate the interfacial tension between PET and p(E-s-MA)37

Figure 4. Stress−strain curves of quenched p(E-s-MA) copolymers taken at room temperature. The dog-bone samples were strained at a rate of 5 mm/min. Depending on the MA content, the copolymer exhibits thermoplastic to elastomeric behavior. See Figures S8 and S16 for a logarithmic version of this plot and the full list of tensile runs.

strain curve changes from thermoplastic to elastomeric behavior with an increase in FMA. The elastic modulus decreases significantly from 355 MPa for FMA = 0 wt % (Mn = 20 kg/mol) to 1.4 MPa for FMA = 50 wt % (Mn = 39 kg/ mol). The complete set of elastic moduli, ultimate tensile strengths (UTSs), and elongation at break values for the series of p(E-s-MA) copolymers is listed in Table 1. The elongation at break increases significantly with MA content from 184% for FMA = 0 wt % (Mn = 20 kg/mol) to 543% for FMA = 10 wt % (Mn = 23 kg/mol). The UTS for the higher-molar mass FMA = 10 wt % sample is greater than that for the corresponding lower-molar mass sample with the same composition, which we attribute to the number of tie molecules that interconnect adjacent lamellae. A linear dependence on the number of entangled tie molecules with molar mass has been reported for polyethylene.35 Tie molecules significantly affect the strength of the amorphous regions, and higher-Mn copolymers are more likely to have longer chains that can connect adjacent lamellae; note the long spacing of the lamellar stacks is comparable as shown in Figure S4. The concentration of tie molecules can be calculated on the basis of the approach reported by Huang and Brown, who assume that the tie molecules will form if the endto-end distance of the polymer chain is longer than the critical distance determined by the thickness of the lamellae and amorphous domain.36 Figure S7 shows the thickness of the crystalline lamellae and amorphous composite calculated from the long spacing and the crystallinity measured by DSC.38 It is interesting to note that while the thicknesses of the crystalline domains decrease, the thickness of the amorphous regions remains relatively constant. The fraction of tie molecules, which increases with molecular weight for FMA > 0 wt %, is shown in Figure S6. On the basis of our calculation, for FMA = 10 wt % with similar crystallinities, the tie molecule fraction increases from 5% to 27% with an increase in molar mass from 23 to 93 kg/mol. This result supports the observed tensile property of the FMA = 10 wt % systems and suggests that not only crystallinity but also the tie molecule content plays a role in controlling the tensile properties. As FMA increases to 20%, the stress−strain curves show lower UTS values compared to those of the FMA = 10% specimens (Table 1). For FMA = 20 wt %, the UTS and elongation at break increase as the molar mass increases from 19 to 33 kg/mol and then decreases for 79 kg/ mol (Table 1). This decrease can be attributed to the reduced crystallinity for the Mn = 79 kg/mol sample, which might occur due to the reduced chain mobility of the longer chains during crystallization. On the basis of our calculation shown in Figure

γ12 = γ1 + γ2 −

4γ1dγ2d γ1d + γ2d

−

4γ1pγ2p γ1p + γ2p

(1)

where γ12 is the interfacial tension, γi is the surface tension, and γdi and γpi are the dispersion and polar components of γi, respectively. Figure 5 shows that, as expected from the surface tension measurements, the interfacial tension between PET and p(E-s-MA) decreases with an increasing level of the polar MA co-monomer in the copolymer. Therefore, the interfacial adhesion between PET and the copolymer is expected to be stronger as the MA content increases. To experimentally measure the interfacial adhesion, T-peel testing was performed. The peel strengths of PET/p(E-s-MA)/ PET trilayer samples are plotted in Figure 6. Peel tests were conducted after samples had been aged at room temperature (20 °C) for 2 days. A picture of a T-peel test is shown in Figure S10. The T-peel strength differs by 2 orders of magnitude depending of the composition and molar mass of the copolymer as illustrated in Figure 6. As expected from the surface tension measurements, the peel strength between PET and p(E-s-MA) increases with an increase in MA content in F


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copolymer interface produced during quenching from the melt for the longer chains. Peel strength, which reflects interfacial adhesion, will be compared with the blend micromechanical deformation behavior described below.38 The morphology of the PET(93 wt %)/p(E-s-MA)(7 wt %) blends was characterized via SEM as shown in Figure 7. A complete set of SEM images for all blends is provided in Figure S12. In all cases, the SEM images reveal a particulate morphology, where the particle diameter depends on both the molar mass and the MA content of the copolymer. The decrease in particle size with an increase in MA content is attributed to the associated reduction in interfacial tension (Figure 5), consistent with the enhanced adhesion measured by peel testing (Figure 6). The molar mass of the copolymer also plays a role in the particle size as it affects the particle viscosity; a lower molar mass results in the formation of smaller particles. Particle size is one of the most important factors affecting cavitation in the blend as the total surface energy contribution becomes less important for larger particles.17 Also, the critical volume strain for cavitation increases with a decrease in the particle size. However, the sizes of the particles in the blends used in this study far exceed the lower limit for cavitation (30 nm) as reported for rubber-toughened nylon6.15 The stress−strain curves for PET and the PET/p(E-s-MA) blends are presented in Figure 8, and the associated mechanical properties are summarized in Table 2. PET showed low crystallinity in both quenched pristine PET and the PET/p(E-s-MA) blends as shown by DSC in Figure S23. However, the p(E-s-MA) phase is characterized by a semicrystalline state as shown in Figure S21 possibly due to the low Tg. The crystallinity of the p(E-s-MA) phase dispersed in the PET matrix was inferred from the quenched pristine p(E-s-MA) data because this quantity could not be extracted from the blend due to the overlap of the p(E-s-MA) melting transition with the glass transition and crystallization temperature of PET. Pristine PET has a Young’s modulus (E) of 1.73 GPa; it yields at ∼35 MPa, undergoes a significant necking with an associated 40% drop in drawing stress in the plateau region, and then plastically deforms. Strain hardening starts at ∼200% strain, and the specimen breaks at ∼340% strain. Addition of copolymer inclusions impacts mechanical properties in various ways depending on the molar mass and composition (FMA) of the copolymer. The stiffness of the blended material is reduced by only 0−15% (E = 1.47−1.73 MPa) due to the low modulus of the copolymer (Table 1); the yield stress drops by 0−30% (σy = 25−38 MPa), and the stress in the plateau region of the plastic zone (defined here at 50% strain) decreases slightly by approximately 2−20%. However, the strain at break is broadly distributed between 146 and 368% with a significant strain hardening for FMA ≥ 10 wt % blends. Tensile toughnesses of a majority of the blends were comparable with that of pristine PET; however, FMA = 0 wt % blends showed 15−60% lower toughnesses. In situ tensile SAXS measurements show an abrupt increase in the overall scattering intensity beyond the yield point in the blend samples as quantified by the scattering invariant (Figure S14). This derives from the development of voids due to either debonding at the particle−matrix interface or cavitation from within the particles. TEM images of strained PET/copolymer blends are shown in Figure 9. Because of the significant difference in density between the copolymer and PET, no staining was required to

Figure 5. Interfacial tension between PET and p(E-s-MA) based on eq 1. The interfacial tension decreases with an increase in the content of the polar MA co-monomer. The number next to the symbol denotes the Mn of the copolymer.

Figure 6. Peel strength of PET/p(E-s-MA)/PET. Two rectangular sheets (0.6 cm × 3 cm, 300 μm thick) of PET were laminated with p(E-s-MA) (100 μm thick) at 280 °C for 1 min and pulled apart at a rate of 10 mm/min. Peel strength increases with an increase in the MA content of the copolymer because the polarity of acrylate improves the compatibility with PET. The dotted line shows the peel displacement beyond which the peel strength becomes a plateau for all samples. Pictures show the laminates of p(E-s-MA(0))-90, p(E-sMA(10))-55, p(E-s-MA(20))-79, and p(E-s-MA(50))-39 after testing. Laminates with low MA content show smooth PET surfaces suggesting low adhesion, while laminates with high MA content show rough and deformed surface indicative of cohesive failure.

the copolymer. Peel strengths (S) determined by the plateau region of the T-peel strength curves are listed in Table 2. The peel strength is also influenced by the molar mass of the copolymer as seen by the results for the p(E-s-MA(20)) specimens in Figure 6. At the higher peel strengths associated with the FMA = 20 and 50 wt % samples, the debonded surfaces become rough, which is indicative of cohesive failure. Smooth surfaces are produced with the FMA = 0 and 10 wt % samples. A complete set of pictures of the surfaces after T-peel testing is shown in Figure S11. The dependence of peel strength on the molar mass of the copolymer appears to be more sensitive for the higher-MA content copolymers. The effect of molar mass on peel strength for the FMA = 20 wt % samples is very distinct. The copolymer with the highest Mn shows 5 times the peel strength of the two lower-Mn specimens. We attribute this feature to greater chain entanglement at the PET and G


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Table 2. Peel Strengths of PET/p(E-s-MA)/PET Laminates and Mechanical Properties of PET/p(E-s-MA) Blends Measured by Tensile Testingg copolymer none P(E-s-MA(0)-20) P(E-s-MA(0)-90) P(E-s-MA(10)-23) P(E-s-MA(10)-55) P(E-s-MA(10)-93) P(E-s-MA(20)-19) P(E-s-MA(20)-33) P(E-s-MA(20)-79) P(E-s-MA(50)-59)

103Sa (N/mm) − 4 8 10 10 14 65 61 313 92

± ± ± ± ± ± ± ± ±

2 2 1 1 1 16 9 67 38

Eb (GPa) 1.73 1.71 1.63 1.63 1.64 1.56 1.61 1.73 1.47 1.57

± ± ± ± ± ± ± ± ± ±

0.04 0.13 0.02 0.05 0.09 0.09 0.10 0.08 0.10 0.07

σyc (MPa) 36 38 30 32 32 31 32 38 25 32

± ± ± ± ± ± ± ± ± ±

3 2 1 2 2 2 3 4 4 3

σpd (MPa) 23 23 17 20 21 20 20 20 17 20

± ± ± ± ± ± ± ± ± ±

8 4 2 1 1 2 1 3 2 3

εbe (%) 344 146 263 350 290 368 334 364 302 328

± ± ± ± ± ± ± ± ± ±

31 64 38 16 31 24 21 29 34 20

toughnessf (MJ/m3) 92 35 51 95 89 95 87 123 66 82

± ± ± ± ± ± ± ± ± ±

18 14 9 10 7 12 5 24 9 3

a

Peel testing was conducted at a rate of 10 mm/min and averaged in the plateau region of the measurement. bElastic modulus; measured from the linear proportion of the stress−strain curve (initial 1%). cPlateau stress. dYield stress. eElongation at break. fToughness determined by integrating the stress−strain curve to the point of break. gError bars denote the range of the data.

Figure 8. Stress−strain curves of PET/p(E-s-MA) blends measured in tension. The inset shows the response at low strain levels. The dogbone samples were deformed at a rate of 5 mm/min. A full list of tensile runs is plotted in Figure S18.

particles as described above. For the FMA = 0 wt % blends, debonding is observed for most of the particles in the blend with an Mn = 20 kg/mol copolymer (Figure 9a). Most of the particles are undeformed, and voids are formed at the interface between the particle and the matrix. However, for blends with a higher copolymer molar mass (Mn = 90 kg/mol), debonding is not as prevalent (Figure 9b). This is attributed to either the lower crystallinity or more intensive interfacial entanglement between the PET matrix and particle due to a larger radius of gyration for this copolymer. A smaller elastic modulus for the Mn = 90 kg/mol copolymer due to a lower crystallinity (Table 1) implies that the particles are more susceptible to deformation rather than debonding during deformation. Elongated particles aligned parallel to the macroscopic strain direction also suggest that the interface is more entangled presumably due to the longer chains. This is supported by the observation by TEM of a thin layer at the interface between PET and the particles. We believe this thin layer is composed of extended and entangled chains of the copolymer and PET. Some debonding is still observed but far less than is evident with the lower-molar mass copolymer. Details about the debonding in the blends with the FMA = 0 wt % copolymer are presented in Figure S15. For the FMA = 10 wt % copolymers, interfacial adhesion appears to be more robust. As one can see in panels c and d of Figure 9, both debonding and cavitation within the particles are observed for the Mn = 23 and 55 kg/mol blends, respectively. Debonding forms voids at the interface between the particle and the matrix due to interfacial failure. Cavitation

Figure 7. Scanning electron microscopy (SEM) images of cryofractured specimens showing the particulate morphology of (a) PET/ p(E-s-MA(0))-20 and (b) PET/p(E-s-MA(50))-59. The weightaverage particle diameters of PET(93 wt %)/p(E-s-MA)(7 wt %) were measured from the SEM images and are plotted in panel c. The error bars denote the standard deviation from measurements of 160 particles (on average) for each blend. Increasing the MA content of the copolymer improves the compatibility with PET, and as a result, the average particle size decreases. The scale bar denotes 2 μm. The full list of particle distributions for the PET/p(E-s-MA) blends is presented in Figure S17. The number-average particle diameters of PET(93 wt %)/p(E-s-MA)(7 wt %) are presented in Figure S24.

distinguish them. Various micromechanical deformation mechanisms are observed, which are closely related to the crystallinity, mechanical properties, and adhesion of the H


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micromechanical deformation behavior of the blend from interfacial debonding to bulk cavitation, with the molar mass playing a critical role on the deformation behavior for low-MA content copolymer blends. Void formation during tensile deformation is reflected in optical properties such as whitening of the gauge area. Pictures of the strained tensile bars are shown in Figure S22. Most of the tensile bars showed whitening and opacity in the gauge area as expected from the TEM images. The PET/p(E-sMA(10))-93 blend displayed the smallest amount of whitening, consistent with our observation by TEM that this sample neither cavitates nor debonds. Micromechanical deformation behavior, including cavitation and debonding of toughened semicrystalline polymer systems, improves the impact strength by facilitating the shear yielding of the matrix.19 Even though we have not observed the significant enhancement of the tensile property as shown in Figure 8, a preliminary result shows enhanced notched Izod impact strength in the PET/p(E-s-MA) blend system. Because of the limited amount of synthesized p(E-s-MA), a commercially available p(E-s-MA) with similar molecular characteristics (FMA = 20 wt %; Mn = 24 kDa; Đ = 5.3) was blended at 7 wt % with PET. The SEC traces and characterization data of Elvaloy 1820 AC are shown in Figure S19 and Table S2. The tensile properties of PET and PET/ Elvaloy 1820 AC are comparable as shown in Figure S19 and Table S1, similar to PET/p(E-s-MA) systems. However, the notched Izod impact strength increased by 50% upon addition of Elvaloy 1820 AC. We believe that this increase results from the cavitation of the Elvaloy 1820 AC particles as observed with the FMA = 20 wt % system. One of the objectives of this study is to understand the molecular and structural parameters that maximize the formation of stable voids in strained PET through the inclusion of 7% p(E-s-MA) particles by weight. Density measurements provide a convenient way to quantify the consequences of void formation through debonding and cavitation. As shown in Figure 10, the density of pristine

Figure 9. Transmission electron microscope images of (a) PET/p(Es-MA(0))-20, (b) PET/p(E-s-MA(0))-90, (c) PET/p(E-s-MA(10))23, (d) PET/p(E-s-MA(10))-55, (e) PET/p(E-s-MA(10))-93, (f) PET/p(E-s-MA(20))-19, (g) PET/p(E-s-MA(20))-33, (h) PET/p(Es-MA(20))-79, and (i) PET/p(E-s-MA(50))-59 after being strained to 40−60%. Dark, gray, and white domains are the PET, copolymer, and void, respectively. The strain rates were 5 mm/min, and the arrow denotes the direction of tensile strain. The insets show highermagnification images. The scale bar denotes 1 μm.

produces voids within other particles in response to the development of hydrostatic stress, in response to triaxial strain, which exceeds the intrinsic strength of the particle. Debonding leaves undeformed particles separated from the matrix by void space. Cavitation is accompanied by affine particle deformation because the interface remains stitched to the matrix. Surprisingly, the Mn = 93 kg/mol (FMA = 10 wt %) copolymer blend displays neither cavitation nor debonding as seen in Figure 9e. We attribute this to the combined effects of enhanced interfacial entanglement and the superior mechanical properties of the copolymer as shown in Figure 4. This copolymer has the highest σy and UTS among the FMA = 10 wt % samples, which apparently reduces its susceptibility to cavitation, complemented by greater interfacial strength due to enhanced entanglements. Increasing the MA content (FMA) to 20 wt % results in cavitated voids within elongated particles and relatively little interfacial failure. As shown in T-peel test, the interfacial adhesion is high for this MA content, which results in less debonding. However, fragmentation and cavitation of the semicrystalline lamellae within the particles appear to be more facile as evidenced in Figure 9f−h. Unlike the FMA = 10 wt % copolymers, increasing the molar mass to 79 kg/mol still results in substantial cavitation, which is attributed to the small and defective crystals associated with the FMA = 20 wt % copolymers; i.e., these are more susceptible to fragmentation leading to a breakdown of the tie molecule network. Increasing the MA content to 50 wt % results in only cavitated particles and no debonded particles as shown in Figure 9i, consistent with strong interfacial adhesion and weak particle cohesive strength. Overall, increasing the MA content transforms the

Figure 10. Density of PET/p(E-s-MA) after different extents of strain. Density measurements were taken during immersion in DI/NaBr solutions (see Experimental Section and Figure S13). The PET density increases with extension due to strain-induced crystallization, while void formation coincident with yielding leads to a decrease in density. Strain hardening above roughly 150% extension leads to an increase in density as a consequence of narrowing of the void space.

PET increases with strain by ∼1.5% compared to that of the undeformed state due to strain-induced crystallization. For the PET/p(E-s-MA) blends, the density initially drops when the blends are strained beyond the yield point due to void formation and then either remains constant or increases slightly as the material strain hardens for an ε of >150% (see Figure 8). In the strain hardening regime, the particles elongate I


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Different compositions and molecular weights of the copolymers were synthesized by ROMP and subsequent hydrogenation to alter the surface chemistry and the mechanical properties of the inclusions. As the MA content in the copolymer increases, the interfacial tension calculated by Wu’s equation decreases. Interfacial adhesion between the PET and copolymer measured by the T-peel test was consistent with the interfacial tension calculated by Wu’s equation. The variance in the interfacial adhesion and mechanical properties of the copolymer has a significant effect on the micromechanical deformation mechanisms of the blend as verified by TEM. As the MA content in the copolymer increases, at first it shows only debonding, the coexistence of debonding and cavitation is observed, and finally only cavitation is observed. Also, the molar mass of the copolymer plays a significant role as it alters the interfacial entanglement and the number of tie molecules in the copolymer. The density measurement of the blend after strain was an efficient tool for characterizing the extent of voiding, and maximum voiding was observed for the PET/copolymer blend that showed both debonding and cavitation.

and decrease in diameter in response to the reduction in the width of the gauge section of the tensile specimen. A host of factors contributes to the overall final density of the material, including the copolymer composition, molecular weight, and average particle size. Several useful design principles emerge from these results. We focus first on the blend that produced virtually no change in density during drawing in tension. The PET/p(E-sMA(10))-93 blend teaches us that the mechanical properties of the inclusions play an important role in optimizing void formation. The behavior of PET/p(E-s-MA(10))-93 is also associated with absence of void formation, in this case apparently due to a combination of high particle strength associated with a high molar mass (Figure 4) and sufficient particle−matrix adhesion to suppress interfacial debonding due to the MA content. Here, the particle size [d = 1.53 μm (Figure 5)] cannot be implicated as it lies intermediate between those of the two best performing blends with identical MA content (see below). This limiting case provides important bounds on the criteria for designing particles capable of creating voids in strained PET. Additional insight can be gleaned from the other blends, although the trends are subtle and less definitive. PET/p(E-sMA(10))-55 displays the best performance, with the largest drop in density after yielding, which is sustained throughout the strain hardening regime. This material appears to void through both cavitation and debonding (see Figure 9d), perhaps offering a mechanism for continuous generation of cavities during the entire stretching process. We also note that this blend contains among the largest particles [d = 2.1 μm (Figure 5)], which should correlate with the greatest stress concentration, although this alone does not explain the best performance. The PET/p(E-s-MA(0))-90 blend has the same particle size but generates less than half the change in density. The role of molar mass displayed by the three 10 wt % MA copolymers is especially intriguing. We speculate that a fortuitous combination of properties produces the optimal result: (i) intermediate copolymer mechanical strength directly linked to the semicrystalline morphology governed by FMA and molar mass and (ii) intermediate adhesion controlled by the FMA thus permitting both stress transfer between the matrix and particle while under triaxial strain and debonding. Increasing or decreasing FMA around the 10 wt % level leads to a decline in performance. Although FMA = 0 wt % (blends in which debonding dominates) and FMA = 20 and 50 wt % (blends in which cavitation dominates) samples display comparable densities over intermediate strain (before strain hardening begins), the void density is more retained in the FMA = 0 wt % blend than in the FMA = 20 and 50 wt % blends when stretched beyond the strain hardening strain (ε > 150%). The ability to retain the void density during the strain hardening regime depends on whether the void has internal structures to prevent the closing of the void. In blends in which cavitation dominates, the void is formed by the failure of the copolymer particle, and as a result, voids lack any internal structure. However, in blends in which debonding dominates, particles are undeformed and that might prevent the collapse of the void during the strain hardening regime.

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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.iecr.8b06362. 1 H NMR spectra; SEC traces; DSC curves; data related to the crystallinity, morphology, mechanical properties, and surface energy of the polymers and blends; SEM and TEM images and associated particle size distributions; and in situ tensile SAXS measurements of polymer blends (PDF)

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AUTHOR INFORMATION

Corresponding Authors

*E-mail: [email protected]. *E-mail: [email protected]. ORCID

Frank Bates: 0000-0003-3977-1278 Marc Hillmyer: 0000-0001-8255-3853 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The authors acknowledge Kimberly-Clark for financial support. Partial support of this work was provided by the Center for Sustainable Polymers, a National Science Foundationsupported Center for Chemical Innovation (CHE-1413862), at the University of Minnesota. Parts of this work were carried out in the Characterization Facility, University of Minnesota, which receives partial support from the National Science Foundation through the MRSEC program. The authors thank Chris Frethem and Han Seung Lee for helping with the SEM imaging. The authors also thank Dr. Alex Todd for helpful discussion.

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REFERENCES

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Investigation of Micromechanical Behavior and Voiding of

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