Effects of Polyethylene Molecular Weight Distribution on Phase

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Effects of Polyethylene Molecular Weight Distribution on Phase Morphology Development in Poly(p‑phenylene ether) and Polyethylene Blends Jun Wang,† Andy H. Tsou,‡ and Basil D. Favis*,† CREPEC, Department of Chemical Engineering, É cole Polytechnique de Montréal, Montréal, Québec H3T 1J4, Canada Global Chemical Research, ExxonMobil Chemical Company, Baytown, Texas 77520, United States

† ‡

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S Supporting Information *

ABSTRACT: Bimodal molecular weight distribution (MWD) in a polymer can potentially deliver mechanical toughness from the high molecular weight (MW) components without compromising its processability due to the presence of the low MW components. The influence of MW bimodality on the morphology development in a heterophase polymer blend system, however, remains unclear. In this study, high-density polyethylenes (HDPEs) of high unimodal, low unimodal, and bimodal MWDs, namely HDPE-H, HDPE-L, and HDPE-B, were synthesized and then blended with PPE to study the effect of the MWD on morphology development. After blending, the low MW components in HDPE-B are found to populate the HDPE/ PPE interface, and the morphological and continuity behaviors of PPE in HDPE-B/PPE are virtually identical to those in the unimodal HDPE-L/PPE blend. Despite the high interfacial tension between HDPE and PPE, all three HDPEs form fiber-like dispersions in HDPE/PPE blends when PPE is the majority phase. This leads to an early HDPE phase continuity onset and an extremely broad cocontinuity region from 10 to 60 wt % of HDPE in all the HDPE/PPE blend systems. These fiber-like HDPE dispersions in the PPE matrix are likely a consequence of the extremely high PPE viscosity. The study opens up the concept/ potential of using biomodal polymers in such a way that the higher molecular weight component could be used to control mechanical properties while the lower molecular weight component controls the viscous/morphological characteristics.

1. INTRODUCTION Polymer blend morphologies dictate the blend properties. Controlling the blend morphology development to arrive at optimal morphologies in polymer blends is of critical importance.1−6 There are principally two types of morphologies in binary polymer blends; one is the dispersed phase/ matrix morphology, and the other is the cocontinuous morphology.2 Characterized by two phases being interconnected throughout the blends, the cocontinuous morphology is of great interest for many advanced applications such as conductive materials, tissue engineering scaffolds, separation membranes, and so forth.1,3,7−9 The key material properties and mixing conditions that define morphology development in polymer blends include interfacial tension, composition, viscosity ratio, elasticity, shear rate, and flow fields.10 Classic theories on the dispersion of a Newtonian immiscible liquid/liquid system suggest that a droplet more easily deforms and breaks up when the viscosity ratio (p, which is the ratio of the dispersed phase viscosity to the matrix viscosity) of the blend components is in the range 0.1−1.11,12 Grace found that the droplet only shows limited deformation (without breakup) beyond a viscosity ratio of 3.5.12 However, in a polymer melt blend system, this viscosity ratio upper threshold has been shown to be well above 3.5.13,14 The reason © XXXX American Chemical Society

for this phenomenon is still unclear. One possible explanation is due to the elongational flow involved during polymer processing (e.g., from extruders) that is more effective to break up the droplets.12,14 Favis and Chalifoux demonstrated that a decreasing viscosity ratio, p, from 17.3 to 4.5 in a polypropylene/polycarbonate (PP/PC) blend can lead to a phase size reduction of 3−4 fold.13 Nevertheless, a minimum droplet size was achieved at a range of p from 0.15−1, which is close to what Grace found in Newtonian systems. Everaert et al. studied the influence of viscosity ratio on the phase morphology in polymer blends and found that droplet breakup is enhanced when the matrix viscosity is high (i.e., p ≪ 1) due to a combination of effective shear stress transfer from the matrix to the droplet and increased dispersive forces.14 In blends when the matrix viscosity is low (i.e., p > 1), droplets cannot be finely dispersed by the lubricating low-viscosity matrix with insufficient dispersive forces. Breakup and coalescence have been proposed to be responsible for the morphology evolution in a polymer blend. First, the polymer phases are sheared and extended Received: July 7, 2018 Revised: October 20, 2018

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DOI: 10.1021/acs.macromol.8b01445 Macromolecules XXXX, XXX, XXX−XXX

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ties, and reduced cost. Moreover, no reported studies, to our knowledge, have examined the influence of MW bimodality on the morphology development of a heterophase polymer blend. In this work, PPE blended with unimodal low MW HDPE (HDPE-L), unimodal high MW HDPE (HDPE-H), and bimodal HDPE (HDPE-B) were prepared, and the effects of MW and MWD on morphology and phase continuity development in binary HDPE/PPE blends are examined.

into sheets at the initial stage of mixing followed by sheet breakup into connected filament structures.15−17 These extended structures can then either undergo further division through capillary instabilities into droplets or remain as stable fibers to form a cocontinuous structure. Afterward, these droplets or fibrous domains could coalesce.18,19 In this context, Li et al. proposed a model based on the lifetime of the dispersed phase morphologies, droplets or fibers, which, in turn, is determined by the interfacial tension.19 In a high interfacial tension blend, the dispersed phase is dominated by a droplet morphology and the continuity is developed through droplet−droplet coalescence with higher phase continuity onset, whereas in a low interfacial tension blend, stable fibrous structures can be found and the coalescence among fibers (high aspect ratio leads to lower percolation threshold) results in a lower phase continuity onset and a broad cocontinuity region. Hence, the formation of stable fibrous structures is essential in developing a cocontinuous morphology at a low blend composition or phase concentration. Poly(p-phenylene ether) (PPE), or poly(2,6-dimethyl-1,4phenylene oxide), is synthesized by oxidative coupling polymerization of 2,6-dimethylphenol (2,6-xylenol). It is an engineering polymer with many desirable properties of high heat distortion temperature and strength, excellent dimensional stability, and flame retardancy. However, PPE is very viscous in melt state and is difficult to process. Commercially, all PPEs are blended with other polymers to improve their processability. PPE/PS blends are most commonly used since PPE and PS are miscible. However, PPE/PS blends have limited solvent resistance due to the amorphous characteristics of both PPE and PS.20 At the same time, the high heat distortion temperature (HDT) of PPE is compromised by PS plasticization. To overcome these drawbacks, PPE has been blended with other immiscible polymers such as polyamides, polyesters, and polyolefins.21−26 Because of the high viscosity of PPE, it has a tendency to stay dispersed in blends even at relatively high PPE concentrations. Li and Shimizu examined the morphology of a PPE/polyamide (PA) blend and found that PPE is dispersed within the PA matrix at 50%.21 However, by adding 5% clay into this 50/50 PPE/PA blend, a cocontinuous morphology developed. This cocontinuous morphology transformation arises from a high PA phase viscosity due to selective clay partition into the PA phase. Polyolefins (POs) are another commonly used immiscible blend component with PPE which provide low cost, excellent processability, and high chemical resistance.22−25 However, POs and PPE are highly incompatible, and compatibilizers are typically employed to prepare commercial PPE/PO blends. Akkapeddi and VanBuskirk studied compatibilized PP/PPE blends with viscosity ratios (PPE to PP) ranging from 6 to >100.22 In all 50/50 PP/PPE blends, PPE remains dispersed due to its higher viscosity. Gowan reported that the mechanical properties of PPE were virtually unaffected by blending in up to 5% PE.27 However, the morphology and interfacial properties were not examined. Bimodal HDPEs with a bimodal molecular weight distribution (MWD) and composition distribution could provide high-stress-crack resistance and strength, from the high MW and comonomer content components. As well, improved processability and stiffness can be derived from the low MW and comonomer content components.28 Blending PPE with a bimodal HDPE may deliver high-performance PPE blends with excellent processability, low-temperature proper-

2. EXPERIMENTAL SECTION 2.1. Materials. PPE (CAS# 25134-01-4) was obtained from Sigma-Aldrich, and its Mn is 15000 with a polydispersity index (PDI) of 2 as measured by a gel permeation chromatograph (GPC) equipped with a differential refractive index detector (DRI). The calibration was performed using PS standards. Both HDPE-L and HDPE-H were synthesized using a laboratory gas phase reactor with a spray-dried silica supported catalyst formed by combining C2 symmetric metallocene of bis(1-methyl-3-butylcyclopentadienyl)zirconium dichloride with excess methylaluminoxane (MAO) activator. The MW was controlled by the hydrogen amount. The bimodal HDPE (HDPE-B) was produced using a laboratory singlegas-phase reactor system with spray-dried silica-supported dualorganometallic catalyst comprised of bis(2-pentamethylphenylamido)ethyl)zirconium dibenzyl (a non-metallocence) and (tetramethylcyclopentadienyl)(n-propylcyclopendienyl)zirconium dichloride (a metallocene) in a 3 to 1 molar ratio. A modified methylaluminoxane, MMAO, is also fed into the reactor as the catalyst activator. No comonomers were used to prepare HDPE-L and HDPE-H, and both are homopolymers. Their MW and MW distributions were thus determined using GPC-3D (gel permeation chromatography−three detectors) where MWs were measured by DRI. The hexene comonomer was added in HDPE-B at a small amount (2−3 wt %). Correspondingly, a GPC-IR (infrared detector), with the ability to measure comonomer content, was employed to determine the MW and MWD of HDPE-B in addition to hexene comonomer content. 2.2. Rheology and Blend Preparation. All polymers were compression molded into disks for rheological evaluations using a MCR 301 rheometer (Anton Paar, Austria) with a parallel-plate configuration. A strain sweep was first performed to identify the linear viscoelastic region. A subsequent frequency sweep was performed from low frequency to high frequency at a strain of 5−10% under small-amplitude oscillatory shear mode. An N2 purge was used to minimize the degradation of the polymers during the rheology tests. A gap of 1 mm was employed, and a typical test time was ∼10 min. All the HDPE/PPE blends with different weight fractions were meltprocessed on an Xplore MC5 conical miniature twin-screw extruder equipped with a recirculation channel running at 260 °C at 100 rpm for 5 min. After processing, filament samples with a diameter of ∼2.5 mm were cooled in air and collected for characterization. 2.3. Phase Morphology and Continuity Characterization. All blends were either cryo-microtomed or cryo-fractured for morphology examination. Prior to evaluation using atomic force microscopy (AFM) and scanning electron microscopy (SEM), samples were cryofaced using a cryo-microtome (Leica) at −120 °C. Because of the hardness of these samples, peak force tapping mode was selected as the primary AFM imaging mode using an Icon AFM (Bruker) whereas a desktop SEM (Phenom) was employed to obtain their electron micrographs without metal coating. To prevent charging on uncoated, cryo-faced samples during desktop SEM imaging at 5 keV, no high magnifications >10000× were used. A Scanning Probe Image Processing Software, SPIP (Image Metrology), was then utilized to process both AFM and SEM micrographs and to quantify phase sizes and phase size distributions. Some blends were cryo-fractured for field emission SEM (FE-SEM) (JEOL JSM 840) examinations. After cryofracturing, a gold layer was deposited onto the fractured surface using a Polaron SC502 sputter coater to prevent charging followed by morphology characterization with the FE-SEM running at 2−5 keV. B

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Macromolecules For phase continuity determinations, blend samples of ∼100 mg were immersed in 20 mL of chloroform for 30 h at room temperature, with the solvent being refreshed three times, to selectively extract the PPE phase. After extraction, each sample was dried at 80 °C under vacuum for 48 h, and the continuity of PPE was determined by the following equation: continuity of PPE (%) initial weight − weight after extraction = × 100% weight of PPE calculated from formulation

Table 1. MWs and MWDs of HDPE-H, HDPE-B, and HDPE-L polymers

Mn

Mw

Mw/Mn

HDPE-H HDPE-Ba HDPE-L

84000 15000 10000

268800 367500 78000

3.2 24.5 7.8

a

The peak MWs in the bimodal HDPE-B are 23840 (61.8 wt % low MW components) and 372910 (38.2 wt % high MW components).

(1)

2.4. Dynamic Mechanical Analysis (DMA). A DMA 2980 dynamic mechanical analyzer (TA Instruments) was used to evaluate the dynamic mechanical properties of the HDPE/PPE blends under the temperature range from −50 to 260 °C at a heating rate of 3 °C/ min. The tests were performed under the dual cantilever clamp mode. A frequency of 1 Hz and an oscillation amplitude of 20 μm were employed.

the gas phase reactor which lowers the MW and the uneven distribution of hydrogen inside the reactor. Because hexene comonomer was added in synthesizing the HDPE-B, GPC-IR, with its ability to measure the comonomer content, was utilized to determine MW, MWD, and hexene contents. In HDPE-B, low MW components account for 62 wt % with a peak MW at 24000 and 2 wt % hexene, whereas the high MW components are about 38 wt % with a peak MW at 373000 and 3 wt % hexene. It should be mentioned that the miscibility of polyolefins depends on its coil dimension (when no polar interactions are in play). HDPE with 2−3% hexane has its coil dimension being about 98−99% of the pure HDPE without any hexane. Hence, it is expected that the impact of 2−3% hexane comonomer on miscibility and interfacial tension of the HDPE-B with other polymers would be only around a few percent which would result in a negligible influence on morphology in a polymer blend. 3.2. Rheology and Interfacial Tension. The complex viscosity and storage modulus of all neat polymers as a function of angular frequency are plotted in Figure 2. All PPE and HDPEs examined have relatively high viscosity, except for HDPE-L. The Cox−Merz rule where η* (complex viscosity) = η (shear viscosity) at ω (angular frequency) = γ̇ (shear rate) is applied to estimate the shear viscosity.29 The relation is empirical but has been shown to be widely valid in polymer melt systems. A variety of PEs with different structures have been found to closely follow the Cox−Merz rule.30−33 For PPE, no study has examined the validation of the Cox−Merz rule, possibly due to the very high melt viscosity which makes it difficult to obtain the shear viscosity data. However, a PPE/ PP blend has been reported to obey this relation.34 Therefore, we assume that the Cox−Merz rule is applicable for the PPE and HDPE polymers used in this study. The maximum shear rate in the miniature twin-screw extruder during processing is estimated to be around 200 s−1.35 As it is difficult to accurately predict the shear rate in the extruder, in this study, we consider the rheological properties of the polymers under the shear rate range 50−200 s−1, which lies in the typical shear rate range in extrusion polymer processing. In this region, PPE has the highest viscosity followed by HDPE-H and then HDPE-B (Figure 2 and Table 2). The HDPE-L viscosity is significantly lower than all the others. The zero-shear viscosity was estimated by applying the Cross model, and very high values were obtained for PPE, HDPE-H, and HDPE-B as shown in Table 2. The storage modulus, a measure of the melt elasticity, follows a similar pattern as the viscosity with reference to the material type. It is very difficult to measure the interfacial tension between HDPE and PPE directly due to the very high PPE viscosity. Therefore, the interfacial tension of HDPE/PPE is estimated using the following harmonic mean equation:36

3. RESULTS AND DISCUSSION 3.1. MWs and MWDs of HDPEs. GPC results for HDPEH, HDPE-L, and HDPE-B are shown in Figure 1 and Table 1.

Figure 1. MW and MWD of (a) HDPE-H, (b) HDPE-L, and (c) HDPE-B.

Despite the use of a single-site organometallic catalyst to synthesize HDPE-H and HDPE-L, both HDPEs have PDIs (or Mw/Mn) much greater than the expected value of 2 due to the multisite characteristics introduced by supporting single-site catalysts on a silica support. The much broader MWD found in HDPE-L is the result of high hydrogen amounts introduced to C

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Figure 2. Complex viscosity and storage modulus curves of PPE, HDPE-H, HDPE-B, and HDPE-L at 260 °C. The dashed lines indicate the angular frequency range from 50 to 200 rad/s.

HDPE and PPE by the harmonic mean equation may be quite reasonable. 3.3. Morphology and Continuity. The morphology development of the HDPE-B/PPE blends with increasing PPE concentration from 10 to 90 wt % can be found in Figure 3. It should be mentioned that both SEM and AFM analyses were performed on the samples, and the results correspond well. However, because of the large phase size variation at different compositions in different blends (also see Figure 5), we have selectively used SEM images at HDPE/PPE 50/50 and AFM images at other compositions to clearly show the morphology. PPE remains dispersed in blends with up to 30 wt % PPE. In the 50/50 HDPE-B/PPE blend, connected PPE domains are evident (see Figure 3d). A cocontinuous morphology is observed in blends with 70 and 80 wt % PPE and potentially extended to 90 wt % PPE (Figure 3e−g). It should be noted that when cryo-microtoming across the extrudate filament, phase morphologies reside on the plane perpendicular to the extrusion flow direction limiting a full three-dimensional cocontinuity determination.17 For this reason, the extraction gravimetry method was applied to quantify the phase continuity which will be discussed in the following section. As shown in Figure 3, when HDPE-B is the minor phase, highly percolated HDPE-B phases can be found even at very low HDPE concentrations. Continuous HDPE-B phase is clearly evidenced at 20 wt % HDPE (Figure 3f). At 10 wt % HDPE-B in the blend, a high HDPE continuity is likely considering the observed connected HDPE structures in Figure 3g. With HDPE being highly incompatible with PPE, a result of their high interfacial tension, it is expected to find large and micrometer-size dispersed PPE phase dispersions in blends with HDPE as the major phase. As indicated in Figure 3a−c, large PPE dispersions can be seen in these blends. However, in blends where HDPE is the minority phase, HDPE-B appears to form highly percolated phase structures with fine submicrometer phase sizes (see Figure 3 and Table 4). As shown in Figure 4, an HDPE-B/PPE 10/90 blend filament did not disintegrate after 1 day in chloroform. It floats while maintaining its structural integrity without clouding the solvent, demonstrating fully self-supporting HDPE structures after PPE removal. By cryo-fracturing this extracted filament, a continuous HDPE phase morphology is indicated in this extracted filament by the SEM micrograph (see Figure 4b). A SEM micrograph of a cryo-fractured HDPE-B/PPE 5/95 blend with PPE extracted is shown in Figure 4c where the early stage of HDPE continuity development with HDPE-B forming fiber phases is suggested.

Table 2. Viscosities (Cox−Merz Rule Applied) and Storage Modulus of the Raw Polymers polymers

zero-shear viscosity (Pa·s)

viscosity at 50− 200 s−1 (Pa·s)

storage modulus at 50−200 s−1 (Pa)

PPE HDPE-H HDPE-B HDPE-L

172980 151510 103380 707

8250−3250 4860−1770 2890−1101 321−221

395000−603000 220000−270000 146000−203000 6030−21000

d d γ pγ p zy ji γ γ γ12 = γ1 + γ2 − 4jjjj d 1 2 d + p 1 2 p zzzz jγ + γ γ1 + γ2 z (2) 2 k 1 { where γ12 is the interfacial tension between polymer 1 and 2; γdi and γpi are the dispersive and polar components, respectively, of the surface tension γi for polymer i. As shown in Table 3, the

Table 3. Surface Tension Parameters of HDPE and PPE (260 °C) polymer a

HDPE PPEb

γ (mN/m)

polarity

γd (mN/m)

γp (mN/m)

22.0 25.8

0 0.2

22.0 20.6

0 5.2

Data extrapolated based on the data tabulated in ref 36: γ(180 °C) = 26.5 mN/m; −(dγ/dT) = 0.057 mN/(m °C). bEstimated from ref 37 using Macleod’s relationship.

a

surface tension of HDPE at 260 °C is extrapolated from the data in the literature.36 The measurement of the neat PPE surface tension at high temperature is very difficult due to the very high viscosity. Therefore, the PPE surface tension at 260 °C was calculated with Macleod’s relationship in a previous study.37 Note that the polarity is independent of temperature. Using the parameters listed in Table 3, the interfacial tension between HDPE and PPE is estimated to be 5.2 mN/m, suggesting high incompatibility. This value is similar to a previously reported value of 5.6 mN/m for HDPE/PS (at 195 °C).19 The close values are expected since PS and PPE are miscible, and a previous study shows that PPE and PS have similar surface tensions.37 Moreover, the interfacial tension between polypropylene (PP) and a mixture of polystyrene/ PPE with various PPE concentrations lies in the range 4.5−6.5 mN/m (γPP/PS ≈ γPP/(PS/PPE mixture) ≈ γPP/PPE).37 Because of the chemical similarity, PP and PE typically show comparable interfacial tensions against a third polymer. Therefore, the estimated interfacial tension value of 5.2 mN/m between D

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Figure 3. Cryo-microtomed AFM (a−c and e−g) and SEM (d) images showing the morphology plane perpendicular to the extrusion direction for the HDPE-B/PPE blends from 90/10 to 10/90. The light phase is PPE, and the dark is HDPE-B. The inserted micrograph in (g) is an amplified AFM image with monochrome color scheme (the scale bar is 2 μm).

should not create fibers or connected HDPE structures if HDPE shows a dispersed morphology. The morphology development of HDPE-H/PPE and HDPE-L/PPE blends with increasing PPE concentration from 10 to 90 wt % is illustrated in Figure 5. The viscosity of HDPE-H is similar to that of HDPE-B with the former being slightly higher (Figure 2), and their viscosity values are about 5−15 times higher than those of HDPE-L within the shear rate range 50−200 s−1. Finer PPE dispersions are found in the HDPE-H/PPE blends, when PPE is the minority blend phase, as compared to those large PPE dispersions observed in the HDPE-B/PPE and HDPE-L/PPE blends (see Figure 5 and Table 4). In blends where the PPE composition is raised beyond 50 wt %, their cocontinuous PPE phase size differences diminish gradually among these three HDPE/PPE blend systems. Viscosity ratios of the dispersed phase to the matrix phase for the HDPE/PPE blends are listed in Table 4 where a correlation between the viscosity ratio and the dispersion size in blends having PPE dispersions is indicated. The high viscosity ratio of 16−25 between PPE and HDPE-L leads to large PPE dispersions found in the 90/10 HDPE-L/PPE blend as compared to those observed in the HDPE-H/PPE blend.10,13 Surprisingly, despite the similar viscoelastic properties of HDPE-H and HDPE-B, PPE dispersion sizes in the

Table 4. Phase Sizes of PPE in 90/10 HDPE/PPE Blends and HDPE in 10/90 HDPE/PPE Blends with Various Viscosity Ratios HDPE/PPE (90/10) viscosity ratio at shear rates of 50−200 s−1 diameter of the PPE phase (μm) HDPE/PPE (10/90) viscosity ratio at shear rates of 50−200 s−1 diameter of the HDPE phase (μm)

HDPE-H/ PPE

HDPE-B/ PPE

HDPE-L/ PPE

1.7−2.0

2.8−2.9

16−25

0.43 HDPE-H/ PPE

3.1 HDPE-B/ PPE

2.9 HDPE-L/ PPE

0.50−0.59

0.34−0.36

0.06−0.04

0.29

0.29

0.34

Both early continuity onset and fine fiber-like morphologies of HDPE-B seem to imply compatibility between PPE and HDPE-B,19,38 which is contrary to the known incompatibility between PPE and polyolefins (e.g., PE and PP).14,24,25 This unusual and unexpected morphology development of minority HDPE-B in the PPE matrix major phase may be related to their viscoelastic properties and will be discussed in the following section. It should be mentioned that cryo-fracture may induce some defects, but the process was performed with the samples being immersed in liquid N2, and therefore it

Figure 4. (a) An HDPE-B/PPE (10/90) blend sample immersed in chloroform for 1 day to extract PPE. Morphology of HDPE-B after matrix dissolution: (b) HDPE-B/PPE (10/90) and (c) HDPE-B/PPE (5/95). E

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Figure 5. AFM and SEM micrographs on the cross-sectional surface perpendicular to the extrusion direction of the HDPE-H/PPE and HDPE-L/ PPE blends from 10/90 to 90/10. The light phase is PPE, and the dark phase is HDPE. Photos (d) and (k) are SEM images, while the others are AFM images. Inserted AFM micrographs shown in (g) and (n) are magnified with monochrome color scheme (the scale bar is 2 μm).

Figure 6. (a) 10/90 HDPE-H/PPE and HDPE-L/PPE blend samples immersed in chloroform for 1 day to extract PPE. (b, c) Morphology (SEM images) of 10/90 HDPE-H and HDPE-L after matrix dissolution.

extraction. This strongly suggests their cocontinuities. The diameters of these HDPE domains in the 10/90 blends are about the same among the three HDPEs at ∼0.3 μm (Table 4). Here, when PPE is the matrix, the low and different viscosity ratios among the three HDPE/PPE blends seem to have little effect on dispersing HDPEs to fine dispersions and on the HDPE dispersion size. This is different from Newtonian systems where the ease of droplet breakup is heavily dependent on viscosity ratio. However, it should be noted that previously studied Newtonian fluids typically possess low viscosities (for example, 0.05−300 Pa·s in the two classic studies from Grace

HDPE-B/PPE blends are similar to those found in the HDPEL/PPE blend, but not in the HDPE-H/PPE blend (see Figures 3 and 5 and Table 4). This finding suggests that the low MW components in HDPE-B may have populated the HDPE-B/ PPE interface during melt processing. On the other hand, the same cannot be said when HDPE is the minor component in the blends. The 10/90 HDPE-H/PPE and HDPE-L/PPE blends after matrix dissolution are shown in Figure 6, and continuous HDPE structures can also be clearly observed. In those 10/90 HDPE/PPE blends, HDPE-H, HDPE-B, and HDPE-L are all self-supporting, without falling apart, after PPE F

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Macromolecules as well as from Karam and Bellinger12,39). It was shown that with the same viscosity ratio droplets become easier to break up (i.e., droplets break up at lower shear rates) with increasing matrix viscosity. The viscosity of PPE in this study is extremely high with a zero shear viscosity of 172980 Pa·s and an estimated viscosity of 8250−3250 Pa·s during processing. Therefore, the limited influence of viscosity ratio on HPDE morphologies in PPE matrix may be due to a dominant effect of high PPE viscosity. It should be noted that mixing time may also have a significant influence on the polymer blend morphology, but it generally requires only a few minutes to achieve the final morphology.40,41 In this work we used 5 min mixing (in addition to about 2 min loading time) which should be sufficient enough to achieve a final equilibrium morphology according to previous studies.42−44 We have also processed some HDPE/PPE (10/90) blends for a longer time (10 min) and examined the morphologies (see Figure S1 in the Supporting Information), and it can be seen that there is no significant difference in morphology by increasing the processing time. It is important to address the migration rate of the low MW HDPE to the interface of HDPE-B/PPE. Under static conditions, this rate may be slow. A previous study examined the interfacial tension evolution in a polydisperse system of polyisobutylene (PIB)/poly(dimethylsiloxane) (PDMS) using the pendent drop method, and they found that it could take over a day to achieve a constant interfacial tension, indicating a slow migration rate of the low MW components to the interface.45 However, in this study, intensive mixing is involved which can greatly facilitate the migration process. The calculation of the migration time in the present system is very complicated due to the complex flow field during the processing. However, it may be reasonable to consider a similar interfacial migration process in polymer blends as a guideline: the migration of a compatibilizer (or interfacial modifier) to the interface of two immiscible polymers. In previous work, it was also shown that the migration of a copolymer compatibilizer to the PIB/PDMS interface under static conditions is even slower (1−2 orders of magnitude slower) than that of low MW component migration.45 It has been generally recognized that in polymer blend systems only a few minutes of mixing are needed for a compatibilizer to migrate to the interface and effectively reduce the interfacial tension. Therefore, we may assume that the mixing time required for the migration of low MW species to the interface during dynamic mixing can also be low. Also, as mentioned previously, selected HDPE/PPE blends at a longer mixing time were studied, and no noticeable morphology change was found (Figure S1). Therefore, the mixing time used in this work should be long enough for the low MW HDPE species to effectively migrate to the HDPE-B/PPE interface. Most polymer pairs are immiscible since the enthalpy of mixing (ΔHmix) is normally positive (absence of a specific interaction) and entropy (ΔSmix) is small and negligible due to their high molecular weights. Nevertheless, it has been found that MWs and MWDs of polymers can modify the interfacial tension in polymer blends.46,47 Using a mean-field theory to examine MW and MWD effects on interfacial properties, Broseta et al. suggested, through computation, that the low molecular weight species entropically populate at the interface to lower the interfacial tension.46 Experimentally, Nam and Jo found that the interfacial tension of PS/polybutadiene (PBD)

increases with increasing MW and with decreasing PDI (polydispersity index, a measure of MWD) of PS.48 In particular, a lower interfacial tension is reached by using a bimodal PS which was attributed to the migration of the low MW PS species to the PS/PBD interface. Arashiro and Demarquette showed that increasing the PS MW raises the interfacial tension of PE/PS until the MW of PS surpasses the PS entanglement MW of 45000.49 Similar results were reported by Kamal et al.50 and Ellingson et al.51 These observations suggest that once the polymer MW exceeds the entanglement MW, the influence of MW on interfacial tension diminishes and it will become more difficult for low MW species to migrate to the interface. The entanglement MW of polyethylene was determined to be in the range 500−3000 g/mol with different techniques.52 In this study, as shown in Figure 1, only a very small portion of the HDPE-H molecules have a MW less than 10000 g/mol. However, large amounts of HDPE-B and HDPE-L molecules have a MW below 10000 g/mol with some portion of them even below 1000 g/mol. During melt mixing, the migration of low MW HDPEs (most likely below a few thousand) in HDPE-B and HDPE-L to the interface of HDPE/PPE is clearly favorable since it reduces the interfacial tension of the system. It is noted that the polymer migration toward interfaces under a flow field is also widely acknowledged in the literature.53−56 While an entropy effect is the predominant argument, there are others who maintain that the migration of low-viscosity components to the high-shear regime is simply due to the tendency to minimize viscous dissipation in the system.53,57−61 It is suggested that the low MW components in HDPE-B migrate toward the HDPE-B/PPE interface, either entropically or energetically (viscous energy minimization), or both, to create an interfacial rheological environment that is similar to that found in the HDPE-L/PPE blend and hence yielding similar dispersed PPE morphologies. The migration of the low MW HDPE species is favorable due to the reduction of interfacial tension. The migration lowers the viscous forces acting on the dispersed phase which are dominantly exerted by the low MW HDPE components populating the interface. It appears that the low-viscosity effect dominates over the interfacial tension influence on blend morphology since the phase size of the dispersed PPE in HDPE-B/PPE is larger than that in HDPE-H/PPE. The partition of low MW species in HDPE-B at the interface and its influence on morphology development are clearly supported by the fact that HDPE-B/PPE and HDPE-L/PPE with significantly different HDPE viscoelasticity, but high similarity in HDPE MW and MWD in the low MW region (typically 90% or if the phase is self-supporting after removing the other phase. As indicated in Figure 7, a very broad cocontinuity region from about 40 to 90

Figure 8. SEM images of the cryo-fractured samples: (a) HDPE-H/PPE 90/10, (b) HDPE-B/PPE 90/10, and (c) HDPE-L/PPE 90/10. H

DOI: 10.1021/acs.macromol.8b01445 Macromolecules XXXX, XXX, XXX−XXX

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Macromolecules blends should be negligible (it may still have occurred but likely to a very small degree with the processing conditions used in this study). When HDPE is the matrix in all three HDPE/PPE blends, a droplet morphology of PPE dispersions is observed, and a late continuity onset of the PPE phase is consistent with expected behaviors in an incompatible polymer blend. But, when PPE is the major phase, very fine, submicrometer, HDPE phase structures and extremely low HDPE phase continuity onset are observed. This highly asymmetric phase continuity diagrams and phase morphology in HDPE/PPE blends suggest that the viscoelastic properties of the matrix polymer may play a key role in controlling the morphology development since the compatibility between HDPE and PPE does not change as a function of blend concentration. During the dispersion of one liquid in an immiscible matrix liquid, their droplets can be broken up through two mechanisms, either by stepwise droplet breakup or by breakup of extended threads into a series of droplets (capillary instability). The latter case has often been observed in polymer blending.17,71 Although in Newtonian systems a high viscosity ratio (typically >4) can effectively suppress droplet breakup, this threshold has been found to be well extended in polymer blend systems. Many studies have shown that polymer droplets can be deformed and disintegrated even at viscosity ratios of 15−20.13,14,72 However, the underlying mechanism has not been definitively resolved. The presence of elongational flow during melt processing of polymers, which is more effective in deforming and disintegrating the droplets, is considered a major factor. Other factors may include viscoelasticity effects, complex transient shear flows, and viscosity/temperature variation during processing. Additionally, the molecular entanglement phenomenon in polymers may also be important since it can enhance the stress transfer at the interface as compared to Newtonian fluids. In the low viscosity region, Karam and Bellinger showed in Newtonian systems that a low viscosity ratio results in high deformation before breakup.39 Below a critical viscosity ratio value, the dispersed phase will be highly deformed without breakup. When HDPEs are dispersed in PPE, the blend viscosity ratio goes from 0.5 to 0.05 which could lead to highly deformed HDPE dispersions without substantial breakup. Correspondingly, HDPE phase domains may form highly extended fiberlike dispersions. The breakup time (tb) of the dispersed HDPE in a PPE matrix and the PPE dispersions in a HDPE matrix can be estimated using the equation73−75 tb =

ij 1.39γR 0 2 yz zz lnjjj z Ω(x , p)γ jk kT z{

Table 5. Morphologies and Fiber Breakup Time of PPE in 90/10 HDPE/PPE Blends and HDPE in 10/90 HDPE/PPE Blends blend type HDPE/PPE (90/10)

HDPE/PPE (10/90)

PPE morphology breakup time tb (s) HDPE morphology breakup time tb (s) breakup time t′b (s)b

HDPE-H/ PPE

HDPE-B/ PPE

HDPE-L/ PPE

fiber

droplet

droplet

continuous

continuous

continuous

587

484a

89

11

12

3.5

48

a

Estimated with the measured viscosity without considering the low MW species populating at the HDPE/PPE interface. bBreakup time is estimated with a lower PPE viscosity of 1000 Pa·s, which is close to that of HDPE-L.

appears to be long enough to form fibrous HDPE structures. Note that the tb for HDPE-B in a PPE matrix was estimated with the measured viscosity, but the real breakup time would be close to that of the HDPE-L/PPE blend considering the low MW species of HDPE-B populating the HDPE/PPE interface. The PPE fibers in HDPE-H/PPE show a moderate breakup time of 48 s, which is about half of the time of HDPE-L in a PPE matrix. This moderate breakup time may explain the formation of a fibrous PPE morphology in the HDPE-H matrix at low compositions (Figure 8). The breakup times for 90/10 HDPE-B/PPE and HDPE-L/PPE are not calculated since PPE are present as droplets (no fiber diameters are available) which indicate that the breakup time is very fast. If the PPE matrix viscosity were to be lowered to that of HDPE-L, then the HDPE fiber breakup times in all three HDPE/PPE blends would be significantly reduced by 30−50 times to 5−10 s (see t′b in Table 5). This demonstrates the importance of the high PPE matrix viscosity in deforming and stabilizing HDPE fiber-like dispersions. Following Tomotika’s theory which is based on mixing of Newtonian liquids, not non-Newtonian polymer melts, the thread breakup time increases with increasing matrix viscosity and viscosity ratio.75 With the high PPE viscosity, HDPE dispersion morphologies in all three HDPE/PPE blends are similar although their viscosity ratios vary from 0.05 to 0.5. This suggests that PPE matrix viscosity has surpassed a critical value and that viscosity ratio is no longer relevant. If one uses the same argument, this may explain the observation of some fiberlike dispersions of PPE in HDPE-H matrix in 90/10 HDPE-H/ PPE blend as a result of the high HDPE-H matrix viscosity (see Figure 8a). 3.4. Dynamic Mechanical Analysis. The dynamic mechanical properties of HDPE/PPE blends were measured and can be used to define PPE and HDPE phase continuities. In a cocontinuous binary blend, the blend modulus comes from the contributions of both phase moduli, whereas in a dispersed binary blend, the blend modulus is dominated by the matrix phase modulus.76,77 In Figure 9, the storage moduli of the HDPE/PPE blends at 150 °C are plotted against PPE concentration. This 150 °C was chosen since, at this temperature, HDPE is a melt whereas PPE is a solid. The storage modulus is very small (