Role of Crystallization on Polyolefin Interfaces: An Improved Outlook

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Role of Crystallization on Polyolefin Interfaces: An Improved Outlook for Polyolefin Blends Alex M. Jordan,† Kyungtae Kim,† Diego Soetrisno,† Jennifer Hannah,† Frank S. Bates,† Shaffiq A. Jaffer,‡ Olivier Lhost,§ and Christopher W. Macosko*,† †

Department of Chemical Engineering and Materials Science, University of Minnesota, Minneapolis, Minnesota 55455, United States Total American Services, Inc., 1201 Louisiana St., Suite 1800, Houston, Texas 77002, United States § Total Research and Technology Feluy, Zone Industrielle Feluy C, 7181 Seneffe, Belgium ‡

S Supporting Information *

ABSTRACT: Polyolefins including linear low density polyethylene (LLDPE), high density polyethylene (HDPE), and isotactic polypropylene (iPP) account for nearly 2/3 of the worldwide plastics market. With wide-ranging applications, often short term in nature such as packaging, recycling of polyolefins is becoming increasingly important in developing a sustainable worldwide plastics market. However, it is difficult to separate polyolefins in mixed recycle streams; it would be advantageous to melt blend them, but their immiscibility leads to blends with poor properties. Here we demonstrate the role of synthetic history (i.e., site specific metallocene vs heterogeneous Ziegler−Natta catalyzed) on the oligomer content of HDPE, LLDPE, and iPP and its influence on adhesion between PE and iPP. Using a range of polymers and processing conditions, we identify four classes of such interfaces with a wide range of interfacial adhesion strengths (GIC): excess oligomer (GIC < 30 N/m), easy chain pullout (GIC ≅ 100 N/m), kinetically trapped entanglements (GIC ≅ 600 N/m), and crystallization across the interface (GIC > 1200 N/m). Using molecular weight distribution data, we identified a critical oligomer content where the interfacial failure mechanism transitions from cohesive failure (GIC > 1200 N/m) to adhesive failure (GIC ≅ 100 N/m). Transmission electron microscopy (TEM) and atomic force microscopy (AFM) highlight distinct interfacial semicrystalline morphologies for each class of polyolefin interface which are defined by molecular parameters and processing conditions. Polyolefin blends were compression molded to highlight the role of interfacial strength in blends formed from mixed polyolefin streams; weak interfaces resulting from excess oligomer buildup yielded brittle failure while superior interfacial adhesion resulted in ductile blend failure. plastics results in materials with reduced tensile, flexural, and impact properties.7 It is widely recognized that the nature of the interfaces in heterogeneous polymer−polymer blends plays a critical role in determining the associated mechanical properties.8 A number of strategies have been employed to improve the interfacial and bulk mechanical properties of PE/iPP blends derived from mixed stream recycling. These include physical compatibilization with date palm leaf fibers9 or modified natural zeolite,10 reactive compatibilization,11−13 inclusion of various olefin

1. INTRODUCTION Polyolefins command a dominant position in today’s plastics market with PE alone capturing a market share of 35% in the U.S. during 2014.1 These market trends are attributable to low cost and a remarkable diversity of applications and properties derived from the methods used to synthesize and process the materials.2,3 Despite the generation of over 33 million tons of plastic from these fossil resources, the overall recycling rate for plastics within the U.S. in 2014 was only 9.5%.4,5 In addition to the low recycling rate, plastic recycling is also limited by existing sorting technologies for mixed stream recycling, which typically do not achieve perfect separation of PE and iPP.6 Subsequent melt processing of recycled blends containing these © XXXX American Chemical Society

Received: January 27, 2018 Revised: March 14, 2018

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DOI: 10.1021/acs.macromol.8b00206 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules Table 1. Manufacturer Supplied Data for Polyolefins Used in This Work

LLDPE HDPE iPP

manufacturer

grade name

sample code

density, ρ [g/cm3]

melt flow index (190 °C/2.16 kg) [g/10 min]

Exxon Exxon Total Total Total Exxon Total Exxon

LL3003.32 Exceed 3518 HD6081 M6040 MR2001 Achieve 3854 PPH9059 PP4052

zlE mlE zhE mhE miPa miPb ziPa ziPb

0.918 0.918 0.960 0.960 0.905 0.900 0.905 0.900

3.2 3.5 8 4

melt flow index (230 °C/2.16 kg) [g/10 min]

25 24 25 2.0

copolymers,14−18 and inclusion of iPP-b-PE block copolymers.19 Poon and colleagues demonstrated the role of comonomer on PE/iPP adhesion. Using coextruded LLDPE/ iPP films, they showed that inclusion of octene with mole fraction increasing from 2.4 to 7.4% led to a systematic increase in PE/iPP adhesion.20 However, our understanding of the structure and properties of the interfaces between immiscible semicrystalline PE and iPP plastics is incomplete. Considerable effort has been directed at understanding the structure of chain-folded PE lamellae at the PE−iPP interface. Epitaxial growth of crystalline PE from oriented semicrystalline iPP has been determined to result in an angle of ±50° between the b-axis of the iPP unit cell and the chain-folded PE lamellae.21−25 Further work on the epitaxial growth model by Gohil26 showed this leads to PE crystalline lamellae lying parallel to the PE−iPP interface. In contrast, Chaffin et al.27 demonstrated deviation from the epitaxial growth model, which were traced to the localization of noncrystallizable oligomer accumulated at the interface in linear low density polyethylene (LLDPE) and iPP laminates and blends. LLDPE and iPP produced with heterogeneous Ziegler−Natta catalyst are characterized by a broad molar mass distribution and a relatively large amount of oligomer. In one departure from the epitaxial growth model, amorphous polymer migrates to the interface, mechanically decoupling the two semicrystalline materials. Conversely, homogeneous site-specific metallocene catalysts produce a narrower molar mass distribution and relatively little oligomer, which leads to intimate contact between semicrystalline LLDPE and iPP, including penetration of lamellae crystals across the interface. The penetration of LLDPE crystals across the interface then appears as crystals perpendicular to the PE−iPP interface which deviates significantly from the epitaxial growth model. This form of interfacial interaction was shown to improve interfacial adhesion in laminates by at least a factor of 2027 and resulted in a dramatic improvement in the mechanical properties of LLDPE−iPP blends.28 Recent work by Nui et al. demonstrated deviation from the idealized epitaxial structure in LLDPE and iPP crystallized under the effects of a flow field,29 and Zhou and colleagues have reinforced this conclusion based on experiments with injected-molded blends of PE and iPP.30 Taken together, these results suggest that factors such as molecular architecture and processing have a significant impact on polyolefin interfacial crystal structure, which governs interfacial adhesion and the macroscopic mechanical properties of polyolefin blends.27−30 Through the combined use of lamination (analogous to industrial heat sealing processes) and continuous coextrusion, we have investigated the thermodynamic and kinetic aspects of formation of the PE−iPP interface. This study has employed

iPP, LLDPE, and high density polyethylene (HDPE) with a variety of molar mass distributions and oligomer contents to probe the role of molecular architecture on interfacial crystalline morphology. By fabricating bilayers from a continuous coextrusion process, we have explored a range of processing conditions and probed the kinetic driving forces associated with interfacial crystallization. Several distinct interfacial morphologies developed while systematically varying the material and processing parameters. These have been correlated with interfacial adhesion, a critical factor in designing blends of immiscible polymers. Understanding the complex interplay between molecular architecture and processing conditions on interfacial morphology and adhesion provides an avenue to more efficiently use mixed stream recycling feeds for specific product development.

2. EXPERIMENTAL SECTION 2.1. Materials. Eight commercial polyolefins were used in this study, including 4 iPP, 2 HDPE, and 2 LLDPE grades supplied by either Total or ExxonMobil (Table 1) and used as received. 2.2. Molar Mass Determination. Weight-average molar mass (Mw), number-average molar mass (Mn), dispersity (Đ), and fraction of chains below 3 times entanglement weight (3Me) were determined via high temperature size exclusion chromatography (HT-SEC, PLGPC220). Each PE resin was dissolved in 1,2,4-trichlorobenzene (TCB, Fischer Chemical, HPLC grade O484-6) at a concentration of 1.5 mg/mL at 135 °C. Each iPP resin was dissolved in TCB at a concentration of 1.5 mg/mL at 160 °C. Retention volume curves were obtained at an eluent flow rate of 1 μL/min at temperatures of 135 °C (PE) and 160 °C (iPP). Retention volume was converted to molar mass distributions via calibration with narrow polystyrene (PS) standards and Mark−Houwink−Sakurada constants.31−33 2.3. Rheology and Fitting. Small-amplitude oscillatory shear (SAOS) measurements for were obtained over a frequency range 0.01−100 s−1 using a 25 mm parallel plate rheometer (Thermal Analysis, ARES) in the linear viscoelastic shear strain region at temperatures (T) from 180 to 220 °C. Temperatures between 140 and 180 °C were used to characterize sample mlE due to its low melting temperature. Steady shear viscosities (η) were obtained on a dual bore capillary rheometer (Malvern R7) over a nominal shear rate (γ̇) range of 20−2000 s−1. Both rheometer barrels were fitted with 1 mm diameter orifices. The first barrel was fitted with an orifice of 15 mm length. The second orifice was 0.25 mm in length and was used to obtain a value for the entrance pressure used in the Bagley correction applied to obtain a true viscosity measurement.34 Steady shear viscosity and SAOS viscosity were overlaid using the Cox−Merz rule to obtain viscosity vs shear rate data over the range 0.01 to ∼2000 s−1 and fitted with the Cross−Arrhenius model (eqs 1 and 2).35,36

η=

η0 1 + (τrγ )̇ 1 − n

η0 = Ae Ea / RT B

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Macromolecules Fitting with the Cross model provides a zero-shear viscosity (η0), critical shear stress (τr), and power-law index (n). Fitting with the Arrhenius model provides an estimate of the flow activation energy (Ea) and an Arrhenius constant (A). 2.4. Melting Temperature and Crystallinity. Heating and cooling curves were obtained for 5 mg samples of each polyolefin in hermetically sealed aluminum pans using differential scanning calorimetry (DSC, TA Q1000). Samples were equilibrated at −90 °C for 5 min before heating to 200 °C at a rate of 10 °C/min to erase any processing history. Samples were then cooled at 10 °C/min to −90 °C to obtain a crystallization temperature range (Tc), which was defined over the range of nonlinearity of the heat capacity of each polyolefin (Figure S6). A second heating to 200 °C at 10 °C/min was used to obtain the melting enthalpy (ΔHf) and melting temperature (Tm). The crystalline fraction (Xc) of each polyolefin was estimated (eq 3) by comparing ΔHf to the enthalpy of melting for a 100% crystalline sample (ΔHf0) of PE (293.6 J/g) and iPP (207 J/g).37,38

Xc =

ΔHf × 100% ΔHf 0

during T-peel testing. It was not possible to operate the gear pumps at a more extreme ratio than 34:3 due to the unreliability of pump rotation at speeds of 1 and 2 rpm. 2.7. T-Peel Adhesion Testing. T-peel testing was based on slight modifications to ASTM standard D1876.43 Bilayer films produced via coextrusion were cut to 12.5 mm wide (w) strips to remove any edge wrap-around resulting from viscosity mismatches during coextrusion. Each trimmed bilayer sample was then initially separated with a razor blade to produce one PE and one iPP tab to be held in the tension fixtures during testing. Laminated bilayer films were fabricated with exposed tabs as described above. For consistency, PE was placed in the top movable grip of the peel testing apparatus (TA Instruments RSAG2) while iPP was placed in the stationary bottom fixture. Force (F) was recorded over a minimum peel length of 25 mm with constant peel velocity of 5 mm/min. Plateau force values were normalized by sample width (F/w) which approximated the critical energy release rate (GIC) during delamination of each individual polymer pair and set of processing conditions.41 2.8. Delaminated Surface Analysis. The delaminated surfaces were investigated by scanning electron microscopy (SEM) without further treatment. The films were examined by a Hitachi S4700 field emission scanning electron microscope (FE-SEM) with accelerating voltage of 1.5 kV at a working distance of ca. 25 mm and 45° sample tilt angle with respect to the incident electron beam. 2.9. Transmission Electron Microscopy (TEM). Coextruded PE−iPP were bulk-stained by immersing in RuO4 aqueous solution (prepared by dissolving 0.2 g of RuCl3 into 2 mL of sodium hypochlorite solution, both purchased from Sigma-Aldrich) for 3 h. Cross sections (thickness: ca. 70−100 nm) were obtained by cryomicrotoming at −140 °C using a Leica EM UC6 Ultramicrotome equipped with cryogenic chamber. The slices were cut using a Diatome diamond knife and transferred to TEM grids using saturated sucrose solution and a perfect loop (Electron Microscopy Sciences, Hatfield, PA). Samples were imaged using a Tecnai G2 Spirit BioTwin microscope with an accelerating voltage of 120 kV. 2.10. Atomic Force Microscopy. Polyolefin interfaces were also imaged using atomic force microscopy (AFM; Nanoscope III with Multimode system, Digital Instruments, Santa Barbara, CA). Films were cut perpendicular to the extrusion direction to image edge-on using the same cryo-ultramicrotome at −120 °C with a series of progressive cuts. Initial cutting was carried out at a step length of 1 μm and a glass knife velocity of 6 mm/s. Each cutting step reduced the step length and knife velocity until a final setting of 50 nm step length and 1 mm/s velocity were reached. A series of finishing cuts were also made at 50 nm step length and 1 mm/s velocity with a Diatome diamond knife. The microtome process ensured a smooth imaging surface for AFM in tapping mode. Samples were examined in the repulsive regime using a silicon tip (resonant frequency = 166 kHz, spring constant = 2 N/m, and radius = 8 nm). 2.11. Heat Transfer Simulation. The cooling profile in the coextruded films was solved as a transient two-dimensional heat transfer problem using an ANSYS Workbench 17.0 software. Typical values based on iPP for thermal conductivity (0.27 W/(m K)), melt density (0.7 g/cm3), and thermal diffusivity (0.1 mm2/s) were assumed for all polymers.44 Heat capacity (cp) was taken from the DSC data described above. Upon leaving the die, heat transfer from the extrudate was assumed to occur due to convection with air moving at a velocity equal to the speed at which the extrudate was being taken up by the chill rolls. This resulted in an estimated heat transfer coefficient (h) of 10 W/(m2 K).45 The measured time for the extrudate to leave the exit die at 180 °C before reaching the counter rotating chill rolls was 2.8 s. During this 2.8 s period, h was assumed constant, 10 W/(m2 K), along both x and y directions of the extrudate with ambient air temperature of 22 °C. The extrudate spent approximately 0.75 s sandwiched between the two chill rolls. During this time, the heat transfer coefficient between both surfaces of the polymer film and the rolls was taken as 1500 W/ (m2 K) with a surface temperature of 5 °C.46 Finally after passing through the chill rolls, the extrudate is again cooled by air with the same heat transfer coefficient (i.e., h = 10 W/(m2 K)).

(3)

2.5. Polyolefin Lamination. Individual PE and iPP films were first prepared via compression molding (Carver Auto Series Compression Press) at 180 °C under 6670 N of force for 5 min in a 120 × 60 × 0.4 mm (length × width × thickness) picture frame mold between Teflon sheets. After 5 min, the polyolefin films were cooled under ambient conditions and rinsed with chloroform (CHCl3, Fischer Chemical C2984) to remove any residue transferred from the Teflon sheets. Following film fabrication, bilayer films were formed via compression molding. Care was taken during the lamination process following CHCl3 wash to keep the PE and iPP film surfaces clean. To laminate each bilayer, a Teflon sheet was placed on a metal plate. A single iPP film was laid on the Teflon sheet. A second Teflon sheet was then used to cover a 10 mm portion across the length of the iPP film before a single PE sheet was aligned with the iPP sheet and laid in place (Figure S1). This allowed two nonlaminated tabs to be preserved for T-peel testing, while also allowing sufficient peel length to determine an equilibrium peel force. A third Teflon sheet was placed above the PE film, and a second metal platen was then placed on the Teflon sheet. This stack was then heated again to 180 °C and held under 6670 N of force for 1 min and then allowed to cool under ambient conditions. 2.6. Polyolefin Coextrusion. Bilayer PE−iPP films were also fabricated via a continuous lab-scale coextrusion process. This equipment and process have been described elsewhere.39−42 In brief, one stream of PE was layered above a single stream of iPP to produce a bilayer film with a final exit temperature of 180 °C. Overall volumetric flow rate was kept constant at 32 cm3/min. Melt contact time was calculated based upon the continuity equation. With die dimensions 50 mm × 1.2 mm and volumetric flow rate ∼32 cm3/min, the average linear flow velocity was calculated to be 8.9 mm/s. The die land length was measured to be 63 mm, giving an approximate melt interfacial contact time in the die of 7 s. After leaving the exit die, bilayer films were quenched on counter-rotating chill rolls with a gap of 0.8 mm between the cooled steel rolls. A gap between the exit die lip and rolls required ∼2.8 s before quenching, giving a total melt interfacial contact time of ∼10 s. The final dimensions of the extrudate after quenching were 22 mm × 0.8 mm. The gear pump metering iPP flow displaced 1.6 cm3/rotation, while the gear pump metering the PE flow displaced 0.8 cm3/rotation. The pump RPM were set to give 16 cm3/min for each stream. This setup enabled bilayer film fabrication with individual layer thickness comparable to those of each laminated film. These 50:50 (PE vol:iPP vol) films were fabricated from mhE− miPa, mlE−miPb, and zhE−ziPa for comparison with lamination with gear pumps speeds of 20 rpm for PE and 10 rpm for iPP. By varying volumetric ratios, we shifted the interface position during extrusion. Nominal volume ratios were chosen as 60:40 (19.2:12.8 cm3/min; 24:8 rpm), 70:30 (22.4:9.6 cm3/min; 28:6 rpm), 80:20 (25.6:6.4 cm3/min; 32:4 rpm), and 85:15 (27.2:4.8 cm3/min; 34:3 rpm). Isotactic polypropylene was chosen as the thinner layer due to its higher tensile strength to assist in maintaining mechanical integrity C

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Macromolecules 2.12. Bilayer Film Recrystallization. Coextruded bilayer films of mhE−miP and zhE−ziP of each volume composition were heated to 180 °C between two heated plates with no compressive force. A 1 mm thick picture frame mold of lateral dimensions equal to the sample was placed around the coextruded film during recrystallization to limit flow and preserve sample geometry. After 5 min hold time, films were removed from the heat source and allowed to cool under ambient conditions. Recrystallized films were tested using identical T-peel procedure as described above. 2.13. Polyolefin Blends. Pelletized blends of 75/25, 50/50, and 25/75 (PE vol %/PP vol %) were prepared using a twin-screw extruder (Prism, Thermo, 16 mm diameter, L/D = 24/1) for zhE−ziP, mhE−miPa, and mhE−miPb. Tensile samples were molded into dogbone geometries (22 mm × 4.8 mm × 0.6 mm) at 180 °C under a compressive force of 6670 N for 5 min. The dogbone specimens were uniaxially elongated at a testing speed of 5 mm/min (Shimadzu AGSX, 500 N load cell), and the stress vs strain response was recorded.

chains below 3Me in mlE; however, this does not mean there is a complete absence of noncrystallizable chains or chains that melt at low temperature. Results from extraction techniques like temperature rising elution fractionation (TREF)48 or chromatographic cross-fractionation (CCF)49 will differ from a molecular weight cutoff criteria. There are several sources of noncrystallizable chains in polyolefins: low molecular weight chains, waxes, high atactic copolymers and high α-olefin random copolymers. The noncrystallizable fraction in mlE and zlE was found by CCF to be 0.7% and 13.5%, respectively.50 These values are significantly higher than the 3Me values in Table 2, but they scale by a factor of about 10. In addition to molar mass characterization, each polyolefin molecular structure was characterized through a combination of shear and linear rheology (Figures S4 and S5, Table S1). By fitting with a Cross−Arrhenius model, Ea can be estimated, which provides insight into the degree of long chain branching (LCB) in each PE grade. Typical literature values for low LCB density (1−5 LCB per 10 000 carbons) in HDPE fall within the range 22−28 kJ/mol,51 while Ea of unbranched LLDPE is typically below 33 kJ/mol.52 Values obtained from Cross− Arrhenius fitting suggest that all of the PE grades used here are linear. This result is not surprising as HDPE and LLDPE are characterized by their lack of long chain branching. Finally, Tm, Xc, and a crystallization range were determined via DSC (Figure S6), which will be used in subsequent discussion of interfacial crystallization kinetics. 3.2. Molecular Architecture and PE−iPP Adhesion. From the lamination results (Figure 1a) we see that the GIC values fall into three distinct categories roughly an order of magnitude apart from each other. The highest GIC values are for mlE laminated to both grades of metallocene iPP: miPa and miPb. In each, a cohesive failure mechanism was observed during peel testing. GIC for both mlE−miPa and mlE−miPb pairs was greater than the strength of the mlE homopolymer, which is consistent with Chaffin et al.27 A decrease in GIC of at least an order of magnitude was observed when mhE was laminated to either miPa or miPb with no significant difference between the adhesion of the mhE−miPa (GIC = 122.6 ± 7.9 N/ m) and mhE-miPb (GIC = 135.3 ± 6.8 N/m) pairs. Taken together, these results suggest that the grade of metallocene iPP did not influence GIC with either mlE or mhE. With similar

3. RESULTS AND DISCUSSION 3.1. Molecular Characterization. Each polyolefin was characterized to determine molar mass and dispersity via HTSEC (Figures S2 and S3). The first readily apparent trend is the tendency of the Ziegler−Natta polyolefins (zhE, zlE, ziPa, and ziPb) to have a broad molar mass distribution compared to the metallocene polyolefins (mhE, mlE, miPa, and miPb). As a means to estimate the relative content of oligomer for each polyolefin, the fraction of chains below 3Me was determined. On the basis of literature values for Me (Me,HDPE = 1.0 kDa; Me,LLDPE = 1.4 kDa; Me,iPP = 5.4 kDa),47 it is also apparent that each Ziegler−Natta polyolefin has an appreciable fraction of chains below 3Me (Table 2). Interestingly, there were tensile strength of LLDPE) was obtained when site specific metallocene catalysts were used and LLDPE replaced HDPE (Entangled Crystals, Figure 10). This is attributed to the nearly complete absence of oligomer in the mlE material. Furthermore, we have identified a critical threshold of noncrystallizable chains as a condition for degrading the intrinsically high strength interfaces associated with PE and iPP. This low concentration of oligomers is sufficient to disrupt cross-interfacial crystallization and reduce LLDPE−iPP adhesion by at least 1 order of magnitude. This critical finding highlights the importance of molar mass screening when considering the recyclability of PE to prevent oligomer contamination in mixed stream recycle feeds. An intermediate adhesive strength was recorded for laminates of metallocene catalyzed HDPE and iPP (Easy Chain Pullout, Figure 10). We attributed this intermediate GIC to a reduced oligomer content, which allowed small pockets of cross interface crystallization of HDPE lamellae to form and be relatively easily pulled out. Importantly, when crystallization occurred over time scales comparable to the self-diffusion time scale of HDPE and iPP chains, entanglements were trapped across the HDPE−iPP interface which increased GIC by about 275% (Kinetically Trapped Entanglements, Figure 10). The kinetic nature of this increased adhesion was demonstrated by recrystallization under ambient cooling conditions and a return to the equilibrium semicrystalline morphology and adhesive strength (Easy Chain Pullout in Figure 10). These findings demonstrate the importance of microscopy in understanding how molecular structure, oligomer content, and processing conditions influence adhesion between immiscible polyolefins. Finally, the role of interfacial adhesion in polyolefin blends is highlighted: poor PE−iPP adhesion lead to brittle failure of polyolefin blends while excellent adhesion lead to mechanical synergism between PE and iPP. It is envisioned that understanding concepts such as a critical oligomer content will influence polyolefin catalysis and reaction conditions when designing polyolefins. Knowledge of critical processing parameters may also inform the blending process of polyolefins for tailored applications and smarter use of mixed recycle streams.

Figure 9. Uniaxial tensile response of zhE−ziP (a), mhE−miP (b), and mlE−miP (c) compression-molded blends at 25/75, 50/50, and 75/25 volume fraction along with each respective homopolymer.

homopolymers (Figure S16). Similar trends were observed in injection-molded zhE−ziP blends (Figure S17). Conversely, tensile strength and toughness of mlE−miP extruded blend sheets were greater than each constituent homopolymer with elongation comparable to mlE (Figure S16); however, in the injection-molded blends this trend was only observed for mlE25/miP75 (Figure S17). I

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Macromolecules

Figure 10. Proposed schematic representation of the four classes of polyolefin interfaces (PE: red; iPP: blue) discussed in this paper based upon microscopy.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acs.macromol.8b00206. Lamination schematic, molar mass distributions, rheological fitting, thermal characterization, delaminated surface analysis, heat transfer simulation solution, and results for extruded and injection-molded blends (PDF)

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AUTHOR INFORMATION

Corresponding Author

*E-mail [email protected]; phone (612) 625-0092 (C.W.M.). ORCID

Frank S. Bates: 0000-0003-3977-1278 Christopher W. Macosko: 0000-0002-2892-3267 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This research was supported by a grant from Total S. A. with partial support by the Industrial Partnership for Research in Interfacial & Materials Engineering (IPRIME). Polymers were provided by Total and ExxonMobil Corporation. We thank Dr. Pat Brant of ExxonMobil for helpful discussions and the CCF data. Parts of this work were carried out in the Characterization Facility, University of Minnesota, which receives partial support from NSF through the MRSEC program.

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REFERENCES

(1) Witter, D. IBISWorld Industry Report 32521: Plastic & Resin Manufacturing in the US; Industry Report; IBISWorld Inc.: 2015. (2) Vasile, C., Ed.; Handbook of Polyolefins, 2nd ed.; CRC Press: 2000. (3) U.S. Energy Information Administration. Petroleum & Other Liquids - Product Supplied; US Department of Energy: Washington DC, February 28, 2018. (4) United States Environmental Protection Agency. Advancing Sustainable Materials Management: 2014 Fact Sheet; EPA: Washington, DC, November 2016. (5) The Association of Plastic Recyclers. 2015 United States National Postconsumer Plastic Bottle Recycling Report; American Chemicstry Council: Washington, DC, 2016. (6) Luijsterburg, B.; Goossens, H. Assessment of plastic packaging waste: Material origin, methods, properties. Resour. Conserv. Recy. 2014, 85, 88−97. J

DOI: 10.1021/acs.macromol.8b00206 Macromolecules XXXX, XXX, XXX−XXX

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DOI: 10.1021/acs.macromol.8b00206 Macromolecules XXXX, XXX, XXX−XXX