UHMW Ziegler–Natta polyethylene: Synthesis, crystallization, and melt behavior

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UHMW Ziegler–Natta polyethylene: Synthesis, crystallization, and melt behavior Muhammad Atiqullah a,∗, Sagir Adamu b, Abdul-Hamid M. Emwas c a

Center for Refining and Petrochemicals, Research Institute, King Fahd University of Petroleum & Minerals, Dhahran 31261, Saudi Arabia Department of Chemical Engineering, King Fahd University of Petroleum & Minerals, Dhahran 31261, Saudi Arabia c NMR Core Laboratory, King Abdullah University of Science & Technology, Thuwal 23955-6900, Saudi Arabia b

a r t i c l e

i n f o

Article history: Received 17 January 2017 Revised 6 April 2017 Accepted 7 April 2017 Available online xxx Keywords: Ziegler Natta catalyst UHMW PE Lamellar thickness distribution Melt behavior Nonisothermal crystallization kinetics Flory’s equilibrium theory

a b s t r a c t The fabrication of normal and UHMW PE end-products involves melting and crystallization of the polymer. Therefore, the melt behavior and crystallization of as-synthesized UHMW PE, and NMW PE and E-1-hexene copolymer have been studied using a new nonisothermal crystallization model, Flory’s equilibrium theory and ethylene sequence length distribution concept (SLD), Gibbs–Thompson equation, and DSC experiments. By using this approach, the effects of MW, 1-hexene incorporation, ethylene SLD, the level of undercooling θ , and crystal surface free energy D on crystallite stability, relative crystallinity α , instantaneous crystallinity χ , the crystallization kinetic triplet, crystallization entropy, and lamellar thickness distribution (LTD) have been evaluated. Consequently, this study reports insightful new results, interpretations, and explanations regarding the melting and crystallization of the aforementioned polymers. The UHMW PE results significantly differ from the NMW PE and E-1-hexene copolymer ones. Ethylene sequences shorter than the so called minimum crystallizable ethylene sequence length, irrespective of E-1hexene copolymer MW, can also crystallize. Additionally, the polymer preparation shows that the catalyst coordination environment and symmetry, as well as achiral ethylene versus prochiral α -olefin steric encumbrance and competitive diffusion affect the synthesis of UHMW PE, particularly the corresponding UHMW copolymers. © 2017 Taiwan Institute of Chemical Engineers. Published by Elsevier B.V. All rights reserved.

1. Introduction Ultrahigh molecular weight polyethylenes (UHMW PEs) are polymers that have MW greater than a million (106 ) g mol−1 . On the other hand, the MWs of normal (standard) molecular weight polyethylenes (NMW PEs) vary from 104 −105 g mol−1 . UHMW PEs are resistant to abrasion wear, chemical attack, corrosion, and moisture absorption as well as to cyclic fatigue and radiation. They have low coefficient of friction (working as a self-lubricating material) and high impact strength. They retain key physical properties even at very low temperature such as −30 °C. These special properties allow UHMW PEs to make end-products that can be used in several high performance applications. MipelonTM, Polystone®, TIVAR®, UTEC, etc. list selected examples of commercial UHMW PEs [1]. The fabrication of UHMW PE end-products involves melting, shaping, and nonisothermal cooling during a typical processing cycle. The cooling mode determines the product crystallinity which

∗

Corresponding author. E-mail address: [email protected] (M. Atiqullah).

influences the physical, mechanical, and chemical properties. Therefore, we review the relevant literature in this area. BarralesRienda and Fatou [2] noted that at high undercoolings, MW > 105 does not appreciably affect isothermal crystallization rate. However, an UHMW PE and a NMW PE show %crystallinity χ c of 35 and 85, respectively. Fatou et al. [3] found that MW significantly affects PE crystallization regimes. For very high MW PE, nucleation rate much exceeds the growth rate. Zhu et al. [4] studied UHMW PE melting and crystallization behaviors. The peak crystallization and melting temperatures Tpc and Tpm of UHMW PE show to be higher than those of a normal high density polyethylene (HDPE). However, the heats of crystallization Hc and fusion Hf vary in a reversed fashion. Parasni and Ramani concluded that UHMW PE crystallization is cooling rate-dependent [5]. The nucleation versus growth rate relation matches that reported by Fatou et al. [3]. The melt onset and peak temperatures decrease linearly with cooling rate. Now, we list the specific limitations of the above reports: i. The melt and crystallization behaviors were not modeled from a microscopic viewpoint using the corresponding complete melting and cooling differential scanning calorimetric (DSC) data.

http://dx.doi.org/10.1016/j.jtice.2017.04.011 1876-1070/© 2017 Taiwan Institute of Chemical Engineers. Published by Elsevier B.V. All rights reserved.

Please cite this article as: M. Atiqullah et al., UHMW Ziegler–Natta polyethylene: Synthesis, crystallization, and melt behavior, Journal of the Taiwan Institute of Chemical Engineers (2017), http://dx.doi.org/10.1016/j.jtice.2017.04.011

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ii. The lamellar thickness distribution (LTD) was not evaluated. Particularly, the crystallization kinetics was not modeled by simultaneously considering activation energy, crystal growth, and frequency factor of the process involved. Note that frequency factor relates thermodynamics (entropy) to kinetics [6]. Consequently, the effects of UHMW PE, having extremely long chain backbone, on distributive melt and crystallization characteristics are not known. iii. The dynamic crystallization and melting behavior of NMW PE and UHMW PE were not studied from Flory’s equilibrium theory perspective [7–9], particularly considering temperaturedependent level of undercooling θ , crystal surface free energy D, and critical stable crystallite sequence number n∗ . iv. The overall UHMW PE thermal behaviors were not investigated considering the synthesis history, particularly the catalyst that designs the PE backbone and microstructure. A probable reason, in this context, may be the challenge facing the synthesis of UHMW PE. This requires a catalyst having exceedingly high chain propagation rate and very low chain termination rate. We have recently shown that investigating PE thermal behaviors from the catalyst structural viewpoint can add new insight into this subject [10–12]. As summarized above, we overall note that melting and crystallization of UHMW PEs, unlike those of normal MW PEs, have not been comprehensively investigated. This prompted us to undertake the present study. We design this work as follows: i. Evaluate the potential of TiCl3 •1/3AlCl3 Ziegler–Natta catalyst to synthesize UHMW PE, poly(1-hexene), and ethylene-1hexene copolymer, particularly at near ambient temperature; and look into the polymerization mechanism. ii. Model the melting behavior and crystallization kinetics of the above UHMW PE and compare them with those of NMW PE using nonisothermal DSC data. iii. Assess the effects of temperature-dependent level of undercooling, crystal surface free energy, and critical stable crystallite sequence number on crystallization. We chose TiCl3 •1/3AlCl3 because it is much cheaper than metallocene and post-metallocene precatalysts and commercially available. It easily dissolves in toluene. From the structural viewpoint, it is a mixed crystal and has interestingly dual Lewis acid sites (Ti3+ and Al3+ ). We shall pursue the proposed modeling work using, as appropriate, Flory equation [7–9,13], Gibbs–Thompson (GT) equation [13–17], and a new nonisothermal Avrami–Arofeev crystallization model that we published in 2013 [18]. We trust that this study will add new insight to the melt behavior and crystallization of UHMW and NMW PEs (including copolymer), and broaden our comprehension of this important subject. 2. Materials and methods 2.1. Materials Silica PQ 3030, having surface area of 322 m2 /g, an average pore volume of 3.00 cm3 /g, and a pore size of 374 Ǻ, was used as the catalyst support. (n BuCp)2 ZrCl2 and MAO (30 wt% in toluene) were bought from Chemtura, Germany. TiCl3 •1/3AlCl3 was purchased from Mcat GmbH, Germany. Analytical grade toluene, n-hexane (both 99.999% pure), molecular sieves, 0.05 w/v% 2,6-di-tert-butyl-4-methyl phenol (BHT), 1,2,4 trichlorobenzene (TCB) (analytical grade and deuterated), triisobutylaluminum (TIBA), diethylaluminum chloride (DEAL)— all were purchased from Sigma-Aldrich. Ethylene (99.999% pure) was procured from a local vendor; and oxygen trap

(OT-4-SS) and moisture absorber (500CC 316-SS), from Agilent and Parker, respectively. 2.2. Synthesis of the supported catalyst The supported catalyst— silica/MAO/(n BuCp)2 ZrCl2 — was synthesized as follows by doing all the manipulations under argon, using standard Schlenk technique. Toluene was dried using 4A activated molecular sieve, and it was used as appropriate. The required amount of silica was dehydroxylated at 250 °C for 4 h using a Thermocraft furnace equipped with a vertical quartz glass tube, a digital temperature indicator and a controller, a gas flow meter, and a vacuum pump. The silica was continuously fluidized during dehydroxylation using nitrogen. Upon dehydroxylation, it was stored in an inert glove box. The required amounts of dehydroxylated silica and the dried toluene were slurried in a specially designed Schlenk flask. Next, a calculated volume of MAO was added drop by drop to a given amount of dehydroxylated silica-toluene slurry under argon at constant stirring and room temperature. Then this was heated for several hours for complete impregnation. Following this, the mixture was cooled to room temperature, and a calculated amount of (n BuCp)2 ZrCl2 was added to it and heated at 80 °C for 1 h This final catalyst was dried under vacuum to free-flowing particles and saved in a glove box. The following compositions— Si (wt%), Al (wt%), Zr (wt%), and Al: Zr molar ratio— were calculated from mass balance. These are 72.20, 26.18, 1.62, and 54.80, respectively. 2.3. Acidity and total acid strength of TiCl3 •1/3AlCl3 The acidity and total acid strength of TiCl3 •1/3AlCl3 were measured using ammonia temperature programmed desorption (NH3 -TPD) technique. The NH3 -TPD experiment was conducted using a Micromeritics AutoChem-II 2029 Analyzer. About 0.03 g of solid TiCl3 •1/3AlCl3 sample (a dual Lewis acid compound) was placed in a U-shaped quartz sample holder, and degassed for 2 h at 500 °C, using argon. The sample was then cooled to 120 °C and saturated for 1 h with NH3 (a monodentate Lewis base ligand) by flowing a mixture of 5% NH3 and 95% He (helium) at 50 ml min−1 . During this process, NH3 adsorbs on TiCl3 •1/3AlCl3 via Lewis acidLewis base coordinate bonding. After that NH3 flow was stopped; He was purged at the above flow rate for 1 h at 120 °C to remove NH3 (if any) trapped in the sample. Then the temperature was increased to 750 °C using a heating rate of 10 °C min−1 to desorb NH3 . The concentration of the physically and chemically desorbed NH3 was recorded using a TCD (thermal conductivity detector). The NH3 -TPD profile of the TiCl3 •1/3AlCl3 sample shows two desorption peaks— one at 163 °C (low temperature) and the other at 540 °C (high temperature). The 163 °C and 540 °C peaks are due to the weak and strong coordination acid sites, respectively [19]. See Fig. 1. The corresponding acidities were calculated by integrating the area under each deconvoluted peak, using the AutoChem-II 2029 Analyzer software. The total concentration of the Ti3+ transition metal acid site turned out to be [1.0 mmol(g catalyst)−1 ], which is much higher than that of the Al3+ post-transition metal acid site [0.1 mmol(g catalyst)−1 ]. This finding can be attributed to the fact that the level of coordination of NH3 with Ti3+ is much greater than that with Al3+ . Ti3+ (1s2 2s2 2p6 3s2 3p6 3d1 ), because of incomplete d orbital, appreciably dominates over Al3+ (1s2 2s2 2p6 ) in this NH3 coordination reaction [20]. The ratio of ionic charge 3 3 density of Ti3+ to that of Al3+ equals to 0.067 : 0.053 = 1 : 1.264 where 0.067 nm and 0.053 nm are the ionic radii of Ti3+ and Al3+ , respectively. Therefore, the effect of ionic charge density, in this particular context, on coordination of NH3 with Ti3+ and Al3+ is less significant [20].

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3

dissolved in 10 ml dried toluene. This catalyst solution and 1 ml of 1.0 M diethylaluminum (DEAL) cocatalyst were introduced into the reactor. Next ethylene fed the reactor at 7 bar(g) pressure, and the impeller speed was raised to 700 rpm. The polymerization run was continued for 5 h Following this, ethylene was vented and acidic methanol was added to deactivate the catalyst. The synthesized polyethylene was washed three times with 20% water–methanol solution, and dried overnight under the fume hood. The homopolymerization of 1-hexene and copolymerization of ethylene with 1-hexene, 1-octene, 1-decene, and 1-dodecene were essentially conducted as summarized above. In each case, 182 ml dried toluene was used. However, the amounts of 1-hexene, 1octene, 1-decene, and 1-dodecene, applied in the copolymerization trials, were 18, 24, 29, and 35 ml, respectively. The 1-hexene homopolymerization also used 18 ml monomer; no ethylene was fed. Fig. 1. NH3 -TPD-determined dual Lewis acid sites (Ti3+ and Al3+ ) of TiCl3 •1/3AlCl3 .

2.5. Molecular weights and polydispersity indices 2.4. Polymerization trials 2.4.1. Metallocene-mediated ethylene homo- and copolymerization with 1-hexene Ethylene and 1-hexene were homo- and copolymerized using a computer-interfaced, AP Miniplant laboratory-scale reactor set up. It consists of a fixed top head and a one-liter jacketed Büchi glass autoclave. The glass reactor was baked for 2 h at 120 °C. Then, it was purged with nitrogen four times at the same temperature. The reactor was cooled from 120 °C to 40 °C. About 200 ml of dried n-hexane was transferred to the reactor. Then, 1.0 ml of 1.0 M TIBA was added to scavenge the impurities that may poison the catalyst, and the mixture was stirred for 10 min. n-hexane was dried by contacting it with activated 4A molecular sieves at room temperature over night that decreased the moisture level to less than 10 ppm. The molecular sieve was activated at 230 °C. At this stage, for the copolymerization, 15 ml 1-hexene was added. The resulting mixture was stirred at 50 rpm for 10 min. For the homopolymerization, no 1-hexene was used. The polymerization trial, in either case, was conducted as follows. An estimated amount of the supported catalyst was slurried in 50 ml of n-hexane. The whole volume was siphoned into the reactor under mild argon flow. No MAO cocatalyst was separately added to the reactor. Ethylene was passed through oxygen- and moisture-removing columns. Then it fed the reactor at a constant flow rate of 50 0 0 Nml/min to maintain approximately the same level of macromixing. The polymerization temperature and stirrer speed were set at 50 °C and 700 rpm, respectively. An anchor stirrer (impeller) was used. The trial was continued for 1 h The polymerization was terminated by closing the ethylene flow and venting the post-polymerization ethylene (in the reactor) to the atmosphere. Then, the data acquisition was stopped; the stirrer speed was reduced to about 100 rpm; and the reactor was gradually cooled to room temperature. After conducting the polymerization trial, the reactor was opened; the resulting polyethylene was dried under ambient condition in a hood, and the dried polymer was weighed to determine the yield. The polymer yield was subsequently used to determine the corresponding catalyst productivity which has been reported in Table 1. 2.4.2. Ziegler–Natta catalyst-mediated ethylene and 1-hexene homopolymerizations and copolymerization with selected α -olefins The glass reactor vessel was baked as before. 200 ml dried toluene and 1 ml TIBA were successively added to it, and stirred at 450 rpm for 10 min. Then the resulting mixture was saturated with ethylene at 2 bar(g) pressure. After this, the ethylene introduced was released. About 50 mg (0.25 mmol) TiCl3 •1/3AlCl3 was

The synthesized polyethylenes were characterized in terms of molecular properties [weight average molecular weight (MW ) and polydispersity index (PDI)] using a Polymer Laboratory gel permeation chromatography (GPC) instrument. The column temperature was set at 135 °C. Polyethylene sample (about 1.0 mg), taken in a 1 ml vial, was dissolved in 1.0 ml butylated hydroxy toluene BHT-stabilized 1,2,4 trichlorobenzene (TCB) as follows. The vial was shaken in the warming compartment of the GPC instrument at 135 °C for about 5 h to completely dissolve the sample. Before injecting the samples, the differential refractive index (DRI) detector was purged for 4 h using TCB (1 ml min−1 ) to obtain stable baseline. Also, the inlet pressure (IP) and the differential pressure (DP) outputs were purged for 1 h The above flow rate of TCB was used, and each sample was analyzed for 35 min. The instrument was calibrated using nine polystyrene standards whose peak molecular weights ranged from 1530 to 15 million. The polystyrene calibration curve was converted into the corresponding polyethylene calibration curve using the Mark–Houwink constants of both polymers [21]. Table 1 reports the measured average molecular weights and the polydispersity indices. 2.6. Thermal properties and density The thermal properties of the as-synthesized polyethylenes were measured in terms of peak melting (Tpm ) and crystallization (Tpc ) temperatures, and % crystallinity (χ c ) using a differential scanning calorimeter (DSC Q20 0 0, Texas Instrument). The instrument was calibrated using indium. The experimental procedure reported in the literature was followed [10,22]. After completing the Cycle 1 DSC melting run, the sample was additionally heated isothermally at the final melting temperature for 5 min to (i) remove the influence of the thermal history and unmelted crystals (that could cause heterogeneous crystallization), and (ii) ensure equilibration. The data were acquired for each cycle and handled using the TA explorer software. The measured Cycle 3 χ c was subsequently used to calculate the polymer material density ρ polym , following the rule of additivity of volumes of polyethylene amorphous and crystalline phases [23]: Xc = (1/ρ polym − 1/ρ a )/(1/ρ c − 1/ρ a ); where ρ = density; a = amorphous phase; c = crystalline phase; and polym = polymer. For polyethylene, ρ c = 1.004 g ml−1 and ρ a = 0.853 gml−1 . Next, the amorphous volume fraction φ a was estimated using the relation φ a = (ρ c − ρ )/(ρ c − ρ a ). Table 1 reports the above thermal properties, densities, and amorphous volume fractions of the as-synthesized polyethylenes. The glass transition temperature Tg of the as-synthesized poly(1-hexene) was determined to be −49.54 °C, using Cycle 3 DSC run and following the procedure reported in the literature [24,25]. This value matches the cited Tg range (−50 to −58 °C) [24–27].

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Table 1 Catalyst productivity and summary of as-synthesized polyethylene properties. Catalyst productivity and polyethylene properties

Catalyst productivity (g polymer/g cat hr) Weight average molecular weight Mw (gmol−1 ) Polydispersity index (PDI) Glass transition temperature Tg (°C) Isotacticity index mmmm% Peak melting temperature Tpm (°C) Peak crystallization temperature Tpc (°C) Crystallinity χ c (%) Polymer material density ρ polym (g/cm3 ) Amorphous volume fraction φ a Weight average lamellar thickness Lwav DSC GT (nm) Most probable lamellar thickness LMPDSC GT (nm) Variance σ lamellar thickness (nm) Half crystallization time t1/2 (min) (min−1 ) Average crystallization rate t1−1 /2 Level of undercooling at t1/2 (θα = 0.5 ) Level of crystal surface free energy at t1/2 (Dα = 0.5 ) × 103 Critical stable crystallite sequence number at t1/2 (n∗α = 0.5 ) Level of undercooling at most probable mass fraction of LTD (θ MP ) Level of crystal surface free energy at most probable mass fraction of LTD (DMP ) × 103

Mixed crystal Z−N catalyst: TiCl3 •1/3AlCl3

Supported metallocene catalyst: SiO2 /MAO/(n BuCp)2 ZrCl2 NMW E−1-hexene copolymer

UHMW poly(1-hexene)

UHMW PE

NMW PE

50.00 2.18 × 106 3.32 −49.5 77.10

24.00 3.08 × 106 3.14

101.82 2.29 × 105 1.44

116.36 1.90 × 105 1.79

133.0 118.0 56.40 0.932 0.476 22.61 23.41 16.68 6.25 0.16 0.182 2.19 98 0.095

126.9 110.0 54.06 0.937 0.505 11.68 12.68 3.19 1.24 0.81 0.144 2.40 120 0.012

112.0 96.5 31.62 0.905 0.725 7.41 8.47 2.43 1.28 0.78 0.350 1.39 34 0.189

2.70

2.53

2.05

2.7. Microstructure of ethylene−1-hexene copolymer and poly(1-hexene) The microstructural parameters, including average ethylene and 1-hexene mol% in the synthesized copolymer, were determined using 13 C NMR spectroscopy. For this, a Bruker 600 MHz AVANAC III spectrometer equipped with Bruker Broadband Observe (BBO) PROBE (Bruker BioSpin, Rheinstetten, Germany) was used. A typical sample was prepared by dissolving about 20 mg of the polymer in 0.6 ml of deuterated 1,2,4 trichlorobenzene (TCB) using a 5 mm NMR tube. The spectra were recorded using a 30° flip angle 1D sequence with power-gate decoupling using standard 1D 90° pulse sequence and standard Bruker (zgig30) program. The spectral width was set at 282 ppm corresponding to 42.613 kHz and digitized into 64 k data points. The excitation pulse duration was 8.8 μs at 59.6 W. Exponential line broadening of 1 Hz was applied before Fourier transformation. To achieve good signal to noise ratio, each spectrum was recorded by collecting at least 12 k transients with 7 s recycle delay. Chemical shifts were adjusted using TCB signal as the internal chemical shift reference; then the spectra were visually phased; base lines were adjusted manually. Bruker Topspin 2.1 software (Bruker BioSpin, Rheinstetten, Germany) was used in all NMR experiments to collect and analyze the data. Using the NMR data, the monad, diad, and triad mole fractions, average E−1-hexene copolymer composition, and the copolymer microstructural parameters of our interest were calculated using the procedure and relations reported in the literature [28,29]. Table 2 lists the triad and monad mole fractions (average copolymer compositions), and the copolymer microstructural parameters. The NMR sample solutions for polyethylene and poly(1-hexene) were prepared and analyzed as summarized above. The tacticity parameters of poly(1-hexene) were determined following the procedure reported in the literature [25,30–33]. 2.8. Flory’s equilibrium theory polyethylene sequence length distribution, melt behavior, and crystallization The thermodynamically predicted theoretical distribution of ethylene sequence length (between the pendant n-butyl side chains), that is, the equilibrium crystal length was calculated using

Table 2 Microstructural parameters of the as-synthesized NMW metallocene E−1-hexene copolymer. Copolymer microstructural properties

NMW metallocene E−1-hexene copolymer

[EEE] [EEC] [CEC] [ECE] [ECC] [CCC] Average [E] mole fraction Average [1-hexene] mole fraction Ethylene reactivity ratio rE Comonomer reactivity ratio rH Average reactivity ratio product a Ethylene perpetuation probability p Copolymer type (based on rE rH  values) Average ethylene sequence length nE av NMR nE MPNMR Flory (most probable number of ethylene units)

0.757 0.102 0.0 0 0 0.031 0.110 0.0 0 0 0.860 0.140 10.24 0.59 5.99 0.930 Blocky 19 10

E = Ethylene; H = 1-hexene. nE MPNMR Flory and nE MPDSC GT are determined from Figs. 4 to 11, respectively. rE = kEE /kEH , and rH = kHH /kHE , where kEE and kHH are ethylene and 1-hexene terminal model homo-propagation rate constants. kEH and kHE are the cross-propagation rate constants. a Estimated by using the relationships listed in References 28 and 29.

the Flory equation [7,13]. According to this, the normalized weight fraction Wn of the sequence of n ethylene units (bounded at either end by 1-hexene) is related to the ethylene perpetuation probability p, that is, the probability that a crystallizable unit is succeeded by another such unit, through the following expression:

Wn = n(1 − p)2 pn−1

(5)

where p is also defined as sequence propagation probability. The plot of Wn versus n, as per Eq. (5), is called Flory’s sequence length distribution. For a statistical copolymer with very long chains, p is related to the experimental reactivity ratio product rE rC  and ethylene mole fraction XE as follows [13,34]: 1/2

p=1−

1 − [1 − 4(1 − rE rC  )XE (1 − XE )] 2(1 − rE rC  )XE

rE rC  = 1

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bl ocky and alternating

(6)

p = XM + (1 − XM )/2 = XE + (1 − XE )/2

(7) where XM is the mole fraction of the monomer. Subscript c stands for comonomer. XE is the mole fraction of ethylene monomer. The term (1−XM )/2 or (1−XE )/2 is the correction due to the incorporation of one −CH2 − unit (or half C2 H4 unit) per insertion of one α -olefin comonomer and exclusion of the pendant n-butyl group from the chain fold. Note that the perpetuation probability p shows the aliased interaction of rE rC  and XE . Therefore, varying combinations of rE rC  and XE may give the same value of p. Consequently, the definition of block copolymer for p> XE and alternating copolymer for p< XE may be misleading [35,36]. See Ref. [37] for the details. The ethylene perpetuation probability p can be used to study copolymer melt behavior and crystallization by calculating the critical (limiting) sequence number n∗ of the stable crystallite, that equilibrates with the melt at a given temperature, using the following equations [8–10]:

n ∗ (T ) =

 

θ

−1 DXE ln + ln p p



+ 2 ln

 1 − p  1 − e −θ

T RTm0 Hu T = (dimensionless ) RTm0 T

θ

ln

D 4 ( 1 − e −θ )

∞

1 n=1 n

2

+

(9)



e−θ 2 (1−e−θ )

−

1−p 1−e−θ

2

+ n∗



1 1−p

−

1 1−e−θ



[ n ∗ ( 1 − e −θ ) + e −θ ]

(12)

where XE is the mole fraction of the crystallizable ethylene units. Eq. (12) is the revised version of the original derivation by Flory [7]. It adds the second term in the denominator to account for the reduction in the amount of amorphous copolymer due to the crystalline phase [8,38–40]. For a polyethylene homopolymer, XE = 1 and p = 1. This reduces Eq. (12) to fc = 1. 2.9. Modeling of polyethylene crystallization kinetics A new nonisothermal crystallization model [18] for crystalline polymer using the Avrami–Erofeev equation was published by us in 2013. This model, with detailed assumptions and mathematical derivation, and experimental evidence, is reported in the above reference. Here, a summary is provided as follows. The nonisothermal Avrami–Erofeev polymer crystallization rate can be written as:





Ea 1 1 dα k0 = × exp − − dT β R T T0



× n (1 − α ) [− ln (1 − α )]

n−1 n

(13) where we define the following:

R

T

1 T0

−

n−1 n

(14)

 (15)

Ea (apparent crystal l ization energy ) = Egrow − Enucl

(16)

k0 (overal l crystal l ization f requency f actor ) (10)

∞ 21 θ n

p (1−p)2

 E 1 a

In Eq. (9), Tm0 is the melting point of a polyethylene perfect crystal, Hu is the heat of fusion for a C2 H4 repeat unit, σ e is the polyethylene basal surface energy, and ao is the cross-sectional area of a polyethylene chain. For a constant cooling rate nonisothermal crystallization process in a typical DSC experiment, Eq. (8) evaluates the temperature-dependent dynamic critical crystallite stability. Only ethylene sequences, greater than n∗ , form stable crystals in copolymers. With increasing temperature, the crystals melt and n∗ increases, and n∗ → ∞ represents the thickest possible crystals [8]. A particular feature of Eq. (8) is that it includes p which is a function of and XE . Catalyst mediates rE rC  and XE . Therefore, this equation indirectly investigates the effect of catalyst on copolymer melt behavior and crystallization. Eqs. (8)–(10) can be extended to a polyethylene homopolymer by assuming that it can be represented by a random copolymer with highly sparse incorporation of the comonomer. Therefore, using p → 1, the equivalent series expansion of ln(1− p), and (1−1/2+1/3−1/4+1/5− ……) = ln 2, we may temptingly derive the following expression:

1



1 − XE pn∗ −1

k (T ) = k0 × exp −

D (cr ystal sur f ace f ree energy effect )

n ∗ (T ) = −

−1

f (α ) = n (1 − α ) [− ln (1 − α )]

Hu Tm0 − T

= e−2σe ao /RT (dimensionless )

∗

(dimensionless ) (8)

θ (level o f undercooling) =

2

fc =

rE rC  = 1 random

XE (1 − p) pn

5

(11)

n=1

However, is non-convergent. Therefore, for the polyethylene homopolymer, we propose to use Eqs. (8)–(10) with p = 0.9999. This value of p corresponds to the existence of a temperature-independent common inflexion point, which we found using computer simulation. For the random copolymer of ethylene with 1-hexene (rE rH  = 1), Flory’s equilibrium crystallinity fc is given by [7,8,38–40]:



=

ks N0 V0



kgrow,0 = f ( T ) knucl,0

(17)

 E 1 grow

kgrow (T ) = kgrow,0 × exp −

R

T

 E 1 nucl

knucl (T ) = knucl,0 × exp −

R

T

−

−

1 T0



 1 T0

(18) (19)

In the above equations, f(α ) is called Avrami–Erofeev nonisothermal crystallization function, and α is the temperature- or time-dependent volume fraction of the molten polymer solidified due to cooling. Therefore, α concerns the phase morphology of the whole sample (melt plus solid). It is called relative crystallinity or degree of crystallization. β is the cooling rate. n is the dimension of the growing crystal. No is the number of germ nuclei, that is, the potential nucleus formation sites/defects. Vo is the initial volume of the molten polymer. Ks is the shape factor for the growing nuclei. kgrow,0 and Egrow are the frequency factors and activation energy of crystal growth, respectively. Knucl,0 and Enucl represent the corresponding terms for nucleation, respectively. R is the universal gas constant, and To is the centering temperature. The Avrami index n, in Eq. (13), illustrates two aspects—the crystal dimension and the nature of nucleation process. Therefore, n is written in terms of the following two components [41]:

n = nd + nn

(20)

where nd shows the dimension of the growing crystals. Theoretically speaking, nd can be only integers— 1, 2, and 3— that correspond to one-, two-, and three-dimensional crystals formed, respectively; and nn represents the nucleation process. In principle, it should be 0 or 1, where 0 refers to instantaneous (athermal/heterogeneous) nucleation; and 1, to sporadic (thermal/homogeneous) one. For many systems, the model-predicted

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n turns out to be a non-integer which is attributed to 0≤ nn ≤ 1. This means that instantaneous and sporadic nucleations may co-occur [42]. A balance between thermodynamic and kinetic factors influences the value of nn [35]. α is related to the corresponding weight fraction relative crystallinity α w through the following expression [18]:

α=

αw αw + ρc /ρa (1 − αw )

(21)

where ρ c and ρ a are the densities of the crystalline and amorphous phases, respectively. For polyethylene, the values of ρ c and ρ a have been reported earlier. The relation α versus T is called relative crystallinity or degree of crystallization profile. α w can be calculated from a typical constant cooling rate nonisothermal DSC experiment by using Eq. (22):

T dH  dT H ( T ) T dT αw = = Ti  Htotal f dH dT Ti

(22)

dT

where Htotal is the maximum enthalpy value reached at the end of the nonisothermal crystallization process and H(T) is the enthalpy evolved as a function of crystallization temperature T. Ti and Tf represent the start and completion temperatures of crystallization, respectively. Htotal and H(T) both can be acquired through the software of a standard differential scanning calorimeter (DSC). Using Eq. (21), the DSC-generated α w can be converted to the corresponding volume fraction α . The experimental confirmation of the above new nonisothermal crystallization model is available in one of our recent publications [43]. Also, note that some of the literature models require both isothermal and nonisothermal experiments to determine the designated parameters. See Harkin-Jones et al. that comprehensively reviews the aforementioned literature models [44]. 3. Results and discussion 3.1. Polyethylene synthesis, backbone structures, and bulk thermal properties In this section, we discuss the synthesis and backbone structures of the experimental polyethylenes, which include (i) MAOmodified silica-supported (n BuCp)2 ZrCl2 -mediated PE homopolymer and E−1-hexene copolymer as well as (ii) TiCl3 •1/3AlCl3 mediated UHMW PE and poly(1-hexene). We address all the polyolefin backbones in terms of molecular chain length, that is, weight average molecular weight Mw . However, E−1-hexene copolymer is additionally described in terms of the following microstructural properties: i.

13 C

NMR-determined average ethylene sequence/block length nE NMR ; and ii. The Flory ethylene sequence length distribution (SLD) between the pendant branched n-butyl short side chains, and its most probable value nE MPNMR-Flory (number of ethylene units corresponding to the peak of the distribution). We particularly address the influence of ethylene SLD on copolymer thermal properties and phase morphology. Table 1 reports that the Mw s of metallocene PE homopolymer and E−1-hexene copolymer are 2.29 × 105 and 1.90 × 105 gmol−1 , respectively. Therefore, the introduction of 1-hexene comonomer decreases the Mw by 60.67%. This shows that 1-hexene acts as a strong chain-transfer agent. The related chain transfer reactions, according to the literature [23], include the following: i. Route A: With ethylene as the last inserted unit, β -hydrogen transfer to Zr+ and/or to ethylene or 1-hexene.

ii. Route B: With ethylene as the last inserted unit, 2, 1 misinsertion of 1-hexene, followed by β -hydrogen elimination to the Zr+ active sites; and iii. Route C: With ethylene as the last inserted unit, 1, 2 insertion of 1-hexene. On the other hand, the following chain transfer reactions relate to ethylene or 1-hexene homopolymerization [23]: i. β -H transfer to the active metal center and the monomer. ii. Chain isomerization followed by β -H transfer to the monomer; and iii. The re-insertion of a macromer into the growing polyolefin chain, followed by β -H transfer to the monomer and/or to the catalyst transition metal Zr+ . The relative rate between chain propagation and chain termination controls the molecular weight distribution (MWD) and Mw . Therefore, from the kinetics point of view, 1-hexene increases the (n BuCp)2 ZrCl2 -mediated rate of the overall chain transfer reaction (that comprises the aforementioned Route A−Route C elementary reactions) over that of the chain propagation reaction. Hence, Mw drops. Note that chain propagation and termination kinetics is strongly catalyst-specific. The Z-N TiCl3 •1/3AlCl3 catalyst is a mixed crystal and has dual Lewis acid sites (Ti3+ and Al3+ ). The NH3 -TPD experiment shows that the strong Ti3+ acid site concentration (22.48 cm3 g−1 ) is much higher than that of the weak Al3+ acid sites (2.22 cm3 g−1 ). Also, see Fig. 1. TiCl3 •1/3AlCl3 particularly synthesizes only UHMW PE (Mw = 3.083 × 106 gmol−1 , %crystallinity χ c = 56.40, Tpm = 133.01 °C) and partially isotactic amorphous UHMW poly(1hexene) (Mw = 2.18 × 106 gmol−1 , glass transition temperature Tg = − 49.5 °C). Therefore, under both situations, the rate of the overall chain transfer reaction, consisting of the above homopolymerization chain termination reactions, relative to the chain propagation reaction, drastically reduces to extremely low value. Consequently, UHMW PE and poly(1-hexene) chains are produced. The as-synthesized UHMW poly(1-hexene) microstructure is characterized in terms of its pentad isotacticity index; mmmm% = 77.10, which indicates the occurrence of a fair amount of misinsertions. Hence, it is partially isotactic. The mmmm% was calculated by dividing the peak area of mmmm with the corresponding sum of all the carbon α -pentades (mmmm, mmmr/rmmr/mmrr, mmrm/rmrr, mrmr, rrrr, mrrr, and mrrm) [30–33]. See Fig. 2 which shows that the 13 C NMR spectrum of the experimental poly(1hexene) has broad peak at δ = 34.00−35.80 ppm, and has no sharp peak at δ = 34.57 or 35.00 ppm [isotactic poly(1-hexene)] [30,31,33,45] and at δ = 34.00 ppm [syndiotactic poly(1-hexene)] [31]. The calculated mrrm% (3.17) is higher than mmrm% (0.65%). Therefore, 1-hexene, in the presence of TiCl3 •1/3AlCl3 , polymerizes following the enantiomorphic site control stereoselective mechanism. This means that the catalyst-dictated coordination environment and symmetry, unlike the last-inserted monomer, decides poly(1-hexene) microstructure [46]. See Scheme 1 for the details [47]. Using TiCl3 •1/3AlCl3 , we also performed polymerization to synthesize copolymers of ethylene with α -olefins such as 1-hexene, 1-octene, 1-decene, and 1-dodecene. The resulting products were analyzed using 13 C NMR spectroscopy. Fig. 3 shows the NMR spectra. Each spectrum overlaps with that of pure polyethylene, and differs from that of E-1-hexene copolymer prepared by the supported (n BuCp)2 ZrCl2 . Therefore, TiCl3 •1/3AlCl3 does not coinsert any of these comonomers with ethylene in the growing polyethylene backbone. Consequently, these α -olefins do not copolymerize with ethylene. This happening can be attributed to the following two factors (one of which relates to the catalyst and the other to the comonomer):

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13

Fig. 2.

C NMR spectrum of the TiCl3 •1/3AlCl3 -synthesized poly(1-hexene).

Stereoerror

P Ti

7

Ti

P

P

LnTi m

Ligand controls stereoselectivity

m r

Enantiomorphic site control mechanism

r

m m m

Isotactic

P Pm

P

LnTi

Ti

m m m

P

r m m m Isotactic

Ti

P P

Pr

LnTi

Ti

r

Chain-end control stereoselectivity

r r

Chain-end control mechanism

m r

r

r

Syndiotactic

Scheme 1. Stereoselectivity mechanism of TiCl3 •1/3AlCl3 -mediated 1-hexene polymerization [9].

i. The unfavorable coordination environment of the dual Lewis acid TiCl3 •1/3AlCl3 , particularly for the copolymerization of the above prochiral linear α -olefins with achiral ethylene; and ii. The steric encumbrance and competitive diffusion limitation, introduced by the above α -olefins, which hinder them from accessing to the active metal-carbon (M−C) catalyst site. The 13 C NMR-determined enantiomorphic site control stereoselectivity in 1-hexene homopolymerization favors the first argument. Also see Scheme 1. The copolymerization of ethylene with alike higher linear α -olefins, using supported Z-N catalysts such as MgCl2 (THF)2 /VCl4 and MgCl2 (THF)2 /TiCl4 by Białek and Czaja [48], complements the second argument. The above figure also confirms that the as-synthesized UHMW and NMW PE backbones are linear and the NMW E-1-hexene copolymer backbone is characterized with n-butyl branch. Based on the 13 C NMR-determined rE rH  value of 5.99

1, we conclude that the supported C2v symmetric (n BuCp)2 ZrCl2 produces a blocky E-1-hexene copolymer. Fig. 4 shows its Flory sequence length distribution (SLD) which has been calculated using the above rE rH  value, average ethylene mole fraction XE , and Eqs. (5) and (6). We also plot the reference random copolymer SLD with the same XE and rE rH  = 1 (Eq. 7). We define the reference copolymer as an ideal copolymer having XE equal to that of the experimental copolymer. Note that Eq. (12) calculates the

Flory’s equilibrium crystallinity fc of this reference copolymer. The comparison of these two distributions evaluates to what extent the positioning of ethylene and 1-hexene along the copolymer backbone deviates from the ideal situation when the probability of inserting ethylene into the polymer backbone equals that of 1-hexene, that is, rE rH  = 1. This identity relation, in other words, means that the relative rates at which ethylene and 1-hexene are incorporated into the growing copolymer chains are the same [49]. The above reference ethylene SLD shows that the NMW metallocene blocky E−1-hexene copolymer appreciably deviates from random sequence distribution. rE rH , greater than 1, broadens the Flory SLD. Now, we compare the SLD of a random UHMW E−1-hexene copolymer, synthesized by Zhu et al. [50], using a C2 symmetric fluorinated bis(phenoxyimine) titanium(IV) catalyst (also called an FI catalyst) with that of our reference NMW E−1-hexene copolymer. This post-metallocene FI catalyst has the coordination environment and symmetry that favor UHMW E−1-hexene copolymerization. Its ethylene perpetuation probability p is 0.955 whereas that of our reference copolymer, also having rE rH  = 1, is 0.930. This shows that a minor change in p drastically widens the SLD. Therefore, the change in rE rH  or p, which is dictated by the specific catalyst type, can broaden SLD. This UHMW copolymer is characterized as follows: XE = 0.910, Mw = 1.28 × 106 gmol−1 , PDI = 1.18 (living copolymerization), %crystallinity χ c = 20.28, Tpm = 62.4 °C, and Tg = −53.5 °C (elastomer) [50].

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Fig. 3. Comparison of

13

C NMR spectra of the as-synthesized PEs.

The effect of the preceding drastic broadening of ethylene SLD on the UHMW E−1-hexene copolymer thermal properties and phase morphology is very pronounced. χ c and Tpm significantly decrease, and crystalline and amorphous phases co-develop. Consequently, an UHMW elastomer is produced. A similar effect of SLD broadening on the experimental metallocene NMW E−1-hexene copolymer thermal properties and phase morphology is also noticed. Hence, ethylene SLD plays a more important role in this context than other structural factors such as average comonomer content and molecular weight, and molecular weight distribution (PDI). Therefore, PE copolymers with new properties and applica-

tions can be synthesized by manipulating the ethylene SLD. This is a significant finding. Refer to Table 1 and compare its Tpm , Tpc , χ c , and amorphous volume fraction φ a with those of the NMW and UHMW PEs. Tpm , Tpc , and χ c decrease while φ a increases. Therefore, 1-hexene induces structural/enchainment defects and partially disrupts the crystal package of the polyethylene chains. The resulting n-butyl side chains, because of steric hindrance, block disentanglement processes in the amorphous phase. This prevents the whole of the regular ethylene sequences from incorporating into the crystallization lamella; resists crystallization lamellar thickening; and most likely changes intermolecular topological structure. Consequently the crystalline phase decreases and the amorphous phase increases [12]. Now, we discuss the crystallization phenomena of the above metallocene NMW and the nonmetallocene UHMW E−1-hexene copolymers from the viewpoint of minimum crystallizable ethylene sequence length nmc . The literature cites that a minimum (critical) ethylene sequence length is required for the backbone to crystallize [51–53]. For an E−1-hexene copolymer, nmc has been reported to be 32 [53]. Fig. 4 shows that for the metallocene NMW E−1hexene copolymer (χ c = 31.62%), sequence length n < 32 mostly contributes to crystallization. On the other hand, for the UHMW copolymer (χ c = 20.28%), n < 32 as well as n > 32 comparably contribute. Therefore, both short and long ethylene sequences crystallize. For either copolymer, the most probable n is less than 32. Similar is the remark about the average ethylene sequence length nE av NMR of the NMW E−1-hexene copolymer. See Table 2.

3.2. Melting behavior and crystallization of the as-synthesized polyethylenes: flory’s equilibrium theory perspective In this section, we discuss the melting behavior and crystallization of the as-synthesized polyethylenes from Flory’s equilibrium theory perspective. Hence, we particularly consider the effects of temperature-dependent level of undercooling θ , crystal surface free energy D, and critical stable crystallite sequence number n∗ on the above subject. See Eqs. (8), (9), and (10). We consider a heating and cooling rate of 10 °C min−1 , which is mostly reported to be used in the literature.

Fig. 4. Comparison of Flory ethylene sequence length distributions of E−1-hexene copolymers. Zhu et al. 2015 corresponds to Reference 50.

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1

0.8

0.8

Relative crystallinity α

Relative crystallinity α

M. Atiqullah et al. / Journal of the Taiwan Institute of Chemical Engineers 000 (2017) 1–15

0.6

0.4 UHMW PE NMW PE

0.2

9

UHMW PE NMW PE

0.6

NMW E–1-hexene copolymer 0.4

0.2

NMW E–1-hexene copolymer 0

0 0

0.2

0.4

0.6

0.8

1

Level of undercooling θ Fig. 5. Variation of as-synthesized PE relative crystallinity with level of undercooling.

1

UHMW PE

Relative crystallinity α

NMW PE

0.8

NMW E–1-hexene copolymer

0.6

0.4

0.2

0 0.0005

0.001

0.0015

0.002

0.0025

0.003

Crystal surface free energy D Fig. 6. Variation of as-synthesized PE relative crystallinity with crystal surface free energy.

The crystallization profiles— relative crystallinity α versus temperature T relations— feature very fast primary and slow (impinged) secondary crystallizations. Hence, we first evaluate the influence of θ and D on them for the as-synthesized polyethylenes. θ and D were calculated using Eqs. (9) and (10), respectively, and the Cycle 2 DSC data. Both PEs were approximated by using p = 0.9999. In these calculations, the following parametric values were used [8,54]: Tm0 (equilibrium melting temperature of a perfect crystal) = 145.5 °C, Hu (heat of fusion for a C2 H4 repeat unit) = 8.284 kJ mol−1 , and σ e (crystallite specific surface free energy) = 90 mJm−2 . Fig. 5 plots relative crystallinity α as a function of the level of undercooling θ . A common variational trend is noticed. A minor increase in θ sharply enhances the primary crystallization of the UHMW PE as well as the NMW PE and E−1-hexene copolymer. Consequently, this crystallization phase proceeds very fast. On the other hand, the slower secondary crystallization experiences a milder impact of θ on α . Structural factors, such as striking difference in MW or the incorporation of 1-hexene in the polymer backbone, also show similar effect of θ on α . This is how primary and secondary crystallizations are differently affected as θ increases while the crystallization begins and proceeds to completion. Fig. 6 evaluates how crystal surface free energy D affects the concerned crystallization. The overall effect is converse to that of the level of undercooling θ on relative crystallinity α . On the con-

0

100

200

300

400

500

600

Critical stable crystallite sequence number n* Fig. 7. Variation of as-synthesized PE relative crystallinity with critical stable crystallite sequence number.

trary, the slower secondary crystallization incurs a sharp decrease in D. This happening deviates from what occurs during primary crystallization. The effect of θ and D on apparent crystallization energy Ea will be addressed later on. From the above findings of Figs. 5 and 6, we conclude that the very rapid primary crystallization occurs due to a relatively small increase and decrease in the level of undercooling and crystal surface free energy, respectively. On the other hand, the very slow secondary crystallization happens because of the very sharp increase and decrease in the level of undercooling and crystal surface free energy, respectively. These phenomena are not affected by the structural variations of the PE backbone. To the best of our knowledge, these are insightful new explanations for the observed characteristics of the fairly fast primary and very slow secondary crystallizations. Fig. 7 pursues the progression of the critical (limiting) stable crystallite sequence number n∗ as the crystallization proceeds. Therefore, it plots the relative crystallinity α as a function of n∗ . We calculated n∗ , applying Eqs. (8)–(10), and using the following parametric values: Tm0 = 418.7 K, Hu = 8.284 kJ mol−1 , σ e = 90 mJ m−2 , and ao = 1.10 × 105 m2 mol−1 [8,54]. Note that as per Eq. (8), n∗ originates from the combined effect of the level of undercooling θ and crystal surface free energy D. The following interesting findings can be listed: i. At α = 0.5, n∗ can be rated as NMW PE > UHMW PE > NMW E−1-hexene copolymer whereas t1/2 oppositely varies as UHMW PE > NMW PE ≈ NMW E−1-hexene copolymer. t1/2 is inversely related to the average crystallization rate. Therefore, the above rating shows the influence of MW and average crystallization rate on n∗α =0.5 . ii. For the UHMW and NMW PEs, at α ≥ ∼0.6, n∗ is approximately same; the α versus n∗ plots practically overlap. Medium to small crystallites develop. At α < ∼0.6, for the UHMW PE, large to medium crystallites appear. Also, the decrease of n∗ with respect to increasing α is relatively gradual. For the NMW PE, dα /dn∗ remains fairly steep. Therefore, n∗ moderately decreases as α increases, and medium crystallites consistently predominate from the beginning. In conclusion, the effect of the remarkable difference in MW on n∗ is only manifest at α < ∼0.6. iii. For the NMW E−1-hexene copolymer, the α versus n∗ plot, benchmarked against that of the NMW PE, shows that the introduction of 1-hexene into the polymer backbone (that introduces ethylene SLD), combined with MW, significantly affects the critical stable crystallite sequence. The growth of moderate to small crystallites prevails during the course of crystallization.

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0.7

0.6 0.5

0.6 0.5

0.4

0.4

0.3

0.3

0.2

0.2

0.1

0.1

0

0 318

338

358

378

398

418

438

The 20−83 range of n∗ intriguingly falls within the ethylene SLD for the above copolymer reported in Fig. 4. In other words, the above plot demonstrates the impact of ethylene SLD on n∗ . As discussed above, the variation of n∗ with α concludes that nonisothermal crystallization grows small, medium, and large crystallites. The polarized light microscopy experiment conducted by Sajkiewicz et al. [55] fairly supports this model prediction. Structural factors such as average molecular weight and 1-hexene content, and ethylene SLD affect this crystallite size distribution. Fig. 8 compares the temperature-dependent instantaneous (dynamic) crystallinity χ profiles for the as-synthesized PEs. χ was estimated using the following relation:

dH 

Ti

dT



H 0f

dT

NMW PE

0.5

NMW E–1-hexene copolymer

0.4 0.3 0.2 0.1

0.0

Fig. 8. Temperature-dependent as-synthesized PE instantaneous crystallinity profiles versus the corresponding modified Flory’s equilibrium crystallinity profiles. The broken line above the solid circles indicates the equilibrium crystallinity for the assynthesized NMW E−1-hexene copolymer.

T

UHMW PE

0.6

0.0

Temperature (K)

H ( T ) χ (T ) = = H 0f

Temperature-dependent instantaneous crystallinity χ

0.7 Flory's equilibrium crystallinity for PE homopolymer = 1 UHMW PE NMW PE NMW E–1-hexene copolymer

Flory's equilibrium crystallinity fc

Temperature-dependent instantaneous crystallinity χ

0.7

(23)

where Hf ° is the heat of fusion (290 Jg−1 ) of the perfect PE crystal (of infinite lamellar thickness and molecular weight). Therefore, χ (T), unlike relative crystallinity α (T), is heat of fusion-based crystallinity that concerns the phases of the material solidified from the cooling melt. The Cycle 2 DSC crystallization data were applied to the above calculation. For each PE, the χ versus T relation shows the following common trend. χ initially increases fairly sharply as T decreases with continued cooling. However, below a critical cooling temperature Tc, critical , it asymptotically flattens (which indicates hindrance (impingement) to the further development of crystallinity), and is not any further affected by the decrease in T and hence, by the level of undercooling θ and crystal surface free energy D. The χ profiles for UHMW and NMW PEs are similar. But the existing striking MW difference shows significant influence. Tc, critical of the former is less than that of the latter. Also, the effect of enchainment defect due to 1-hexene incorporation on χ profile is very pronounced. This defect spectacularly lowers the χ profile. Also, the Tc, critical is less than that of UHMW and NMW PEs. This means that reaching the asymptotic plateau is delayed. Only above Tc, critical , the following happens: i. χ shows to be temperature-dependent; and ii. χ increases as θ increases, and D and critical stable crystallite sequence number n∗ decrease with the decrease in T. The asymptotic value of χ equals to the Cycle 2 DSC %crystallinity of the PE samples. χ (asymptotic or non-asymptotic) is always much less than the corresponding Flory’s equilibrium crystallinity fc . This remark equally applies to the UHMW and

0.2

0.4

0.6

0.8

1.0

Relative crystallinity α Fig. 9. Relation between as-synthesized PE temperature-dependent instantaneous crystallinity and relative crystallinity.

NMW PEs as well as the NMW E−1-hexene copolymer. For a PE homopolymer fc , irrespective of difference in MW, is equal to unity. See Eq. (12) and the text below it. Hence, the linear PE fc , unlike the branched copolymer analogue, is not affected by θ and D. The predicted crystallinity difference of NMW E−1-hexene copolymer may be attributed to the topology and the eventual kinetic restraint with reference to Flory’s equilibrium theory [9,56–59]. Therefore, crystallinity may be improved by decreasing the topological and kinetic restraints. By topology, we mean the crystallizable ethylene sequence length distribution SLD (structural defect due to 1-hexene incorporation, Fig. 4), the density of chain entanglement, and the configuration of the folding lamellae. Note that according to the equilibrium theory, only sequences exceeding a certain critical length crystallize. The SLD decreases the average crystallite thickness Lwav DSC GT (Table 1) and reduces the longer sequences required for equilibrium. On the other hand, the resulting kinetic constraint imposes hindrance to nucleation and crystal growth, and impinges the growing centers. The crystallization of particularly the very long sequences becomes especially difficult. Consequently, the above copolymer does not achieve the structural topology stipulated by the equilibrium requirements, and its heat of fusion and crystallinity decrease. Fig. 9 correlates the temperature-dependent instantaneous crystallinity χ to the corresponding relative crystallinity α for each as-synthesized PE. All of them show that χ increases linearly as a function of α . PE Mw or 1-hexene mole fraction affects only affects the slope dχ /dα of this relation. dχ /dα increases as UHMW PE > NMW PE > NMW E−1-hexene copolymer. χ represents the heat of fusion-based crystallinity that concerns the phases of the material solidified from the cooling melt. On the other hand, α characterizes the phase morphology of the whole sample (melt plus solid). The linear relation signifies the following. α and χ occur and progress under thermodynamic equilibrium. Therefore, the effects of level of undercooling θ , crystal surface free energy D, and critical stable crystallite sequence number n∗ on α and χ are equivalent. Hence, Figs. 5–7 also map how θ , D, and n∗ affect χ . Now, we address the effects of the level of undercooling θ and crystal surface free energy D on the melting behavior of the as-synthesized PEs. θ and D were calculated using Eqs. (9) and (10), respectively, and the Cycle 3 DSC melting data. As per Flory’s thermodynamic equilibrium theory, melting and crystallization are both reversible. Fig. 10 shows that for all the PEs, melting first starts with the smaller lamellae at lower temperatures, and higher undercooling favors this. On the other hand, the larger lamellae melt later at higher temperatures, which corresponds to lower

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Fig. 10. Variation of lamellar thickness, level of undercooling, and crystal surface free energy as a function of cooling temperature. Plots for all the as-synthesized PEs fairly overlapped.

Fig. 11. Variation of the as-synthesized PE mass fraction of crystal lamellae melted, level of undercooling, and crystal surface free energy as a function of lamellar thickness.

level of undercoling θ . All the plots essentially overlap. In either case, the lamella melting temperature is practically less than the equilibrium melting temperature Tm o of the perfect crystalline PE (418.7 K). However, the effect of crystal surface free energy D on melting behavior of the above PEs differs from that of θ . Here, D, unlike θ , increases as the melting temperature increases, and consequently the smaller to larger lamellae sequentially melt. Overall, θ and D oppositely influence PE melting phenomenon. These are insightful findings, and to the best of our knowledge, have not been reported earlier in the literature. They can be explained as follows. Eq. (9) shows that θ linearly increases as a function of T (where T = T 0 − T ), and it mathematically expresses the m T undercooling profile of the experimental PEs with reference to the polyethylene perfect crystal (having infinite MW and lamellar thickness, and θ = 0). Hence, it eventually shows the temperature gradient effect on PE melting. On the other hand, D decreases exponentially as a function of non-dimensional crystal surface e ao free energy σRT . See Eq. (10). The above variational trends of θ and D are fundamentally related to the topologies of the lamellae and crystallite surface, which are affected by the architecture of the as-synthesized PE backbones. Refer to the influence of MW, PDI, and SLD (structural defect due to 1-hexene incorporation) on the varying values of Lwav DSC GT , LMPDSC GT , and σ lamellar thickness (Table 1). The latter three terms quantify the lamellar chain topology which is linear for all the above PEs.

Fig. 11 demonstrates that the mass fraction of lamellae melted at a given temperature/time, that is, the lamellar thickness distribution (LTD) cannot be directly related to the level of undercooling θ and crystal surface free energy D. Nevertheless, the LTD plots interestingly illustrate the influence of θ and D particularly on the progression of melting. This feature varies with the PE type. For the UHMW and NMW PEs, a significant mass fraction of the lamellae melts at lower θ and higher D. See the LTD at the right side of the corresponding vertical lines. However, for the NMW E−1-hexene copolymer, the situation differs. The population of lamellae melted at higher and lower θ s are approximately comparable while the whole mass melts at low to high Ds. All these findings follow the succssive leftward positioning of the LTDs of UHMW PE, NMW PE, and NMW E−1-hexene copolymer, respectively. This order of LTD positioning shows to be related to their MWs. But particularly for the copolymer, the MW effect is aliased with the 1-hexene incorporation, that is, the SLD. The plots of θ and D, as a function of lamellar thickness L, practically overlap for all the experimental PEs. θ decreases whereas D increases. We calculated the LTDs applying Gibbs–Thompson (GT) equation, the related literature expressions [10,13–17] and the Cycle 3 DSC data. We evaluated the most probable lamellar thickness LMPDSC GT from the peak of each lamellar thickness distribution. In these calculations, the following parametric values were used [8,54]: Tm0 (equilibrium melting temperature of a perfect crystal) = 145.5 °C, H 0f (heat of fusion per unit volume for the perfect crys-

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tal) = 290 Jcm−3 , and σ e (crystallite specific surface free energy) 0 = 90 mJm−2 . However, Tm,copolym (equilibrium melting temperature of the copolymer) was estimated as described in the literature [13–17]. The weight average lamellar thickness Lwav DSC GT was determined using the LTD. From the above figure, the most probable properties, for the experimental PEs, can be ranked as follows: i. Most probable laminar thichness LMPLTD : UHMW PE > NMW PE > NMW E−1-hexene copolymer. ii. Most probable level of undercooling θ MP : UHMW PE < NMW PE < NMW E−1-hexene copolymer (opposite to that of LMPLTD ); and iii. Most probable crystal surface energy DMP : UHMW PE > NMW PE > NMW E−1-hexene copolymer (same as that of LMPLTD ). The above ranking of LMPLTD and DMP matches that of the Tpc s and Tpm s of the as-synthesized PEs. See Table 1. Accordingly, the observed variation in Tpm s can be accounted for. But the difference in %crystallinity χ cannot be viewed likewise; χ of UHMW PE is comparable with that of NMW PE. Hence, the remarkable difference in chain lengths (MWs) hinders crystallite development. This explanation agrees with the DSC experiment of Barrales-Rienda and Fatou [2]. The report by Zhu et al. [4] supports our finding that the Tpc and Tpm of the UHMW PE are higher than those of the NMW PE. 3.3. Crystallization kinetics of the as-synthesized polyethylenes In this section, we address the nonisothermal crystallization kinetics of the as-synthesized polyethylenes in terms of apparent crystallization activation energy Ea , frequency factor ko , crystal dimension nd , and time-dependent nucleation mode nc , using a cooling rate of 10 °C min−1 . We particularly compare the UHMW PE crystallization model parameters with those of the NMW PE homo- and copolymer. Solution of crystallization model equation and estimation of kinetic triplet We solved Eq. (13) numerically, and estimated the kinetic triplet (ko , n, and Ea ) as follows. First, we modified it through separation of variables, and integrated the left hand side (LHS) analytically, and transformed the right hand side (RHS) into the Ea well-known temperature integral exp (− RT ) dT . See Eqs. (24) and (25). Next, we converted the temperature integral into a real part and an exponential integral part [60]. See Eq. (26). This is how we integrated the temperature integral, which we finally transformed to the non-linear algebraic form. We solved this modified model equation using the Mathematica 8.0 Nonlinear Model Fit software. To = 370 K was used as the centering temperature. The LHS of Eq. (26), containing − ln (1 − α ), has a point of discontinuity at α final = 1. This was resolved by approximating α final ∼ = 0.9999. Depending on the PE types, 30 to 45 experimental data points were considered for kinetic parameter estimation.





dα n(1 − α ) [− ln (1 − α )]

n−1 n

k0

=

β

 E 1 a

×exp −

R

T

−

1 T0



dT (24)

1 n

[− ln (1 − α )] = 1 n

[− ln (1 − α )] =

k0 exp

β k0 exp

β

Ea  

 E  a

RT0

exp −

Ea   RT0

RT

dT

 E  a

T exp −

RT

(25)

+

Ea Ea Ei − RT



R (26)

where

Ea Ei(− RT

) is the exponential integral of

Ea − RT

.

1

Relative crystallinity α

12

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0.8

0.6

0.4 UHMW PE 0.2

NMW PE NMW E–1-hexene copolymer

0 300

320

340

360

380

400

420

Temperature T (K) Fig. 12. Comparison of the model-predicted as-synthesized PE relative crystallinity profiles with those determined using DSC experiments.

We estimated the model parameters considering the following statistical criteria—95% confidence interval, coefficient of determination (R2 ), estimated variance, and standard error. The aforementioned Mathematica software eventually generates them. Convergence was accepted when the objective function changed less than the specified tolerance of 10−8 . For the sake of brevity and sufficiency, we list only R2 for each of the estimated kinetic parameters. See Table 3. 3.4. Evaluation of model performance and significance of the predicted results Now we evaluate the present nonisothermal Avrami–Erofeev model performance, and discuss the significance of the major findings. Fig. 12 compares the model-predicted relative crystallization profiles, determined using Eq. (13), with the corresponding DSC profiles for the experimental UHMW PE and the NMW PE homoand copolymer. It shows that the E−1-hexene copolymer crystallization profile, because of the defects introduced by the pendant n-butyl groups, shifts leftward with respect to the PEs. In other words, 1-hexene incorporation partially disrupts the polyethylene crystal package. The primary crystallization of UHMW PE is slower than that of NMW PE. This manifests the effect of UHMW (Mw = 3.08 × 106 gmol−1 ). The crystallization half-time can be rated as follows: t1/2 | UHMW PE > t1/2 | NMW PE ≈ t1/2 | NMW E−1-hexene PE . See Table 1. t1/2 is inversely related to the bulk (average) crystallization rate [61]. Hence, the Z-N UHMW PE, with far greater MW, crystallizes slower than the metallocene NMW PE. This prevents the further growth of crystallinity. See Table 1 χ (%crystallinity) values. In each of the above cases, the agreement between the model and experiment is excellent. This report is the first successful application of the present crystallization model to an UHMW PE. Notably, a single n accounts for the whole crystallization regime [induction period which corresponds to α ≤ 0.1, as well as very fast primary and slow (impinged) secondary crystallizations]. This result signifies that the mechanism of nucleation and crystal growth holds, irrespective of molecular weights, all throughout the nonisothermal PE crystallization. This illustrates how the present model overcomes the drawbacks and limitations of the arbitrary parameter-based nonisothermal crystallization models, which we summarized in our first study [18]. Also, see Harkin-Jones et al. [44] that reviews the various parametric nonisothermal crystallization models. Therefore, the assumption of change in crystallization mechanism, as reported in the literature, [44,62–65] does not hold. Fig. 12 also confirms that a single value of apparent activation Ea fits the well-known isokinetic Avram–Erofeev crystalliza-

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13

Table 3 Summary of model-predicted crystallization kinetic triplet. Nonisothermal crystallization parameters

UHMW PE

NMW PE

NMW E−1-hexene copolymer

k0 (min−1 ) Ea (kJ/mol) n R2 (Coefficient of determination)

0.316 ± 0.013 26.56 ± 1.07 3.03 ± 0.13 0.996

0.520 ± 0.025 57.81 ± 2.80 3.10 ± 0.15 0.994

0.571 ± 0.027 24.08 ± 1.16 3.06 ± 0.14 0.997

tion mechanism throughout each PE crystallization process (primary plus secondary). Therefore, Ea is essentially constant of crystallization time or temperature, relative crystallinity α , and the two competitive temperature-dependent equilibrium theory parameters— the level of undercooling θ , and the crystal surface free energy D. Also, see Figs. 5 and 6. Now we compare this finding with that published in the literature as follows. Supaphol et al. [66] modeled the nonisothermal crystallization activation energy of selected aromatic polyesters and showed that it is a function of the relative crystallinity α . Depending on the polymer structure, it either monotonically increases, or it first increases and then it decreases as α increases. This variational trend was ascribed to the dependence of nucleation energy barrier on temperature. However, note that this explanation ignores isokinetic concept, crystal growth, as well as primary and secondary crystallizations. Hence, their model prediction and explanation are insufficient and unacceptable. Table 2 shows that the model-predicted crystal dimension nd = 3, irrespective of PE type and molecular weight. Hence, the resulting PE crystals are spherulites, which involve three dimensional alignments of the polymer backbone lamellae. Further, according to Eq. (20), nn = 0.13−0.16. Therefore, the nucleation process for the subject PE crystallization is more close to instantaneous (athermal/heterogeneous) than sporadic (thermal/homogeneous). The pendant n-butyl groups in the backbone of the corresponding copolymer do not affect this nucleation phenomenon. We also observe that the subject PE Ea s can be ranked as EaNMW PE > EaUHMW PE ∼ EaNMW E−1-hexene PE . The above table also lists the model-predicted nonisothermal crystallization frequency factor k0 . See Eq. (17). Considering the relation between kinetics and thermodynamics, it can be shown that k0 ⇒ exp(S/R) or S ⇒ Rlnk0 where S is entropy of the system and R is the universal gas constant [6]. Therefore, the entropies of the experimental PE crystallization can be rated as SNMW E−1-hexene PE ∼ SNMW PE > SUHMW PE . This means that the Z-N UHMW PE, despite having enormously greater Mw , crystallizes with less entropic disorder than the metallocene NMW PE and NMW E−1-hexene PE. This finding may be ascribed to the difference in solid-state electronic environment and structure of the residual metallocene catalyst (SiO2 /MAO/(n BuCp)2 ZrCl2 ) from those of the Z-N catalyst (TiCl3 •1/3AlCl3 ), and their consequential interactions with the nuclei and polymer backbones. This argument is detailed in one of our earlier publications [61]. Here, we have also shown that the structure of the residual catalyst differs from that of the original catalyst. This structure can be derived from the overall polymerization mechanism. Based on the above overall findings of Fig. 12 and Table 3, this study confirms the invariance of activation energy articulated by Galwey and co-thinkers [67,68], and does not support the concept of variable instantaneous activation energy [62,66,69–73] which is practiced in analyzing nonisothermal crystallization kinetic data. This conclusion originates from the correct application of isokinetic concept and the current nonisothermal Avrami–Erofeev crystallization model, and the appropriate calculation algorithm that we developed. This is how we address in this study the mathematical artefact that exists in the literature.

Following the preceding discussion, we conclude that the nonisothermal primary and secondary crystallizations of the experimental Z-N UHMW PE, including those of the metallocene NMW PE and NMW E−1-hexene PE, occur isokinetically with constant apparent kinetic energy Ea . The crystals are spherullites while the nucleation process is more close to instantaneous (athermal/heterogeneous) than sporadic (thermal/homogeneous). The crystallization entropy seems to be residual catalyst-dependent. The UHMW PE, with far greater MW and less Ea and entropic disorder, crystallizes slower than the NMW PE. The model predicts that it is entropy-controlled. 4. Conclusions UHMW PE fabricates a large number of end-products. The fabrication process involves melting and crystallization of the polymer. Therefore, this report investigates the melt behavior and nonisothermal crystallization kinetics of as-synthesized UHMW PE, and NMW PE and E−1-hexene copolymer using a new nonisothermal crystallization model [18] (which we published in 2013), Flory’s equilibrium theory and ethylene sequence length distribution concept, Gibbs–Thompson equation, and nonisothermal DSC experiments. By applying the above approach to the problem, the relative crystallinity α , temperature-dependent instantaneous crystallinity χ , the crystallization kinetic triplet, and lamellar thickness and melting temperature have been duly correlated, as appropriate, to the level of undercooling θ , crystal surface free energy D, and critical stable crystallite sequence number n∗ . Consequently, insightful new results, interpretations, and explanations, related to the melting and crystallization of UHMW and NMW PEs and E−1-hexene copolymers, have been concluded. The superior performance of the present fundamental mechanistic crystallization model, compared to several existing literature models, has also been emphasized. In particular, the following can be listed: •

•

•

The catalyst coordination environment and symmetry, as well as achiral ethylene versus prochiral α -olefin comonomer steric encumbrance and competitive diffusion limitation affect the synthesis of UHMW PE, particularly the corresponding copolymer. The ethylene sequence length distribution (SLD) of UHMW and NMW E−1-hexene copolymers, compared to other structural factors such as average comonomer content and molecular weight, and molecular weight distribution, more pronouncedly affect the thermal properties and phase morphology. Peak melting temperature Tpm and %crystallinity χ c significantly decrease, and crystalline and amorphous phases co-develop, depending on the broadening of the SLD. Ethylene sequences shorter than the minimum crystallizable ethylene sequence length, irrespective of E−1-hexene copolymer MW, also crystallize. The temperature-dependent instantaneous (dynamic) crystallinity χ increases as the level of undercooling θ increases, and crystal surface free energy D and critical stable crystallite sequence number n∗ decrease with the decrease in T. The UHMW and NMW PE Flory’s equilibrium crystallinity fc = 1. Hence, θ and D do not affect this. But they influence fc of the NMW E−1-hexene copolymer. χ (asymptotic or temperature-

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•

•

•

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dependent non-asymptotic) is always much less the corresponding fc . Average molecular weight and 1-hexene content, and ethylene SLD vary the χ profile. Smaller lamellae first melt at lower temperatures and higher level of undercooling θ . The reverse applies to the larger lamellae. The crystal surface free energy D, unlike θ , increases as the melting temperature increases, and consequently the smaller to larger lamellae sequentially melt. These findings apply to all the experimental PEs, despite they have varying lamellar thickness distributions (LTDs). The nonisothermal primary and secondary crystallizations of the experimental Z-N UHMW PE, including those of the metallocene NMW PE and NMW E−1-hexene PE, occur isokinetically with constant apparent kinetic energy, which is also unaffected by the level of undercooling θ , and crystal surface free energy D. The crystals are spherulites while the nucleation process is more close to instantaneous (athermal/heterogeneous) than sporadic (thermal/homogeneous). The crystallization entropy appears to be residual catalyst-dependent. The UHMW PE, with far greater MW and less apparent crystallization energy and entropic disorder, crystallizes slower than the NMW PE. The very rapid primary crystallization originates from a small increase and decrease in the level of undercooling θ and crystal surface free energy D, respectively. In contrast, the very slow secondary crystallization occurs due to the very sharp increase and decrease in θ and D, respectively. However, θ and D, despite having opposite characteristics, do not change Ea , as a function of cooling temperature T. The variation of n∗ with α shows that nonisothermal crystallization grows small, medium, and large crystallites. Structural factors such as average molecular weight and 1-hexene content, and ethylene SLD affect this crystallite size distribution.

Acknowledgment The authors greatly acknowledge the financial support provided for this study by King Abdulaziz City for Science and Technology (KACST) via the Science & Technology Unit at King Fahd University of Petroleum & Minerals (KFUPM) through Project Number 14-PET283-04 as part of the National Science and Technology Innovation Plan (MAARIFAH). The technical assistance provided by the Center for Refining & Petrochemicals (CRP) at the Research Institute, KFUPM, Dhahran, Saudi Arabia; and NMR Core Laboratory, Thuwal, King Abdullah University of Science & Technology (KAUST), Saudi Arabia is also gratefully acknowledged. The support offered by Prof. J. C. Santamarina, Energy GeoEngineering Laboratory, KAUST to NMR assay is highly appreciated. The technical assistance of Mr. Anwar Hossaen and the gift of the experimental α -olefins by United Petrochemicals, an affiliate of Saudi Basic Industries Corporation (SABIC), are thankfully appreciated. References [1] Stein HL. Ultrahigh molecular weight polyethylenes (UHMWPE). Engineered materials handbook, 2. Materials Park, OH, USA: ASM International; 1998. p. 167–71. [2] Barrales-Rienda JM, Fatou JMG. Effect of molecular weight on the rate of crystallization of polyethylene fractions at high undercooling. Polymer 1972;13:407–11. [3] Fatou JG, Marco C, Mandelkern L. The influence of molecular weight on the regime crystallization of linear polyethylene. Polymer 1990;31:1685–93. [4] Yan Z, Linkai C, Shouzhi Y. Studies of non-isothermal crystallization and melting of ultra high molecular weight polyethylene. J Therm Anal Calorim 1995;45:329–33. [5] Parasnis NC, Ramani K. Nonisothermal crystallization of UHMW PE. J Therm Anal Calorim 1999;55:709–19. [6] Jackson K. Kinetic processes: crystal growth, diffusion, and phase transitions in materials. 2nd ed. Weinheim, Germany: Wiley-VCH Verlag Gmbh; 2012. (revised). [7] Flory PJ. Theory of crystallization in copolymers. Trans Faraday Soc 1955;51:848–57.

[8] Crist B, Howard PR. Crystallization and melting of model ethylene-butene copolymers. Macromolecules 1999;32:3057–67. [9] Ruiz-Orta C, Fernandez-Blazquez JP, Anderson-Wile AM, Coates GW, Alamo RG. Isotactic polypropylene with (3,1) chain-walking defects: characterization, crystallization, and melting behaviors. Macromolecules 2011;44:3436–51. [10] Atiqullah M, Anantawaraskul S, Emwas AM, Al-Harthi MA, Hussain I, Ul-Hamid A, et al. Effects of heterogeneous (n BuCp)2 ZrCl2 catalyst active center distribution on ethylene−1-hexene copolymer microstructure and thermal behaviors. Indus Eng Chem Res 2013;52:9359–73. [11] Atiqullah M, Adamu S, Al-Harthi MA, Hossain MM, Anantawaraskul S, Emwas AM, et al. Effect of supported (n BuCp)2 ZrCl2 catalyst active center multiplicity on crystallization kinetics of ethylene homo- and copolymers. J Taiwan Inst Chem Eng 2014;45:1982–91. [12] Atiqullah M, Adamu S, Malaibari ZO, Al-Harthi MA, Emwas AM. (n BuCp)2 ZrCl2 -catalyzed ethylene-4M1P copolymerization: copolymer backbone structure, melt behavior, and crystallization. AICHE J 2016;62:1688–706. [13] Hosoda S, Nozue Y, Kawashima Y, Suita K, Seno S, Nagamatsu T. Effect of the sequence length distribution on the lamellar crystal thickness and thickness distribution of polyethylene: perfectly equisequential ADMET polyethylene vs ethylene/α ;-olefin copolymer. Macromolecules 2011;44:313–19. [14] Flory PJ. Thermodynamics of crystallization in high polymers. IV. A theory of crystalline states and fusion in polymers, copolymers, and their mixtures with diluents. J Chem Phys 1949;17:223–40. [15] Chen F, Shanks RA, Amarasinghe G. Molecular distribution analysis of melt-crystallized ethylene copolymers. Polym Intl 2004;53:1795–805. [16] Musa RK, Lijun F, Tao H. A generalized equation for the prediction of melting temperatures of homopolymers and copolymers. Can J Chem Eng 2002;80:432–42. [17] Lijun F, Musa RK. Distributions of crystal size from DSC melting traces for polyethylenes. The Can J Chem Eng 2004;82:1239–51. [18] Atiqullah M, Hossain MM, Kamal MS, Al-Harthi MA, Hossaen A, Khan MJ, et al. Crystallization kinetics of PE-b-isotactic PMMA diblock copolymer synthesized using SiMe2 (Ind)2 ZrMe2 and MAO cocatalyst. AICHE J 2013;59:200–12. [19] Lonyi F, Valyon J. On the interpretation of the NH3 -TPD patterns of H-ZSM-5 and H-Mordenite. Micro Meso Mater 2001;47:293–301. [20] Cotton FA, Wilkinson G. Advanced inorganic chemistry: a comprehensive text. 4th ed. New York: John Wiley; 1980. [21] Wood-Adams PM, Dealy JM, DeGroot AW, Redwine OD. Effect of molecular structure on the linear viscoelastic behavior of polyethylene. Macromolecules 20 0 0;33:7489–99. [22] Wignall GD, Alamo RG, Laondono JD, Mandelkern L, Kim MH, Lin JS, et al. Morphology of blends of linear and short-chain branched polyethylenes in the solid state by small-angle neutron and X-ray scattering, differential scanning calorimetry, and transmission electron microscopy. Macromolecules 20 0 0;33:551–61. [23] Atiqullah M, Al-Harthi MA, Anantawaraskul S, Emwas AM. Ethylene homo- and copolymerization chain-transfers: a perspective from supported (n BuCp)2 ZrCl2 catalyst active centre distribution. J Chem Sci 2015;127:717–28. [24] Fengshou L, Haiyang G, Zhilong H, Haibin H, Fangming Z, Qing W. Poly(1-hexene) with long methylene sequences and controlled branches obtained by a thermostable α -diimine nickel catalyst with bulky camphyl backbone. J Polym Sci Part A: Polym Chem 2012;50:3859–66. [25] Irina VV, Sergei VK. The influence of cocatalysts on 1-hexene polymerization with various supported magnesium-titanium catalysts. Polym Bull 2006;57:129–38. [26] Cheremisinoff NP. Materials selection desk book. New Jersey: Noyes Publications; 1996. [27] Ellis B, Smith R. Polymers: a property database. 2nd ed. New York: CRC Press; 2008. [28] Seger MR, Maciel GE. Quantitative 13C NMR analysis of sequence distributions in poly(ethylene-co-1-hexene). Anal Chem 2004;76:5734–47. [29] Hung J, Cole AP, Waymouth RM. Control of sequence distribution of ethylene copolymers: Influence of comonomer sequence on the melting behavior of ethylene copolymers. Macromolecules 2003;36:2454–63. [30] Jayaratne KC, Sita LR. Stereospecific living Ziegler–Natta polymerization of 1-hexene. J Am Chem Soc 20 0 0;122:958–9. [31] Galland GB, da Silva LF, Nicolini A. Tacticity of poly−α -olefins: from poly−1-hexene to poly−1-octadecene. J Polym Sci: Part A: Polym Chem 2005;43:4744–53. [32] Vasilenko IV, Kostyuk SV, Gaponik LV, Kaputskii FN. Polymerization of 1-hexene on catalytic system TiCl4 -Al(C6 H13 )3 -Mg(C6 H13 )2 . Russ J Appl Chem 2004;77:295–8. [33] Tao X, Gao W, Huo H, Mu Y. Achiral Cs -symmetric half-sandwich scandium (III) complexes with imine-cyclopentadienyl ligands catalyze isotactic polymerization of 1-hexene. Organometallics 2013;32:1287–94. [34] Allegra G, Marchessault R, Bloembergen S. Crystallization of markoffian copolymers. J Polym Sci Polym Phys Ed 1992;30:809–15. [35] Muthukumar M. Nucleation in polymer crystallization. In: Rice S, editor. Advances in chemical physics. New York: John Wiley; 2004. p. 128. [36] Mandelkern L. Crystallization of polymers: equilibrium concepts, 1. Cambridge: Cambridge University Press; 2002. [37] Mirabella FM, Crist B. Melting behavior of polyethylene/α -olefin copolymers: Narrow composition distribution copolymers and fractions from Ziegler— Natta and single-site catalyst products. J Polym Sci Part B: Polym Phys 2004;42:3416–27.

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[38] Kilian HG. Partiell kristallisierende Copolymere als “pseudoeutektische Systeme. Kolloid-Zeitschrift & Zeitschrift für Polymere (Kolloid Z Z Polym) 1965;202:97–107. [39] Baur H. Zur Theorie der Kristallisation in Copolymeren. Kolloid-Zeitschrift & Zeitschrift für Polymere (Kolloid Z Z Polym) 1965;203:97–107. [40] Wunderlich B. Macromolecular physics, 2. New York: Academic Press; 1976. p. 259–71. [41] Avrami J. Kinetics of phase change. I. General theory. J Chem Phys 1939;7:1103–12. [42] Lorenzo AT, Arnal ML, Albuerne J. DSC isothermal polymer crystallization kinetics measurements and the use of the Avrami equation to fit the data: guidelines to avoid common problems. J Polym Test 2007;26:222–31. [43] Atiqullah M, Hossain MM, Adamu S, Hossaen A. Thermal behavior of PE-b-PMMA block copolymer: effect of multiple heating and cooling rates versus mathematical artefact. Polym Intl 2014;63:1824–34. [44] Mubarak Y, Harkin-Jones EMA, Martin PJ, Ahmad M. Modeling of non-isothermal crystallization kinetics of isotactic polypropylene. Polymer 2001;42:3171–82. [45] Babu GN, Newmark RA, Chien JCW. Microstructure of poly(1-hexene) produced by ansa-zirconocenium catalysis. Macromolecules 1994;27:3383–8. [46] Segal S, Goldberg I, Kol M. Zirconium and titanium diamine bis(phenolate) catalysts for α -olefin polymerization: from atactic oligo(1-hexene) to ultrahigh– molecular-weight isotactic poly(1-hexene). Organometallics 20 05;24:20 0–2. [47] Coates GW. Precise control of polyolefin stereochemistry using single-site metal catalysts. Chem Rev 20 0 0;10 0:1223–52. [48] Bialek M, Czaja K. The effect of the comonomer on the copolymerization of ethylene with long chain α -olefins using Ziegler–Natta catalysts supported on MgCl2 (THF)2 . Polymer 20 0 0;41:7899–904. [49] Young RJ, Lovell PA. Introduction to polymers. 3rd ed. New York, USA: CRC Press, Taylor & Francis; 2011. [50] Guo S, Fan H, Bu Z, Li B, Zhu S. Thermal and mechanical properties of ultrahigh-molecular-weight ethylene/1-hexene copolymers prepared by living polymerization with fluorinated bis(phenoxy-imine) titanium(IV) catalyst. Polymer 2015;80:109–14. [51] Burfield DR. Correlation between crystallinity and ethylene content in LLDPE and related ethylene copolymers, demonstration of the applicability of a simple empirical relationship. Macromolecules 1987;20:3020–3. [52] Chen Q, Luo H, Yang G. Influence of the sequence distribution on the crystalline structure of ethylene-dirnethylaminoethylmethacrylate copolymers. Polymer 1997;38:1203–5. [53] Luo H, Chen Q, Yang G. Studies on the minimum crystallizable sequence length of semicrystalline copolymers. Polymer 2001;42:8285–8. [54] Zhou H, Wilkes GL. Comparison of lamellar thickness and its distribution determined from DSC, SAXS, TEM and AFM for high-density polyethylene films having a stacked lamellar morphology. Polymer 1997;38:5735–47. [55] Gradys A, Sajkiewicz P, Minakov AA, Adamovsky S, Schick C, Hashimoto T, et al. Crystallization of polypropylene at various cooling rates. Mater Sci Eng A 2005;413:442–6.

[m5G;April 25, 2017;19:59] 15

[56] Alamo R, Domszy R, Mandelkern L. Thermodynamic and structural properties of copolymers of ethylene. J Phys Chem 1984;88:6587–95. [57] Alamo R, Mandelkern L. Crystallization kinetics of random ethylene copolymers. Macromolecules 1991;24:6480–93. [58] Alamo R, Mandelkern L. The crystallization behavior of random copolymers of ethylene. Thermochim Acta 1994;238:155–201. [59] Crist B, Finerman TM. Copolymer crystallization: approaching equilibrium. Polymer 2005;46:8745–51. [60] Gradshteyn IS, Ryzhik IM. Table of integrals, series, and products, 93. 4th edition. New York: Academic Press; 1980. [61] Atiqullah M, Hussain I, Al-Harbi A, Fazal A, Hossaen A. Crystallization kinetics of ethylene homopolymers: a new perspective from residual catalyst and resin molecular weight. J Polym Res 2012;19:9797–810. [62] Papageorgiou GZ, Achilias DS, Karayanndis GP. Estimation of thermal transitions in poly(ethylene napthalate): experiments and modeling using isoconversional methods. Polymer 2010;51:2565–75. [63] Acar I, Durmus¸ A, Özgümüs¸ S. Nonisothermal crystallization kinetics and morphology of poly(ethylene terephthalate) modified with poly(lactic acid). J Appl Polym Sci 2007;106:4180–91. [64] Olmos D, Dominguez C, Castrillo PD, Gonzalez-Benito J. Crystallization and final morphology of HDPE: effect of the high energy ball milling and the presence of TiO2 nanoparticles. Polymer 2009;50:1732–42. [65] Adhikari A, Lozano K. Effects of carbon nanofibers on the crystallization kinetics of polyethylene oxide. J Polym Res 2011;18:875–80. [66] Supaphol P, Dangseeyun N, Srimoaon P, Nithitanakul M. Nonisothermal melt-crystallization kinetics for three aromatic polyesters. Thermochem Acta 2003;406:207–20. [67] Galwey AK. Eradicating erroneous Arrhenius arithmetic. Thermochem Acta 2003;399:1–29. [68] Galwey AK. What is meant by the term ‘variable activation energy’ when applied in the kinetic analyses of solid state decompositions (crystolysis reactions)? Thermochem Acta 2003;397:249–68. [69] Sahay SS, Krishnan K. Analysis of the nonisothermal crystallization kinetics in three linear aromatic polyester systems. Thermochem Acta 2005;430:23–9. [70] Chrissafis K, Paraskevopoulos KM, Stavre SY, Docoslis A, Vassiliou A, Biakiaris DN. Characterization and thermal degradation mechanism of isotactic polypropylene/carbon black nanocomposites. Thermochem Acta 2007;465:6–17. [71] Papageorgio G, Bikiaris DN, Chrissafis K. A different approach for the study of the crystallization kinetics in polymers, key study: poly(ethylene terephthalate)/SiO2 nanocomposites. Polym Intl 2010;59:1630–8. [72] Eloussifi H, Farjas J, Roura J. Non-isothermal model free predictions: applications to the formation of yttria from yttrium trifluoroacetate. J Therm Anal Calorim 2012;108:597–603. [73] Qiu S, Zheng Y, Zeng A, Guo Y. Prediction of non-isothermal crystallization parameters for isotactic polypropylene. Thermochim Acta 2011;512:28–33.

Please cite this article as: M. Atiqullah et al., UHMW Ziegler–Natta polyethylene: Synthesis, crystallization, and melt behavior, Journal of the Taiwan Institute of Chemical Engineers (2017), http://dx.doi.org/10.1016/j.jtice.2017.04.011