Bimodal High-Density Polyethylene: Influence of the Stereoregularity

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Bimodal High-Density Polyethylene: Influence of the Stereoregularity of the Copolymer Fraction on the Environmental Stress Crack Resistance Maila N. Cardoso† and Adriano G. Fisch*,‡ †

Innovation and Technology Center, Braskem S.A., III Pólo Petroquímico, Via Oeste lote 5, Passo Raso, Triunfo, RS 95853-000, Brazil ‡ Chemical Engineering Department, Universidade Luterana do Brasil, Avenida Farroupilha 8001, Canoas, RS 92425-900, Brazil ABSTRACT: In the present study, it is demonstrated that the stereoregularity of the copolymer fraction of bimodal high-density polyethylenes (HDPEs) influences the environmental stress crack resistance. Polymer samples were characterized in terms of their semicrystalline microstructure. A certain stereoregularity was identified for the copolymer fraction, and a linear correlation between the stress crack resistance, which was assessed by the full notch creep test, and the fraction of stereoregular chains in the HDPE was established. Conceptual models to understand the action of the stereoregular copolymer chain in the slow crack growth mechanism are proposed. The results found in the present research will serve as guidelines for tailoring the properties of bimodal high-density polyethylene through optimization of the operating parameters of the reactors and catalyst design.

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INTRODUCTION Polyethylenes are used in many applications in which environmental stress cracking is a relevant property of the material. Pipes, bottles, and containers, for instance, are applications of high-density polyethylenes (HDPEs) that have lifetimes that are dependent on the applied load and aggressiveness of the environment.1−3 As suggested by some earlier studies,2,4,5 failure due to environmental stress is controlled by the slow crack growth (SCG) mechanism of fracture. According to SCG theory, failure starts from the formation of cavities (i.e., voids) in the amorphous phase located between crystalline lamella during the first stages of deformation.6 At small strains, the crystalline phase dominates the mechanism through intra- and/or interlamellar slip processes.7−9 At large strains, the mechanism changes and becomes controlled by characteristics of the amorphous phase, which crystallizes as fibrils (i.e., crazing) as a result of the drawing. Indeed, fragmentation of the crystalline lamellae occurs, and the respective polymeric chains could recrystallize again, forming fibrils as well.7,8 Different concepts have been investigated to address the relationship between the stress crack resistance and the semicrystalline microstructure of the polymer. In this sense, the first concept is based on the population of tie molecules, which are polymer chains that exit one crystal region and enter another after a random walk in the amorphous phase. These tie molecules allow the lamellae to be linked to each other and to the amorphous phase as well. Consequently, the macromolecular network formed by the crystalline and amorphous domains is held together by the tie chains, so that the semicrystalline structure is stiffer than it would be in the absence of tie chains. It is worth mentioning that the crystalline © XXXX American Chemical Society

and amorphous domains can also be tied together by trapped entanglements of chains. It has been suggested10 that the presence of tie chains improves the stress crack resistance by retarding crack initiation, probably avoiding the formation of cavities, according to the SCG mechanism. Previous studies11,12 have shown that the concentration of tie molecules is related to the portion of short chain branches incorporated into long chains, that is, into the high-molecular-weight fraction of the polymer. Additionally, the chain length of the comonomer also increases the effects of the tie molecules on environmental stress resistance; for example, 1-hexene presents better performance than 1-butene in a polyethylene copolymer.2,3,7,11 In this sense, the distributions of both the molecular weight and chemical composition play a relevant role in the development of resins with improved stress crack resistance. Another feature of the semicrystalline microstructure of the polymer that is considered to explain stress crack resistance is the mobility of the amorphous domain of the semicrystalline network.9 An amorphous phase that is more mobile allows the formation of larger and stronger craze fibrils connecting the crack surfaces during polymer deformation. According to this concept, rupture of the fibrils and crack propagation are delayed, resulting in the improvement of the stress crack resistance. The interplay between tie molecules and amorphous-phase mobility was also studied13 in an attempt to explain the full notch creep test (FNCT) values of high-density, linear lowReceived: March 8, 2016 Revised: May 10, 2016 Accepted: May 17, 2016

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DOI: 10.1021/acs.iecr.6b00927 Ind. Eng. Chem. Res. XXXX, XXX, XXX−XXX

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column temperatures were 150 °C. The molecular weight distribution was calculated by the universal calibration method using polystyrene standards as a reference. The melting temperatures (Tm) and crystallinity contents (χc) were obtained by differential scanning calorimetry (DSC) using a TA Instruments apparatus (DSC Q1000). Indium was used as a standard for the calibration of the melting temperature. A rate of 10 °C min−1 in the range from −20 to 160 °C was used for both melting and cooling runs. The heating cycle was performed twice, but only the results from the second run are reported. The Fourier-transformed infrared (FTIR) spectroscopy was performed using a Shimadzu spectrophotometer (IR Prestige 21). The spectra were obtained by coadding 32 scans at a resolution of 1 cm−1 under inert (N2) atmosphere. The samples were analyzed in absorbance mode as films (10−20 μm). The range of 3950−4482 cm−1 was used for normalization because of the differences in the film thicknesses. The fraction of the HDPE that was soluble in n-hexane was obtained using a Soxhlet extractor. The extraction was conducted using about 10 g of sample and lasted for 5 h. The polymer solution was dried by careful evaporation of the solvent. 13 C nuclear magnetic resonance (NMR) spectroscopy was conducted on an Agilent 400 spectrometer operating at 100 MHz and using a probe of 5 mm. The polymer solution was prepared by dissolving ca. 50 mg in 0.6 mL of a solution of odichlorobenzene and deuterated tetrachloroethane (75/25 vol %). The deuterated solvent was used to provide the internal lock signal. The chemical shifts were referenced internally to the major backbone methylene carbon resonance, which was taken as 30.00 ppm. The experiments were performed at 120 °C. The pulse length was set up for a 74° flip angle, the acquisition time was 1.5 s, and the delay time was 4.0 s. These parameters allow the spectra to be 90% quantitative for carbon atoms that have relaxation times (T1) lower than 2.0 s.15 The environmental stress crack resistance was evaluated by the full notch creep test (FNCT) following standard method ISO 16770. The polymer density was obtained in accordance with standard method ASTM D-1505.

density, and low-density polyethylenes. However, this approach also failed in fitting the experimental data using a single equation, even though the parameters had phenomenological explanations. Thereby, the existence of a more complicated mechanism controlling the deformation of a semicrystalline structure under mechanical stress in a chemically aggressive environment is evidenced. The ordinary production process of bimodal HDPE uses two, or even three, reactors in series, allowing the molecular weight and chemical composition distributions to be designed in accordance with the desired properties.1 The properties of the resin depend on the polymerization parameters of the reactors, such as temperature, polymerization rate, partial pressure of monomer and comonomer, and hydrogen-tomonomer (H2/C2) and comonomer-to-monomer (C4/C2) ratios.14 The catalyst nature is also a variable to be considered during resin development, mainly when the balance between mechanical properties, stiffness, and processability is the focus. Preferentially, the comonomer is added to the second and/or third reactor to maximize the product properties. Following this industrial process, the obtained resin is a tailor-made blend of homopolymer (HP) and copolymer (CP) of ethylene.1 The addition of comonomer to the second (or third) reactor increases the probability of generating tie molecules in the semicrystalline polymer structure once this reactor produces the higher-molecular-weight polymer of the blend.2 Depending on the product application, the increase of the tie-chain concentration obtained by adding comonomer to the second/ third reactor is limited to a minimum product stiffness.3 In this sense, the improvement in the stress crack resistance achieved by using such a polymerization technology is also limited. Taking into account the aforementioned operating parameters for product development and previous studies dealing with the slow crack growth mechanism, the influence of the stereoregularity of the side branches incorporated into the copolymer fraction of bimodal HDPE on the failure mechanism, or even on the tie-molecule concentration and the stress crack resistance, has not been completely addressed so far. With the aim of contributing to the development of polyolefin technology, the present work deals with the influence of the stereoregularity of the copolymer fraction of bimodal HDPE on the stress crack resistance. In this study, we present a structural characterization of the polymers and its relationship to the environmental stress crack resistance. Conceptual models intended to explain the importance of the stereoregularity of the copolymer fraction to the stress crack resistance are also proposed.

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RESULTS AND DISCUSSION Bimodal HDPE Characterization. As reported in Table 1, the HDPE samples analyzed in this research presented a narrow range of average molecular weights. On the other hand, the samples showed a broad range of molecular weight distributions (Mw/Mn = 7.8−20.5). These characteristics are attributed to the bimodal process and the unique set of

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EXPERIMENTAL SECTION All samples analyzed in this research were high-density polyethylenes obtained from an industrial bimodal polymerization process (two reactors in series) using 1-butene as the comonomer. The catalyst system used in the polymerization process was the same for all resins. The addition of comonomer to the reaction medium occurred solely in the second reactor. The HDPE samples were obtained as powders, that is, without any incorporation of additives. The polymers were characterized in terms of molecular weight distribution by high-temperature size-exclusion chromatography (HT-SEC) using a Polymer Char apparatus (GPCIR). In a typical analysis, ca. 8 mg of HDPE was dissolved in 8 mL of trichlorobenzene (TCB) at 160 °C for 1 h. The samples were injected at a flow rate of 1 mL min−1. The injector and

Table 1. Features of the Samples of High-Density Polyethylene

a

B

sample

Mwa (kg mol−1)

HDPE-1 HDPE-2 HDPE-3 HDPE-4 HDPE-5 HDPE-6 HDPE-7

228 218 251 222 277 173 165

Mw/Mn Mz/Mw 14.8 16.4 7.8 20.5 9 11.9 12.7

4.2 4.6 3.6 4.1 3.8 4.6 4.1

Tmb (°C)

χc (%)

density (kg m−3)

131.8 134.3 129.7 130.1 133.9 134.1 131.8

70.7 85.8 59.6 68.5 78.3 82.8 70.7

955 960 944 952 952 961 952

From HT-SEC. bFrom DSC. DOI: 10.1021/acs.iecr.6b00927 Ind. Eng. Chem. Res. XXXX, XXX, XXX−XXX

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(copolymer), it is rational to evaluate the copolymer fraction deeply, as it is responsible for the creation of a unique polymer microstructure that favors stress crack resistance. Based on this argument, it was speculated that the stereoregularity of the copolymer fraction has some influence on the FNCT results. Figure 2 shows a typical FTIR spectrum of the HDPE samples in which is possible to observe the ordinary peaks.17,18

conditions used in the polymerization reactors, including the reactor temperatures, the split of production between reactors, the hydrogen-to-monomer ratio, and the comonomer-tomonomer ratio. This allowed a broad range of polymer microstructures to be produced, as indicated by its melting points (Tm), crystalline contents (χc), and densities. Based on previous studies,16 the environmental stress crack resistance can be enhanced by adding a copolymer fraction to the resin recipe as a result of the increase in the tie-molecule concentration. In this sense, this property could be related to the crystalline content and/or to the melting point of the resin. The failure times in the full notch creep test (FNCT) of the HDPE resins considered in the present study correlated partially either with the melting point (Figure 1a) or with the

Figure 2. Typical FTIR spectrum of HDPE (sample HDPE-7).

The doublet located at 720 and 730 cm−1 is attributed to the inphase CH2 rocking vibration in the crystalline domain. Additionally, the small shoulder at 780 cm−1 is due to the inphase CH2 rocking vibration of ethyl branches. The small peaks located at 887 and 907 cm−1 are attributed to terminal unsaturated carbons of the polymer chain. The peak at 1304 cm−1 is related to the CH2 wagging vibration. The small peak at 1350 cm−1 is due to the CH deformation, and its neighboring peak at 1366 cm−1 is attributed to the CH3 scissoring vibration. The peak at 1464 cm−1 is attributed to CH2 scissoring plus CH3 wagging vibrations. The spectrum in Figure 2 also shows peaks ranging from 1000 to 1100 cm−1, which is the traditional range for vibrations related to chain conformations for general polymers.18,19 Specifically, the peaks at 1051, 1079, and 1174 cm−1 in Figure 2 are characteristic of the C−C skeletal vibrations of a polymer chain in a helical conformation, which are similar to those of stereoregular polymers such as isotactic poly(1-butene) and even polypropylene, for instance. From this evidence, it is supposed that the stereoregular fraction of the bimodal HDPE is an ethylene copolymer with a high content of branches. Therefore, we attempted to extract such a fraction from the whole polymer using n-hexane as the solvent in an effort to separate and to characterize the supposed copolymer. The contents of the soluble fraction in the HDPE samples are listed in Table 2. All HDPE samples presented from 0.38 to 2.8 wt % of the soluble polymer in their compositions. These values are higher than those found for commercial ethylene−1-butene copolymers of uniform chemical distributions, suggesting that the soluble fractions found in the samples of HDPE are not exclusively formed by the stereoregular copolymer but also contain a portion of a linear polymer of low molecular weight. This was confirmed by FTIR spectroscopy (see below). Some of the soluble fractions were also characterized by their melting temperatures using differential scanning calorimetry (DSC). The results of a typical DSC run are depicted in Figure 3 and show two well-defined melting temperatures, ca. 100 and

Figure 1. Relationships between the failure time in the FNCT and the (a) melting point and (b) crystalline content of the polymers.

crystalline content (Figure 1b). In fact, the fitting of the regression curves and their correlation coefficients (R2) revealed a better correlation of the FNCT with the melting temperature. It is worth mentioning that the melting point is influenced by the quality (thickness) of the lamella, rather than the crystalline content. The lack of correlation of the data in Figure 1 indicates the existence of variables that are not addressed by ordinary polymer characteristics influencing the environmental stress crack resistance at the microstructure level. Considering the earlier discussion, the semicrystalline microstructure of the bimodal HDPE grades needs to be described accurately, mainly in terms of chain branch distribution.12 Assuming that the HDPE resin is a reactor blend of linear and branched chains C

DOI: 10.1021/acs.iecr.6b00927 Ind. Eng. Chem. Res. XXXX, XXX, XXX−XXX

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1097, and 1260 cm−1 are well-defined and can be attributed to the C−C skeletal vibrations of chains in a helical conformation.19−21 These peaks are somewhat different from those at the respective positions in the FTIR spectrum of the whole polymer, which is shown in Figure 2. The shifts in the peak positions from 1051 to 1024 cm−1, from 1079 to 1097 cm−1, and from 1174 to 1260 cm−1 probably indicate differences in the helical conformations of the chains. These results can be attributed to the environments around the stereoregular copolymer chains, whose conformations depend on the amorphous and/or crystalline (lamellar) phases. Indeed, the peak at 802 cm−1 in the spectrum of the soluble fraction (Figure 4) does not have a similar peak in the respective spectrum of the whole polymer (Figure 2). This could be a direct result of the overlay of the doublet located at 720 and 730 cm−1, which is intense in the whole-polymer spectrum. The soluble fraction was also characterized by 13C NMR spectroscopy, as depicted in Figure 5. A typical spectrum of a linear polyethylene is shown in Figure 5a. The spectrum in Figure 5b is typical of polymers presenting ethyl branches. The signals at 11.10, 24.65, 30.98, 34.70, and 37.60 ppm that appear in Figure 5c indicate the presence of the regular insertion of ethyl branches in sample HDPE-7. Indeed, the presence of blocks of poly(1-butene) could be observed by the signals at 11.10 and 37.60 ppm. These results provide corroboration for the assumption of the presence of a stereoregular copolymer in the composition of the HDPE samples, which was based on the FTIR spectra. Table 3 lists the carbon assignments for the chemical shifts that appear in the 13C NMR spectra of the samples. Based on the characterization results, the copolymer fraction of the bimodal resins presents some content of stereoregular chains. Indeed, the 13C NMR spectra evidence the existence of copolymer chains formed even by blocks of poly(1-butene). Surely, the presence of stereoregularity in the resins is due to the catalyst/cocatalyst system and, to a certain extent, to the operating conditions of the reactors. However, the regularity of the chain structure identified in this study deserves to be studied in greater depth, mainly by employing a better procedure for isolating the stereoregular fraction. Even though the polymer microstructure has not been completely elucidated, by following a pragmatic approach, the relationship between the stereoregular copolymer chain and the environmental stress crack resistance can still be evaluated in the present study. Relationship between Stress Crack Resistance and Stereoregularity. It is important to mention that poly(1butene) is a polymer with outstanding long-term properties, such as resistances to environmental stress cracking and to creep.20 Considering this fact and the fact that the copolymer fraction presented some degree of stereoregularity, we expected to observe some relationship between the full notch creep test (FNCT) and the content of the stereoregular copolymer in the bimodal HDPE resins. In this sense, the area of the FTIR peaks in the range of 1010−1100 cm−1 was employed in an attempt to quantify the stereoregular copolymer portion of the bimodal HDPE samples. The range under consideration was limited to 1010 cm−1, excluding the peak at ca. 802 cm−1, even though it is related to the helical conformation of the polymer chain, because some terminal unsaturated carbons absorb in the range of 887−907 cm−1, which could interfere in the relationship between the FTIR area and the FNCT failure time. As the FTIR analyses were accomplished in film mode, it was necessary to eliminate the interference of the film thickness.

Table 2. Fraction of the HDPE Samples That Is Soluble in nHexane

a

sample

soluble content (wt %)

melting temperatures (°C)

HDPE-1 HDPE-2 HDPE-3 HDPE-4 HDPE-5 HDPE-6 HDPE-7

1.10 1.48 2.70 1.27 0.38 0.60 2.80

99.3/105.7 97.6/103.8 NAa 97.8/104.3 NAa NAa NAa

NA, not accomplished.

Figure 3. Typical differential scanning calorimetry (DSC) curve of the soluble fraction from sample HDPE-2.

105 °C (see Table 2 for details). These melting temperatures are characteristic of ethylene−1-butene copolymers with high contents of comonomer20 and/or stereoregular comonomer insertion, and they are also lower than those of the overall polymer (ca. 130 °C; see Table 1 in the Experimental Section). A shoulder also appears around 60 °C that is attributed to linear polyethylene with a short chain length. A typical vibrational spectrum of the soluble fractions is shown in Figure 4. The doublets at 1470/1462 cm−1 and 730/ 719 cm−1 are due to the effects of crystal splitting.17,18 These results provide evidence that a lamellar phase exists in the soluble fraction that probably arises from a linear polymer chain of short length (low molecular weight). The peaks at 802, 1024,

Figure 4. Typical FTIR spectrum of the HDPE soluble fraction (sample HDPE-7). D

DOI: 10.1021/acs.iecr.6b00927 Ind. Eng. Chem. Res. XXXX, XXX, XXX−XXX

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Figure 5. 13C NMR spectrum of the polymer soluble fraction from (a) linear HDPE (low molecular weight), (b) HDPE presenting no regular insertion of comonomer, and (c) HDPE presenting some portion of stereoregular branches (sample HDPE-7).

mentioning that few of the copolymer chains present stereoregularity, which indicates the influence of these chains on the semicrystalline microstructure under mechanical stress. Conceptual Models. Despite the evidence, it is difficult to rationalize the influence of the stereoregular copolymer fraction on the failure mechanism considering only the polymer characterization techniques used in the present study. One speculative attempt to explain the results is to consider that the stereoregular copolymer chain is excluded from the lamellae and embedded in the amorphous phase. From this consideration, the mobility of the linear chain in forming folds during crystallization would decrease, and consequently, the probability of forming tie chains would increase. In this model, the stereoregular copolymer chain is not participating as the tie chain or in the lamellae. Indeed, the molecular weight of the

Consequently, the FTIR area of the peaks in the range of 3950−4482 cm−1, which are attributed17 to the combined absorption of methyl and methylene groups, was used to normalize the FTIR area in the range of 1010−1100 cm−1. Figure 6 depicts the relationship between the HDPE stereoregularity measured by FTIR spectroscopy and the respective FNCT failure times. As noticed, the variables present a linear dependence, which is supported by the good fit of the linear regression curve to the experimental data and by the respective correlation coefficient (R2). In fact, the FNCT data were found to correlate better with the stereoregularity than with the melting point or crystalline content of the polymers (compare Figure 6 with Figure 1). To some extent, this confirms the interpretation of the polymer characterization obtained by 13C NMR and FTIR spectroscopies. It is worth E

DOI: 10.1021/acs.iecr.6b00927 Ind. Eng. Chem. Res. XXXX, XXX, XXX−XXX

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Industrial & Engineering Chemistry Research Table 3. 13C NMR Chemical Shift Assignments for Ethylene−1-Butene Copolymers22,23 chemical shift (ppm)

carbon assignment

sequence assignment

11.10 11.25 14.10 22.90 24.65 26.85 27.40 30.00 30.50 30.98 32.20 34.10 34.70 37.60 39.80

1B2 1B2 1s 2s ββ 2B2 βδ+ δ+δ+ γδ+ γγ 3s α δ+ Aγ methine methine

EBB + BBE EBE − − BEB EBE BEE + EEB EEE BEE + EEB BEEB − EBEE + EEBE EBEB + BEBE EBB + BBE EBE

Figure 6. Relationship between the failure time in the FNCT and the normalized area for the range 1010−1100 cm−1.

stereoregular chain would not be important in this approach. A similar model was proposed to explain the crystallization behavior and the crystalline microstructure of syndiotactic polypropylene and syndiotactic poly(1-butene) blends.24 It is worth mentioning that nonstereoregular copolymer fractions are also excluded from the crystalline domain, but such chains are not crystalline, so they do not impose restrictions on the mobility of the amorphous chains during crystallization. Figure 7a illustrates this approach. Evidence supporting this conceptual model is the relationship, albeit rough, between the crystalline content and the stereoregular copolymer fraction (FTIR range of 1010−1100 cm−1), which is shown in Figure 8. As a characteristic of the proposed model, a stereoregular copolymer chain hampers other linear chains toward the formation of the lamellae during crystallization. From another point of view, the presence of stereoregular copolymer chains embedded in the amorphous domain could make this phase stiffer. Following this assumption, the strain of the amorphous domain due to applied stress would be more difficult, and thus, the creation of voids and the crack initiation would be hampered as well. In addition, the diffusivity of a chemical agent, which accelerates the cracking phenomenon in environmental stress cracking, through the polymer matrix is reduced, prolonging the failure time. The permeability of water

Figure 7. Illustrations of the stereoregular copolymer chain (red) in the microcrystalline structure of the polymer, in which (a) the stereoregular copolymer chain is not participating as a tie chain (blue) and (b) the stereoregular copolymer chain (block copolymer) is acting as a tie chain.

through HDPE resins was studied10 to explain the respective values of FNCT failure times. In that study, an empirical equation was formulated based on diffusivity and the tieF

DOI: 10.1021/acs.iecr.6b00927 Ind. Eng. Chem. Res. XXXX, XXX, XXX−XXX

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parameters of the reactors. Moreover, the results are also important for improving catalyst design and for tailoring the properties of bimodal HDPEs.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: adriano.fi[email protected]. Tel.: +55 51 3477 4000. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS The author acknowledge Dr. Antonio M. Netto (Braskem S.A., Brazil) for the 13C NMR measurements. Braskem S.A. is acknowledged for financial support (Project ZPIB06640).

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Figure 8. Relationship between the crystalline content and the normalized area for the range of 1010−1100 cm−1.

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molecule concentration that allowed the environmental stress crack resistance to be predicted. A different model could be considered if stereoregular copolymer chains are blocks in a long chain of polyethylene, as suggested by the 13C NMR results. Assuming that the block chains were as long as necessary to serve as tie chains, the stereoregular blocks would hold the lamellae together. According to this approach, the length of the stereoregular copolymer chains would be important. Indeed, tie molecules could be formed not exclusively by linear chains but by linear and copolymer chains, thereby increasing their probability of formation. Figure 7b illustrates this model. In fact, a deep characterization of the resins in terms of stereoregularity is necessary to determine how the stereoregular copolymer fraction influences the environmental stress crack resistance. Even though the stereoregularity of the copolymer fraction acts to enhance the stress crack resistance, the participation of the nature of the catalyst and its kinetic behavior in the resulting polymer microstructure is evident. The stereospecificity of the catalyst is an additional parameter on which research efforts should be focused. Consequently, it is expected that copolymer fractions produced using a stereospecific catalyst (e.g., metallocenes) could perform better in terms of environmental stress crack resistance.

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CONCLUSIONS The influence of the stereoregularity of the copolymer fraction of bimodal HDPE grades on the environmental stress crack resistance was studied in an industrial-scale test, that is, on samples generated using a series of reactors. The results show that the ordinary parameters, such as resin density and melting point, do not explain the performance of the resin in terms of the FNCT. On the other hand, a relationship was established between the content of the stereoregular copolymer chains in the bimodal HDPE formulation and the FNCT results. This pragmatic study led to speculations about the action of stereoregular chains. Conceptual models were proposed to explain the enhancement of the FNCT failure time. However, it is necessary to characterize the resins in depth, obtaining information about both the amorphous and crystalline domains of the semicrystalline polymer microstructure, to allow for model discrimination. The results found in the present research will serve as guidelines for product design and optimization of the operating G

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DOI: 10.1021/acs.iecr.6b00927 Ind. Eng. Chem. Res. XXXX, XXX, XXX−XXX