Influence of Polymerization Conditions on Melting Kinetics of Low

Article pubs.acs.org/Macromolecules

Influence of Polymerization Conditions on Melting Kinetics of Low Entangled UHMWPE and Its Implications on Mechanical Properties Dario Romano,† Niek Tops,‡ Efren Andablo-Reyes,† Sara Ronca,† and Sanjay Rastogi*,†,‡ †

Department of Materials, Loughborough University, Ashby Road, Loughborough, LE11 3TU, Leicestershire, U.K. Research Institute, Teijin Aramid, Velperweg 76, Arnhem, 6802ED, The Netherlands

‡

S Supporting Information *

ABSTRACT: The synthesis of linear ultra high molecular weight polyethylene, with a “pseudoliving” catalyst in various conditions results in samples of different molecular weights (Mw ranging between 2 to 35 million g/mol), all having a reduced number of entanglements to an extent that allows the solid-state uniaxial deformation of such high molar masses without melting. The solid-state processing of these materials shows a clear relationship between mechanical properties and molecular weight. For the adopted polymerization conditions, stretching forces required for the uniaxial deformation increase with the increasing molar mass, ultimately limiting the achievable maximum draw ratio in the polymers having Mw > 10 million g/mol. The increase in the stretching force is attributed to the increasing number of entanglements between the crystals with the molar mass. The estimation of entanglements is established with the help of melting kinetics involved in the “disentangled” crystals, and rheological response of the polymer melt obtained just after melting of the crystals. In spite of the increase in the stretching forces with the increasing molar mass, tensile modulus increases with the increasing draw ratio and the molecular weight. However, above the molar mass of 10 million g/mol, the stretching force required increases to the level that the uniaxial deformation becomes difficult−thus limiting the tensile strength.

1. INTRODUCTION

One of the fundamental drawbacks of UHMWPE is the extremely high melt viscosity, which does not allow the use of the conventional processing techniques like injection molding, blow molding or extrusion, commonly used to process thermoplastics. An alternative route to process UHMWPE was introduced in the 1980s by Smith and Lemstra.7 The processing route, aimed to reduce the viscosity of the polymer by reducing the entanglements between the chains, required the dissolution of the polymer in a suitable solvent in a ratio of 5% v/v (polymer in the solvent) to produce a viscous solution that can be drawn/spun into fibers. Another adopted approach by Smith et al. was to reduce the polymerization rate below the crystallization rate during synthesis.8 Recently, it has been shown9 that using a single-site catalytic system it is possible to synthesize UHMWPE having a reduced number of entanglements, leading to a polymer with an initial melt viscosity lower than the entangled state of the same polymer. The nascent powder having the reduced number of entanglements can be processed in solid-state below its melting temperature, leading to products such as fibers and tapes of UHMWPE.10 This route to process the polymer avoids any use of solvent during processing, uses milder temperature conditions during the

Polyethylene (PE) is a generic name for a class of polymers having a very broad range of applications that vary with crystallinity and thus the associated bulk density of the material, which is strictly related to molecular architecture. Polymer properties can be varied by modifying the molar mass and the branching content of the polymer. For instance, low density polyethylene is flexible and soft, the macroscopic mechanical properties are intimately related to its highly branched microstructure, which does not allow the formation of a densely packed crystalline structure. The properties of polyethylene can be drastically modified by increasing the molecular weight above a million g/mol and removing/ reducing the branching: linear ultra high molecular weight polyethylene (UHMWPE) is used to produce fibers and tapes having unprecedented mechanical properties required for an engineering material rather than a commodity plastic.1−3 Moreover, because of its biocompatibility, high abrasion resistance, low fatigue and adhesive wear it is also used as a synthetic material for hip and knee prostheses and joint replacements.4,5 These unique mechanical properties, related to the high molar mass, make UHMWPE of great relevance and for this reason several studies have been carried out to understand the effect of synthesis conditions6 on the material’s properties. © XXXX American Chemical Society

Received: April 17, 2014 Revised: June 19, 2014

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dx.doi.org/10.1021/ma5008122 | Macromolecules XXXX, XXX, XXX−XXX

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Scheme 1. Graphical Representation of the Protocol Used for the Annealing Experiments

in combination with MAO at pressures higher than atmospheric, as the increased rate of polymerization could not be sustained by the gas control unit. A Büchi 1.6 l jacketed glass vessel and a stainless steel lid equipped with a thermometer probe, gauge, ethylene feeding pipe, two pipes for vacuum and nitrogen, injection switch and a Büchi Cyclone 075 equipped with three 4-bladed (35 mm and two 50 mm) propellers are kept under vacuum overnight at 125 °C by the means of a feedback loop control Huber Unistat 425 thermo-regulator. The reactor vessel is backfilled with nitrogen, and after at least three cycles of vacuum/ nitrogen, anhydrous toluene is introduced at room temperature. The solvent temperature in the vessel is set at 10 °C and controlled by the thermostat probe connected to the thermo-regulator. The desired amount of scavenger is introduced in the reactor vessel and the nitrogen stream is replaced by ethylene gas at 0.1 atm overpressure (1.1 atm absolute pressure). The ethylene pressure is kept at the fixed pressure by means of a gas flow meter Büchi BPC6002 and the reaction medium is stirred at 1250 rpm. After attaining the ethylene saturation at fixed temperature, the catalyst bis[N-(3-tert-butylsalicylidene)pentafluoroanilinato]titanium(IV) dichloride is dissolved in 1−2 mL of toluene in glovebox and activated with 1 mL of MAO. The polymerization is initiated with the injection of the solution and the pressure is immediately increased to the set value. After 60 min, the polymerization is stopped by addition of methanol. Independently from the polymerization procedure, the resulting polyethylene is filtered, washed with copious amount of methanol/ acetone and dried overnight at 40 °C under vacuum. An acetone solution of Irganox 1010 (0.7−1.0 wt % Irganox1010/Polymer) is added to the polymer in order to prevent oxidation during long rheological experiments. The polymer is subsequently dried in a vacuum oven at 40 °C. 2.2. Thermal Characterization Procedures. Thermal characterization is performed in a differential scanning calorimeter (DSC) Q2000 TA Instruments, where nitrogen is continuously purged at 50 mL min−1. Temperature and enthalpy calibrations are performed using certified indium at the same heating rate used in the experimental procedure. In the high precision T-zero pans an amount between 1.0 to 2.0 mg of PE samples is weighed and crimped with a T-zero lid. A heating−cooling−heating temperature ramp from 50 to 160 °C is

process, reducing the polymer degradation, and also gives better mechanical properties.10 In this study we aim to address the effect that reaction conditions (in particular reaction time and monomer pressure) have on the molecular weight capability of the catalytic system and entanglement density of the polymer produced, as these in turn will influence the mechanical properties of the material. The “quasi-living” behavior of the Bis[N-(3-tertbutylsalicylidene)pentafluoroanilinato] titanium(IV) dichloride (FI) catalytic system of Fujita11 allows tailoring the molecular characteristics of the synthesized polymer.12 For the purpose, a series of UHMWPEs are synthesized using the same catalyst precursor as in ref 10 activated with “trimethylaluminum-free” methylaluminoxane (for different reaction times)12 and methylaluminoxane (for different reaction pressures). The adopted synthesis route provides an unique opportunity to investigate the influence of molar mass on tensile strength, tensile modulus, and energy to break of the uniaxially drawn polymers (Mw between 2 to 35 million g/mol), processed in the solid state without using any solvent.

2. EXPERIMENTAL SECTION All manipulations of air and moisture-sensitive compounds are performed under a nitrogen or argon atmosphere using standard high-vacuum Schlenk techniques or in a glovebox. Ethylene (grade 3.5) is purchased from BOC. Bis[N-(3-tert-butylsalicylidene)pentafluoroanilinato]titanium(IV) dichloride catalyst is received from MCat and used as received. 2,4-di-tert-butyl-4-methylphenol (BHT) (≥99.0%), and toluene (anhydrous, 99.8% in 20 L drums) are obtained from Sigma-Aldrich. Methylaluminoxane (MAO) (10 wt % of MAO in toluene solution) is supplied from Albemarle. Irganox 1010, added as antioxidant to the polymer after the synthesis, is purchased from Ciba. 2.1. Polymerization Procedures. The polymerization procedure that uses the cocatalyst modifier at atmospheric pressure is described elsewhere.12 With increasing monomer pressure, the polymerization rate increases significantly, leading to very high yield. Because of technical limitations of the equipment used, we chose not to use BHT B


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Table 1. Ethylene Polymerization Results for Different Reaction Time/Pressuresa run d

1 2d 3e 4 5

time (min)

pressure (atm)

yield (g)

Rpb

Mwc

Mnc

MWD

Tm (°C)

crystallinity (%)18

10 30 60 60 60

1.1 1.1 1.1 2.1 4.1

16.0 32.4 63.6 51.8 90.5

9738 6761 4599 3963 3547

2.3 5.3 9.9 15.3 34.0

1.0 2.0 3.8 5.1 4.8

2.3 2.6 2.6 3.0 7.1

139.5 140.0 139.4 140.1 140.3

77 75 76 75 77

a Other conditions: Al/Ti molar ratio, 1200; catalyst concentration, 8.3 μmol/L; reaction solvent, 0.75 L of toluene; temperature, 10 °C. bCatalyst activity expressed as kgPE/molTi·atm·hour. cAs obtained from rheological characterization, in 106 g/mol. dReaction solvent, 1.0 L; Al/BHT molar ratio, 1.67. eReaction solvent, 1.5 L.

performed at 10 °C min−1. For annealing experiments, 1.500 ± 0.100 mg of PE samples are weighed in a high precision T-zero pan and crimped with a T-zero lid. A preannealing scan at an annealing temperature between 137 and 138 °C for 1 h was required to determine melting temperature. The DSC experiments are performed as follows (Scheme 1): (a−b) From 50 °C to the chosen annealing temperature following the empirical equation: Ta = Tm2p − 3.5

polymer (expected to have molar mass above 6 million g mol−1) is preannealed in a vacuum oven for 48 h at 160 °C. The annealed powder is compression molded at 160 °C. From the molded plate, discs of 12 mm are obtained. Frequency sweep tests are performed in a range of 100 to 1 × 10−3 rad/s, at a constant strain (0.3−0.5%) within the linear viscoelastic region of the polymer. Following the stress relaxation experiments described in ref 15, the frequency range was extended to values as low as 10−4 rad/s. The time dependent relaxation modulus G(t) is Fourier transformed to angular frequency space and used as complementary data for the oscillatory frequency sweeps in the low frequency range (10−3−10−4 rad/s). Using the TA Orchestrator software available on the rheometer, molecular weight and molecular weight distribution were determined. 2.4. Mechanical Characterization Procedures. Shaping of the Synthesized Powder to Uniaxially Oriented Tapes. For determination of the mechanical properties of the uniaxially deformed disentangled UHMWPE the molding of nascent powder without melting is desired. For compression molding of the powder without melting, 25 g of polymer powder is poured into a mold having cavity of 620 mm in length and 30 mm in width. The uniformly dispersed low density powder in the mold is compression-molded at 129 °C and 130 bar for 10 min. The resulting sheet of 1.42 mm thickness is preheated for at least 1 min at a constant temperature of 136 °C and rolled with a Collin calender (diameter rolls, 250 mm; slit distance, 0.15 mm; inlet speed, 0.5 m/min). While rolling (speed 2.5 m/min) the sheet is stretched partially. The rolled and stretched sheet is further stretched in two steps on a 50 cm-long, oil heated hot plate. The draw ratio is obtained by dividing the specific weight of the sheet prior to deformation by the specific weight of the tape after stretching. A typical processing temperature of polyethylene in the two stretching steps ranges between 130 to 154 °C. The higher stretching temperature, above 140 °C, is used for the partially stretched samples where the macroscopic forces can be transferred to molecular level under external constraint. To recall, melting temperature of linear uniaxially stretched UHMWPE can be increased under external constraints.16 The sample is stretched to the desired initial draw ratio in the first stretching step. Parts of the drawn sample are used to measure the mechanical properties, whereas the remainder of the sample is drawn further to the final draw ratio and the mechanical properties are determined subsequently. Determination of the Mechanical Properties. Tensile properties were measured according to ASTM D7744−2011 using an Instron 5566 tensile tester at room temperature (25 °C). To avoid any slippage, the side action grip clamps with flat jaw faces are used. The nominal gauge length of the specimen is 100 mm, and the test is performed at a constant rate of extension (crosshead travel rate) of 50 mm/min. The breaking tenacity (or tensile strength) and modulus (segment between 0.3 and 0.4 N/Tex) are determined from the force against displacement between the jaws.

(1)

where Ta is the annealing temperature and Tm2p is the melting temperature of the high temperature peak after annealing for an hour between 137 and 138 °C. (b−c) Annealing for different times (0.25, 0.50, 0.75, 1.00, 6.00, 12.00, and 24.00 h), at the set annealing temperature Ta. (c−d) Cooling from the annealing temperature to 50 °C at 10 °C min−1. (d−e) Isothermal at 50 °C for 2 min. (e−f) Heating from 50 to 160 °C at 10 °C min−1. (f−g) Isothermal at 160 °C for 2 min. (g−h) Cooling from 160 to 50 °C at 10 °C min−1. (h−i) Isothermal at 50 °C for 2 min. (i−j) Heating from 50 to 160 °C at 10 °C min−1. 2.3. Rheological Characterization Procedures. To perform the rheological studies a procedure similar to the one described in ref 13 is applied. The nascent powder is compressed into a plate of 50 mm diameter and thickness 0.6−0.7 mm at a fixed temperature of 125 °C, under a maximum force of 20 tonnes for an average time of 20 min. From the compressed plate, using a punching device, 12 mm diameter disks are cut for rheological studies. Rheological experiments are performed using a 12 mm parallel plate strain controlled rheometer TA Instruments, ARES G2 or ARES LS2. The disc is placed between parallel plates at an initial temperature of 110 °C. To prevent oxidation, the temperature is controlled by a convection oven under a nitrogen environment. After thermal stabilization at 110 °C, temperature is increased to 130 °C at 30 °C/min. After the stabilization, the sample is heated further to 160 °C at 10 °C/min while maintaining the compression force of 4 N. Subsequently, an oscillatory amplitude sweep test is carried out to determine the range of oscillatory strains in the linear viscoelastic region. The test is performed at a fixed frequency of 10 rad/s. To follow changes in the modulus of the polymer melt, during the transformation from nonequilibrium to equilibrium state, an oscillatory time sweep is performed; the test is executed at a fixed frequency of 10 rad/s and a strain between 0.1 and 0.5%, well within the linear viscoelastic regime. To determine the molecular weight of the polymer, after reaching the equilibrium storage plateau modulus, a small amplitude oscillatory frequency sweep test at a fixed strain between 0.3 and 0.5% is performed. Using the commercial TA Orchestrator software available, based on the method developed by Mead,14 molecular weight and molecular weight distribution are determined for linear polyethylene. However, polymers having molecular weight higher than 6 × 106 g/ mol are not able to achieve the thermodynamic equilibrium state. Since achieving the thermodynamic melt stable state is a necessary condition for molar mass and molar mass distribution determination, a different procedure is applied to obtain the thermodynamic melt. The

3. RESULTS AND DISCUSSION One of the unique features of the adopted synthesis in this study is that the “pseudoliving” character of the precatalyst allows tailoring of the molar mass of the polymers by changing the polymerization time, at atmospheric monomer pressure. Possible deactivation/termination processes due to the C


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Figure 1. (a) Molar masses and molecular weight distribution as a function of time for the polymers synthesized at atmospheric pressure. (b) Molar masses and molecular weight distribution as a function of pressure for the polymers synthesized for 60 min. Black filled squares, Mw; black filled circles, Mn; red filled up-triangles, molecular weight distribution. The lines have been drawn for visual guidance to the data points.

most-likely chain termination process is a β-H transfer to the monomer, at least in the conditions used. The increase in the weight-average molecular weight on increasing the monomer pressure could be related with the formation of longer chains promoted by a secondary catalytic specie generated in the presence of trimethylaluminum,12 that could also contribute to the broadening of the molecular weight distribution. At lower monomer pressure, the secondary species (possibly slightly faster in polymerization rate, but lower in concentration compared to the main catalytic species) could be limited by mass transfer limitation of the monomer to the active site, by the main catalytic specie. Thus, on increasing the monomer pressure, and thereby enhancing the monomer dissolution in the reaction medium, a possible mass transfer limitation is likely to be reduced, and the secondary active specie may start to contribute more to the overall polymerization. It should be pointed out that values of molecular weight exceeding 107 g/mol are affected by high inaccuracy arising from the rheological method adopted for molar mass determination. The accuracy in the molecular weight and molecular weight distribution estimation is based on the determination of the crossover point of the viscoelastic functions (G′ and G″), i.e., on the response of the viscoelastic functions at low frequencies. For molecular weights that exceed 107 g/mol, the crossover point is likely to be located at very low frequencies, below 10−5 rad/s. Following changes in the viscoelastic response at these low frequencies is a challenging task that cannot be achieved by linear oscillatory rheometry in the commercially available rheometers. Beside the limit that the force transducers will encounter to measure changes in the loss and storage moduli, the time required to collect a data point becomes extremely long. To overcome this problem, stressrelaxation experiments are performed, and the time is converted into frequency.15 Some of the frequency sweep and stress relaxation data shown in the supplementary section: Figures S2, S6, and S7, in the supplementary section have been used for estimating the molar mass and molar mass distribution. The comparison of Figures S2 and S6 clearly demonstrate that the

presence of free trimethylaluminum in equilibrium with the MAO cocatalyst are considerably reduced with the addition of a sterically hindered phenol (BHT) in combination with the catalytic system (FI/MAO). The presence of BHT causes reduction in the polydispersity of the synthesized polymer, as shown elsewhere.12 3.1. Influence of Polymerization Conditions on Molecular Characteristics. Table 1 summarizes some of the molecular characteristics of the polymers synthesized at atmospheric pressure for different polymerization times, and at fixed time while changing the monomer pressure. Molar mass and molar mass distribution have been estimated using the rheological studies,15 crystallinity and melting temperatures have been determined with the help of DSC.17 In accordance with the earlier findings,12 the catalyst activity decreases with increasing the polymerization time. The catalyst activity slightly decreases also on increasing the ethylene pressure, a result that is rather counterintuitive, giving that the polymerization rate should increase with increasing the monomer concentration. The decrease in the catalyst activity is likely to be a result of slow diffusion of the ethylene monomer through the fast growing polymer chain.19 Figure 1a depicts the increase in the weight (Mw) and number-average (Mn) molar mass and molecular weight distribution (MWD) as a function of polymerization time. Relatively low polydispersity for these high molar masses is in the expected range of the polymers synthesized using a single site catalytic system (2−4). The molecular weights (both Mw and Mn) and molecular weight distribution data of the polymers synthesized at higher monomer pressure are summarized in Figure 1b. From Figure 1b, it is apparent that the Mw increases significantly with increasing the monomer pressure, while the Mn increases from 3.8 to 5.1 × 106 g/mol when the monomer pressure is increased from 1.1 to 2.1 atm, and does not change significantly if the pressure is further increased to 4.1 atm. Since the Mn of the polymer produced does not increase above ∼5.0 × 106 g/ mol, on increasing the monomer pressure, we can infer that the D


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Figure 2. Evolution of the normalized area of DSC peaks as a function of annealing time. Samples used for the study are summarized in Table 1. Black-filled symbols represent polymers synthesized for different polymerization time; black unfilled symbols represent polymers synthesized at different pressure. Solid line represents commercial (entangled) Sigma-Aldrich UHMWPE, shown for comparison. The arrow indicates the direction of increasing Mw and Mn of the polymers synthesized in our lab.

the surface.22 Following this existing knowhow,9,17,23 characterization of entanglements in the synthesized UHMWPE is performed and linked to ease in the solid state deformation as well as mechanical properties of the polymers. The protocol used for these studies is discussed in section 2.2 and shown in Scheme 1. The annealing time shown in Figure 2 refers to the annealing in the region b−c of Scheme 1. The annealing temperature is determined following the described protocol depicted in eq 1. Depending on the annealing time, on first heating from e to f, two melting peaks (approximately at 135 and 142 °C) instead of one are observed. In a disentangled polyethylene crystal, with the increasing annealing time the area of the low temperature peak increases at the expense of the high temperature peak. The low melting temperature peak is attributed to the melt-crystallized component obtained on cooling from c to d after the annealing protocol b−c. The high melting temperature peak is attributed to the remainder of the nascent crystal. For more details on this analysis please refer to reference by Pandey et al.23 The area of the low temperature melting peak, normalized by the total area of the two peaks, gives an indication on the distribution of entanglements in the amorphous region of the semicrystalline polymer as obtained from the reactor. The initial slope and the highest value reached at longer annealing time are related to the entanglement density of the polymer. Annealing experiments depicting the entangled state of the polymers from Table 1 are summarized in Figure 2. All the investigated polymers show increase in the normalized area as a function of annealing time indicating the synthesis of polymers having sufficiently low degree of entanglements between the crystals.9 With the increasing polymerization time (and consequently the molar mass) the slope of the growth rate of the normalized area decreases. The slope depicts the disentangled state of the crystals obtained during synthesis: in principle, considering the

samples synthesized at higher pressure have the crossover point at much lower frequencies, and have broad frequency region where the storage plateau modulus remains constant. The plateau of this frequency region increases with the increasing polymerization pressure, from 1.1 to 4.1 atm, which is indicative of the increase in Mw. One of the requisites for the solid state deformation of the UHMWPE is the reduction of entanglements residing in the amorphous region of the semicrystalline polymer.20 It is conclusively demonstrated that the solid state processing requires substantial amount of adjacent chain re-entry, thus reducing the entanglement density.20 The reduction in entangled state is achieved either by dissolution and subsequent recrystallization of the polymer7 or by controlled synthesis using a single-site catalytic system12 or heterogeneous synthesis at very low temperatures.8 Because of the near absence of structural coherence in the amorphous region, the quantification of entanglements is challenging. However, recently with the help of DSC we have shown that is possible to estimate the influence of entanglements in the solid state under quiescent conditions.17 To recall, the nascent powder of ultra high molecular weight polyethylene shows a peak melting temperature close to the equilibrium value of 141.2 °C.21 Following the Gibbs− Thomson equation such high temperature should be correlated to extended chain crystals, i.e., crystal thickness above 100 nm. However, the observed melting temperature could not explain the chain folded morphology of nascent UHMWPE having crystal thickness of ca. 20 nm. 12 To overcome this incongruence in melting temperature and the crystal thickness, the influence of amorphous region on melting behavior is invoked. It has been shown that melting kinetics, arising on annealing of the crystals below the equilibrium temperature, causes melting with the consecutive detachment of chains from E


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Figure 3. Schematic representation of the crystals during the DSC annealing experiments. Melting during annealing (below Tm) occurs via chain detachment from the crystal surface.

concept of a monomolecular crystal, for the crystals having the same molar mass, the steeper the slope, the more disentangled the amorphous region will be.22,23 At first instance these findings may appear to be in contrast with ref 6 and ref 23, where it is argued that the entanglement density decreases with the increasing polymerization time, or molar mass. However, in ref 23, the growth rate of the normalized area has been compared between samples of similar molar masses. In our case the comparison on the slope of the growth rate of the normalized area has been made between molar masses that differ by more than a factor of 2 or 4. If we consider the simplistic model of successive chain detachment from the crystal surface (enthalpic relaxation process), the time required for the normalized area to reach the plateau value for a crystal having the highest molar mass will be considerably larger than the crystal of low molar mass. Important to note is that ultimately all polymers reach the plateau value of 1, though the highest molar mass takes the longest time. The normalized value reaching 1 conclusively shows that, if given enough time, the disentangled polymer, independent of its molar mass, will melt at temperatures lower than that predicted by normal heating runs. These findings conclusively demonstrate the kinetic aspect of polymer melting, as the melting process is strongly influenced by the chain topology in the amorphous region rather than the crystalline domain. In addition, it is evident that synthetic conditions, such as polymerization pressure (that also influences heat dissipation during polymerization) and polymerization time, both have a strong influence on tailoring the entangled state in the amorphous region of the semicrystalline polymer and the weight-average molecular weight/distribution. The DSC experiments support the idea that the kinetics of the enthalpic relaxation process on

annealing strongly depends on the molar mass and the constrained state of the chain segments in the amorphous region of the nascent semicrystalline polymer. The SEM image in the top left of Figure 3 exhibits chain folded nascent crystals of UHMWPE. These crystals, obtained from the nascent powder out of the reactor, show a crystal thickness of approximately 12 nm. On heating at 10 °C/min these crystals show a melting temperature peak at 141.5 °C (approximately). On annealing, below the melting temperature (shown in Figure 2), crystals tend to melt by chain detachment from crystal surface (ac or bc crystallographic plane) and reeling into the polymer melt (see Figure 3). Specifics of the mentioned hypothesis have been summarized in the supplementary section (Table S1 and Figures S8−S11) and can be followed in more detail in the references 13, 17, 22, and 23. On cooling to 50 °C the sample contains two populations of crystals, melt crystallized and the remaining nascent crystals. On subsequent heating, at 10 °C/min, the melt-crystallized component melts at 135 °C (approximately) whereas the remaining crystals at 141.5 °C (approximately). With the increasing molar mass, the surface area (ab plane) of the crystals is likely to increase; leading to an increase in the number of entanglements per crystal. Consequence to it, with the increasing molar mass, the time required to melt the crystals by annealing process also increases (Figure 2). On the other hand, if the number of entanglements between the crystals is very large, kinetics involved in the enthalpic relaxation process by subsequent detachment of chains from the surface and their reeling into the melt will be arrested. This will suppress the melting kinetics as observed in the commercial sample (Figure 2). The increasing number of entanglements between the crystals will also have implications on the strain-hardening slope F


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calculated as the force required to stretch the sample normalized by the cross-sectional area for each draw ratio. For all the polymers processed in this publication, the total drawing tension grows quasi linearly as a function of draw ratio as confirmed by the values of coefficient of determination (R2) close to 1. The total drawing tension increases as a function of the molar mass for all the polymers synthesized. For high draw ratio, the total drawing tension ranges between 400 and 1400 N/mm2. Decreasing the hot plate temperature, during the processing, the total drawing tension increases denoting more difficulty in the chain alignment. A common feature for all the polymers investigated is the relatively low stretching force requirement. It is also important to notice that after yielding with the neck formation on the hot plate, the applied force for the specific draw ratio remains constant until a specific draw ratio for the applied force is achieved.24 Figure 5b summarizes the stretching force for the polymers of Table 1 at comparable draw ratio (∼180) as a function of molar masses. It can be conclusively stated that the stretching force increases with the increasing molar mass. In particular, for small changes in the molecular weights the effect on the stretching force is more pronounced for the Mn compared to Mw. Figure 6a shows the breaking tenacity as a function of draw ratio for the polymers of Table 1. The breaking tenacity increases as a function of draw ratio for all the polymers synthesized. High values of breaking tenacity (over 3.6 N/Tex) are found for the polymers stretched to high draw ratios. The breaking tenacity obtained for all the polymers are considerably above the values that can be obtained using a commercial entangled UHMWPE. For run 2, on lowering the hot plate temperature from 153 to 140 °C, no significant changes in the values of breaking tenacity are found. It has to be noticed that the samples of run 1 and run 2 can be stretched beyond the draw ratio of 200, mainly due to the low drawing tension (Figure 5a). The higher draw ratios accessible for the samples run 1 and run 2 provide unprecedented breaking tenacity and consequence to it, high energy to break (Figure 8a). Figure 6b summarizes the breaking tenacity for the polymers of Table 1 at a comparable draw ratio (∼180) as a function of molar masses. There are no clear trends in the evolution of the breaking tenacity. It seems that the breaking tenacity increases as a function of molar mass, for the first three samples, and then decreases. It should be noticed that for the polymers synthesized at higher monomer pressure, the molecular weight distribution broadens possibly due to the synthesis of new chains arising from the chain termination to the monomer and the presence of a secondary active species. The new chains produced from the chain termination (and thus of lower molecular weights) will reduce the maximum breaking tenacity achievable. Figure 7a summarizes the Young’s modulus as a function of draw ratio. The modulus increases with the draw ratio and reaches value in the range of ∼175−210 N/tex. The values of modulus of the polymers synthesized are slightly below (for the polymers synthesized in runs 1, 2, and 3) and above (for the polymers synthesized in run 4 and run 5) compared to the modulus obtained using commercial entangled UHMWPE. To recall, the polydispersity of the commercial sample is much higher than the lab synthesized samples, which implies that the tensile modulus is lesser influenced by Mn compared to the tensile strength to break (Figure 6).

in the stress−strain curve causing the requirement of higher stretching forces for the same draw ratio. This has been depicted for the increasing molar mass in the samples synthesized at higher ethylene pressures (anticipated increase in the surface area of the ab plane and thus the associated number of entanglements) in the following section. From Figure 2, taking the above hypothesis into consideration, it can be stated that the sample of run 5 will take very long time to reach the plateau value. The long time required for the enthalpic relaxation process can be attributed to the entangled state of the amorphous region. This becomes evident on comparing the storage modulus recorded just after melting of the crystals of runs 3, 4, and 5. The sample obtained from run 3 (1.1 atm) shows the lowest starting storage modulus compared to the sample obtained from run 5 (4.1 atm). Higher storage modulus suggests lower molecular weight between entanglements (Me) corresponding to greater number of entanglements in the initial stages of melt. With the chain reptation, as the number of entanglements increases, the Me decreases and the storage modulus continues to increase. The time required for the storage modulus to reach the equilibrium state depends on the molar mass and the molar mass distribution.6,12 In summary, the higher initial storage modulus for run 5 (Figure 4) suggests the presence of higher number of

Figure 4. Influence of ethylene pressure on entangled state between the nascent crystals. The entangled state in the semicrystalline polymer is depicted from the initial storage modulus of the polymer melt of run 5 compared to run 3 and run 4. The data has been recorded at 160 °C and fixed strain in the linear viscoelastic region.

entanglements in the solid state. The presence of the higher entanglement density in the run 5 sample can be attributed to the higher ethylene consumption for the same polymerization time (Table 1), causing increase in the temperature in the localized region due to the exothermic nature of the ethylene polymerization thus influencing the crystallization rate and the associated entanglement density. To recall, one of the requisites for the synthesis of disentangled crystals is that the polymerization kinetics should be slower than the crystallization kinetics.6,8 3.2. Influence of Molar Mass on Mechanical Properties. All the polymers, run 1 to run 5, can be deformed uniaxially at least up to a draw ratio of 190, which confirms the ease in uniaxial deformation of the compressed disentangled UHMWPE powder. Figure 5a shows the total drawing tension as a function of draw ratio. The total drawing tension is G


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Figure 5. (a) Total drawing tension as a function of draw ratio, filled symbols represent polymers synthesized at different time, unfilled symbols represent polymers synthesized at different pressure. For run 2, as anticipated, the stretching force increases on decreasing the temperature from 153 to 140 °C. (b) Stretching force as a function of molecular weights at a fixed draw ratio of 180 (filled squares, Mw; unfilled squares, Mn). For details of the run number see Table 1.

Figure 6. (a) Breaking tenacity as a function of draw ratio; filled symbols represent polymers synthesized at different time, unfilled symbols represent polymers synthesized at different pressure. Solid curve represent the values of breaking tenacity of the commercial entangled UHMWPE, taken from ref 10. (b) Breaking tenacity as a function of molecular weights at a fixed draw ratio of 180 (filled squares, Mw; unfilled squares, Mn). It has to be noticed that the samples of run 1 and run 2 can be stretched beyond the draw ratio of 200, mainly due to the low drawing tension (Figure 5a).

Figure 7b summarizes the values of modulus at a fixed draw ratio (∼180) as a function of molar masses. For the same draw ratio, the modulus increases with the increasing molar mass. Figure 8a shows the energy at break for the polymers synthesized. The energy at break corresponds to the area under the stress−strain curve measured by the extensometer during

mechanical testing at break. As common feature, the energy at break increases as a function of the draw ratio for all the polymers. At high draw ratios (>150), the values of energy at break range from 45 to 60 J/g. No significant differences are found on lowering the stretching temperature from 153 to 140 °C, during stretching for the sample run 2a and run 2b. Higher H


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Figure 7. (a) Modulus as a function of draw ratio; filled symbols represent polymers synthesized at different time, unfilled symbols represent polymers synthesized at different pressure. Solid curve represents the values of modulus of the commercial entangled UHMWPE, taken from ref 10. (b) Modulus as a function of molecular weights at a fixed draw ratio of 180 (filled squares, Mw; unfilled squares, Mn).

Figure 8. (a) Energy at break as a function of draw ratio; filled symbols represent polymers synthesized at different time, unfilled symbols represent polymers synthesized at different pressure. (b) Energy at break as a function of molecular weights at fixed draw ratio of 180 (filled squares, Mw; unfilled squares, Mn).

molecular weight. It seems that the energy at break has a peak corresponding to the run 3. Similar to the breaking tenacity, the presence of low molecular weight chains may lower the energy required to break the specimens. Coupling the mechanical properties with the annealing characterizations, it becomes evident that the samples that reach the plateau value of 1 in the time frame of DSC experiments (Figure 2) can be stretched uniaxially to the

energy at break depicted in Figure 8 implies greater area under the stress−strain curve, having implications in the construction of high impact articles, such as body armor. Figure 8b shows the energy at break at a fixed draw ratio (∼180) as a function of molecular weight for the polymers summarized in Table 1. Similar to the dependence of the breaking tenacity as a function of molar masses (Figure 6b), the energy at break does not show clear dependence with the I


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Overlay of the ultimate mechanical properties at high draw ratio show different tendency for the molar masses (Mw and Mn). Small increase in Mn, compared to Mw, show some rise in the stretching force and tensile modulus, whereas the tensile strength is strongly influenced by the increase in Mn.

ultimate properties, with low deformation forces. On the other side, the sample that did not reach the normalized area value of 1 falls below the ultimate achievable tensile strength and does require higher stretching forces (for example the polymer synthesized at 4.1 atm). It should be noted that, unlike the tensile modulus, the tensile strength to break is strongly influenced by the entanglement density, as well as the molar mass. We can infer from these data that the ultimate tensile strength combined with tensile modulus can be achieved by making a disentangled linear ultra high molecular weight polyethylene having Mn > 5 million g/mol and avoiding the low molar mass component. Beside the difference in the entanglement density of the polymers synthesized, another possible explanation of the observed difference in the dependence of some mechanical properties, as a function of the different molar masses, could be ascribed to the molar mass and molar mass distribution. For example, tensile modulus is found to be more influenced by the weight-average molar mass (Mw), whereas tensile strength is found to be more influenced by the number-average molar mass, i.e., the number of chain ends. This will be further addressed in a following publication where the polymers have been synthesized at atmospheric pressure while the polymerization time is varied, thus removing the influence of entanglement density with the increasing ethylene pressure or influencing the polymerization kinetics.

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ASSOCIATED CONTENT

S Supporting Information *

Catalyst activity as a function of polymerization time, evolution of the storage modulus as a function of time, catalyst activity as a function of ethylene pressure, ethylene uptake profile, reaction temperature profiles, viscoelastic moduli, stress relaxation modulus as a function of time, molecular weight data, DSC studies, and NMR spectra. This material is available free of charge via the Internet at http://pubs.acs.org.

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AUTHOR INFORMATION

Corresponding Author

*(S.R.) E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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REFERENCES

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CONCLUSIONS We have reported the synthesis of a set of disentangled UHMWPE at different polymerization times and different monomer pressures. Molar masses and entanglement density have been characterized using rheology and DSC. The influence of molecular characteristics (Mw and Mn) on the mechanical properties of UHMWPE processed in the solidstate has been addressed. Polymers synthesized at lower polymerization times (low molar masses) require lower stretching force during solid-state processing. The set of polymers synthesized at higher monomer pressure show an increase in Mw and almost constant Mn. This finding suggests a possible prevalence of chain termination to the incoming monomer compared with other possible termination processes. Increase in ethylene pressure results in very high molar mass and higher entanglement density, possibly due to the increase in temperature during polymerization. Increase in the stretching force with the increase in molar mass suggests increase in the number of entanglements between crystals thus restricting the solid-state deformation. The tensile strength decreases above a weight-average molar mass of 10 million g/mol while the modulus and the force required for uniaxial stretching increases. The lower values of breaking tenacity and energy at break achieved with the polymer synthesized at 4.1 atm are attributed to higher entanglement density and/or the presence of lower molecular weight chains. Kinetics involved in melting, while annealing the low entangled UHMWPE samples, provide information on the entanglement density in the amorphous phase, which in turn is related to the uniaxial drawability and mechanical properties. In summary, during annealing in DSC below the peak melting temperature, the samples that reach the plateau value of one (considering the normalized area of the low melting peak) can be uniaxially stretched to provide the ultimate mechanical properties, whereas the samples that do not reach the unitary value show lower mechanical properties. J


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Influence of Polymerization Conditions on Melting Kinetics of Low

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