Correlation between Thermal and Mechanical ... - ACS Publications

Article pubs.acs.org/Macromolecules

Correlation between Thermal and Mechanical Response of Nascent Semicrystalline UHMWPEs Dario Romano,†,‡ Niek Tops,§ Johan Bos,§ and Sanjay Rastogi*,†,‡,§ †

Department of Materials, Loughborough University, Ashby Road, Loughborough LE11 3TU, U.K. Department of Biobased Materials, Faculty of Humanities and Sciences, Maastricht University, Brightlands Chemelot Campus, Urmonderbaan 22, 6167RD Geleen, The Netherlands § Research Institute, Teijin Aramid, Velperweg 76, 6824BM Arnhem, The Netherlands ‡

ABSTRACT: In spite of the abstract nature, entanglements are the folklore of polymer science, and their presence is realized in the semicrystalline and amorphous solid state as well as in the molten state (melt) of polymers. For example, the entangled state in the noncrystalline region is known to be the determining factor in melting kinetics and also in uniaxial as well as biaxial deformation of linear ultrahigh molecular weight polyethylenes (UHMWPEs). In this paper we aim to investigate a possible correlation between the melting of crystals and the ease of deformation of the semicrystalline polymer. Considering that the entanglement density, and the associated topological constraints residing in the noncrystalline region, can be influenced by the polymerization conditions, a series of UHMWPEs have been synthesized. For a set of polymers synthesized at fixed pressure, the characteristic melting time of the as-synthesized nascent crystals (τ1) measured by temperature-modulated DSC (TM-DSC) is found to decrease with an increase in molar mass. The decrease in τ1 with molar mass is indicative of the decrease in the entanglement density. The characteristic melting time of the same set of samples crystallized from their melt state (τ2) is considered to be the reference point, viz. entangled crystals. The characteristic melting time ratio between the two characteristic times (τ2/τ1) is found to increase linearly with molar mass. The linear increase suggests an increasing deviation between the entangled states of the as-polymerized and melt-crystallized samples. The adopted approach helps in estimating the entangled fraction created during polymerization. The drawing tension of these polymers, in their nascent semicrystalline state, increases linearly with increasing draw ratio. The slope defines the ease in processing, processability index PI. The index is found to increase with the ratio τ2/τ1. The relationship between the characteristic melting time ratio and the ease in solid state processing indicates the common role of the entangled state TM-DSC is shown to be a convenient tool to establish the processability of UHMWPE in the solid state. TM-DSC measurements focus on the intracrystalline topology (arrangement of methylene units in the noncrystalline region) whereas in the total drawing process, further along the lines, other factors play a role like intercrystalline entanglements and intermolecular van der Waals forces (slippage of chains). Consequently TM-DSC assesses the indispensable intracrystalline topology required for solid-state processing of intractable UHMWPE.

1. INTRODUCTION

example by dissolving approximately 10 wt % of the polymer into a solvent (e.g., decalin or xylene), the polymer can be solution-spun into fibers having unparalleled high tensile strength and tensile modulus, making it the strongest manmade fiber.5−8 Instead of these inflammable, volatile, and toxic

Ultrahigh molecular weight polyethylene (UHMWPE) is a polyolefin possessing a weight-average molar mass (Mw) greater than 1 million g/mol, with excellent properties required for several demanding applications ranging from prostheses to lightweight body armors.1−3 Because of its high melt viscosity, the polymer cannot be processed into products via conventional methods (i.e., injection molding, blow molding, extrusion, etc.). By reducing the entanglement density,4 for © XXXX American Chemical Society

Received: October 28, 2016 Revised: February 4, 2017

A

DOI: 10.1021/acs.macromol.6b02339 Macromolecules XXXX, XXX, XXX−XXX

Article

Macromolecules solvents, paraffin oil is used in the past9 to dissolve the polymer but with the disadvantage that it is rather cumbersome to remove the oil from the as-spun fibers afterward. Recently, it is shown that the polymer up to 20 wt % can be spun into fibers using more eco-friendly solvents such as vegetable oils.10 Parallel to the solution-spinning routes, solid-state processing of the UHMWPE has been also a subject of interest.11,12 Independent of the processing routes, the basic concept involved in the processing of the intractable UHMWPE evolves from the reduction of entanglement network that resides in the inter- and intra-noncrystalline regions of the semicrystalline polymer. Recently, it has been shown that making use of singlesite catalytic systems it is possible to reduce the number of entanglements in the noncrystalline region of the semicrystalline polymer.13,14 The basic principle is the fast crystallization rate over the polymerization rate.15,16 As a result of the reduced entanglement density, the as-synthesized nascent powder can be processed uniaxially as well as biaxially in the solid-state (below the equilibrium melting temperature of linear polyethylene) without using any solvent. The solid-state processing of the uniaxially drawn tapes, having unprecedented mechanical properties, is economically as well as environmentally attractive over any process using solvents Moreover, tape over fiber geometry is preferred in manufacturing of composites as their performance is enhanced because of lower surface to volume ratio of tape geometry over the conventionally used fiber geometry in terms of packing density. The lower surface area requires lesser amount of adhesive retrospectively enhances the performance of the composite.13 Considering that the tensile strength and tensile modulus of the tapes produced from the consolidated reactor powders are influenced by the draw ratio of the synthesized polymers, which are found to be dependent on the polymerization conditions, it is of interest to have morphological insight on the entanglement network (topological constraints between the crystalline domains).17,18 For the purpose, advanced analytical tools such as solid-state NMR have been used.13,19 The kinetics involved in the melting of the crystals and correlation with the topological constraints has been established recently. The time required for melting of the crystals has been demonstrated as a tool for differentiating topological constraints; for this purpose temperature-modulated DSC (TM-DSC) has been used.18 In this publication we aim to establish a possible correlation between the melting kinetics of the synthesized UHMW-PE reactor powders in relation with the polymerization conditions and the resultant mechanical properties of the drawn tapes. In fact, the approach aims to establish a correlation between the melting kinetics of the crystals in the reactor powder and the ease of deformation. Both the thermal and mechanical properties are found to be influenced by the entangled state in the noncrystalline region of the polymer, which is proven to be dependent on the polymerization conditions. The adopted approach also could help in establishing the entangled fraction in the synthesized polymers. What follows is a brief overview on the anticipated thermal response of crystals subjected to thermal modulation. Melting Kinetics and Thermal Modulation. Different than inorganic and organic molecules, in semicrystalline polymers the melting process is far more complicated. The complication arises because of the translation of chains between crystalline and noncrystalline regions of the same or different crystals. Morphological variations in the noncrystalline region arise because of the crystallization kinetics, molecular weight,

thermal history, and molecular conformation. In its utmost simplicity, chains in the noncrystalline region can be marked to have adjacent20 or nonadjacent re-entry,21 leading to differences in topological constraints and differences in folding. In the past, using conventional differential scanning calorimetry (DSC), studies have been performed to investigate the superheating effect on the melting of nascent entangled, nascent disentangled and melt-crystallized polyethylene. In addition, Toda et al. have successfully shown the efficacy of TM-DSC in studying the melting behavior of polymer crystals.22,23 The basic concept of a TM-DSC experiment is the application of a periodic temperature modulation in addition to a linear heating rate. The resulting average heating rate gives the same information as conventional DSC experiments, while the modulation is used to determine the fraction of the total heat flow that responds to the change in heating rate. This fraction is also called reversing heat flow. The heat flow that does not respond to the modulation in temperature is called the nonreversing heat flow. It is believed that when there is a high response in the reversing heat flow, the material is very sensitive to the modulation. During melting this corresponds to the ease in melting of a crystal, also called the characteristic melting time (τ). This value characterizes the kinetic of the melting transition of the crystallites, which can be derived from the imaginary part of the complex dynamic heat capacity (ΔC) given by the relation ͠ e−iα ≅ Cp + ΔC

−F ̅ /β 1 + iωτ(β)

(1)

where Cp is the real part of the complex heat capacity, F̅ is the underlying heat flow, β is the linear heating rate, ω is the frequency of the temperature modulation, and τ is the characteristic melting time. For instantaneous transitions (ωτ ≪ 1) the real part of the heat capacity is used because the transition follows the change in temperature without any delay, and the heat flow is the same as the heat flow with true heat capacity. For slow transitions (ωτ ≫ 1) the response is controlled by the transition rate which is a function of temperature and is out of phase from the time derivative of the temperature modulation and therefore appears in the imaginary part of the heat capacity equation. The melt transition in semicrystalline polymers is relatively slow, and therefore only the imaginary part is used. Equation 1 can be simplified to ͠ ″ ≅ − F̅ P , ΔC 2πβτ

where 1/ω = P /2π

(2)

The characteristic melting time τ can now be determined from the slope by the linear fitting of ͠ ″ 1 ΔC P(period) ≅ 2πτ |F ̅ / β |

(3)

In the past, the entropic barrier associated with melting of crystals in several polymers has been invoked.24 Recent studies also confirm the role of entanglements in the melting process of polymers; a conclusion is that polymers with more entanglements need more time to melt because of the constraints imposed by the interchain links with other crystals. This results in higher characteristic melting times.25 In this publication, especially designed TM-DSC experiments are used to follow the effect of molecular weight and chain topology in the noncrystalline region on melting behavior. The melting response of crystals is correlated with the ease in solidB


Article

Macromolecules state uniaxial deformation of nascent disentangled, nascent entangled, and melt-crystallized samples of UHMWPE. For the study, using a “pseudoliving” single site catalyst, a series of disentangled UHMWPE samples have been synthesized. The polymers having increasing molecular weights are synthesized by varying the polymerization time and monomer pressure.

Scheme 1. Graphical Representation of the TM-DSC Protocol Used for Studying the Characteristic Melting Time

2. EXPERIMENTAL SECTION 2.1. Materials. 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. Toluene (anhydrous, 99.8%) is obtained from Sigma-Aldrich and used as received. Methylaluminoxane (MAO, 10 wt % of MAO in toluene solution) is supplied from Albemarle and used as received. Irganox 1010, added as an antioxidant to the polymer after the synthesis, is purchased from Ciba. 2.2. Synthesis Procedure. 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 300 equipped with 4 × 50 mm bladedpropellers are kept under vacuum overnight at 125 °C by the means of a feedback loop control Huber Unistat 425 thermoregulator. 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 temperature probe connected to the thermoregulator. The desired amount of scavenger is introduced in the reactor vessel, and the nitrogen stream is replaced by ethylene gas at 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, in a glovebox 5 mg of the catalyst is dissolved and activated with a solution of 1−2 mL of toluene and 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 the desired time, the polymerization is stopped by addition of methanol. 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 DSC and rheological experiments. The polymer is subsequently dried in a vacuum oven at 40 °C until constant weight. 2.3. Differential Scanning Calorimetry Procedures. All experiments are performed using a TA Q2000 differential scanning calorimeter. Polyethylene samples are weighted in low mass aluminum Tzero pans and lids on a XS3DU Mettler Toledo precision balance (sensitivity of ±0.001 mg). 0.250 ± 0.01 mg of sample is used for performing conventional DSC experiments. To avoid any thermal oxidation the experiments are conducted under a nitrogen atmosphere with a flow rate of 50 mL/ min. A heating−cooling−heating temperature ramp from 50 to 180 °C is performed at a linear rate of 10 °C/min. The heat of fusion and peak melting temperature are determined by integrating the melting peak from 100 to 160 °C using the sigmoidal horizontal baseline integration option in the universal analysis 2000 software. For TM-DSC, approximately 0.700 ± 0.01 mg of samples is measured under a nitrogen atmosphere with a flow rate of 40 mL/min. The temperature protocol is schematically shown in Scheme 1. The step details of the protocol used are given below: From A to B: The sample is annealed for 30 min at 110 °C and subsequently heated to 150 °C with a heating rate of 0.8 °C/min. To the heating ramp a sinusoidal temperature modulation is applied with periods ranging from 10 to 100 s. The maximum amplitude used is ±0.2 °C (heat only mode). From B to C: The sample is heated from 150 to 155 °C with a linear heating rate of 5 °C/min without modulation.

From C to D: The sample is cooled from 155 to 130 °C with a linear cooling rate of 5 °C/min. From D to E: The sample is cooled further to 100 °C with a linear cooling rate of 1 °C/min. From E to F: To examine the melting response of melt-crystallized crystals, the sample is heated for the second time to 150 °C with a heating rate of 0.8 °C/min. A heat only sinusoidal temperature modulation is applied with periods ranging from 10 to 100 s. From F to G: The sample is cooled from 150 to 100 °C with a linear cooling rate of 5 °C/min. From G to H: The sample is heated for 30 min at 100 °C without any modulation. From H to I: At 100 °C a modulation with amplitude of 0.2 °C is applied for 15 min with a period of 100 s. Data collected in this step were used for baseline correction. 2.4. Rheological Procedure. The rheological studies to follow the modulus buildup during transformation of a nonequilibrium melt to the equilibrium state are performed following the methods described elsewhere. The polymer after reaching the equilibrium state is used for determining the molar mass and molar mass distribution as has been described in our earlier publications.26,27 To recall, frequency sweep experiments are performed where the low frequency region is accessed by performing stress-relaxation experiments. The stress-relaxation experiments were desired especially for the polymer synthesized at high pressures and longer polymerization times, where the molar mass is expected to exceed 9 million g/mol. 2.5. Processing and Mechanical Characterization Procedures. Shaping of the Synthesized Powder to Uniaxially Oriented Tapes. The mechanical properties of the uniaxially deformed (dis)entangled UHMWPE samples are performed on specimens processed without melting the sample. For this purpose, 25 g of polymer powder is uniformly dispersed into a mold having a cavity of 620 mm in length and 30 mm in width. The mold containing the nascent powder 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 partially stretched. The rolled 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 and 154 °C. The higher stretching temperature, above 140 °C, is used for the partially stretched samples, at higher draw ratios, where the macroscopic forces can be transferred to molecular level under external constraint. To recall, the melting temperature of linear uniaxially stretched UHMWPE can be increased under external constraints.28−30 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 C


Article

Macromolecules sample is drawn further to the final draw ratio and the mechanical properties are determined subsequently. Determination of the Mechanical Properties. Tensile properties are measured according to ASTM D7744−2011 using an Instron 5566 tensile tester at room temperature. 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 (cross-head 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.

melt-crystallized sample during the second heating run is found to be below 50%. 3.1. Melting Kinetics TM-DSC. In Table 1 the values of the characteristic melting times of the nascent (τ1) and melt crystallized (τ2) samples are also summarized. The methods, together with the heating and cooling runs, used for the nascent and the melt-crystallized samples are given in section 2.3. The ratio between τ2/τ1, for the polymers together with their Mw, is reported in Table 1. The ratio represents change in the enthalpic relaxation of the crystals crystallized from the molten state of the nascent disentangled crystals. Considering similar crystal thickness,33 the remarkable difference in the two characteristic times is attributed to the demarcation in the entanglement network established after melting of the nascent crystals. These differences are in agreement with our earlier reported findings, probed by solid-state NMR, on chain mobility of the methylene segments in the noncrystalline region of the nascent and melt-crystallized sample.34 Important to note is that for runs 1−6 the characteristic melting time of the nascent samples decreases with the increasing molar mass, suggesting decrease in the entanglement density with polymerization time. This observation is in agreement with our earlier study, using the same catalytic system, on rheological response of these polymers where the initial storage modulus (at t = 0) of the nonequilibrium polymer melt is the lowest for the highest polymerization time (or molar mass). To recall, Figure 1, which has been reproduced from ref 16, depicts reduction in the entanglement density with increasing polymerization time, i.e., increasing molar mass. The reduction in entanglement density is attributed to the suppression in nucleation barrier after crystallization of the chain segments synthesized in the initial stages of polymerization. Figure 1, on the right, shows the increase in the storage modulus G′ with increasing molecular weight at a fixed temperature and frequency. The time required for the modulus to reach the equilibrium state increases with the increasing polymerization time or molar mass. The initial value of the storage modulus (at t = 0) is also found to be dependent on the polymerization time. The higher value of the initial modulus is attributed to the higher entangled state. The entanglement density is found to decrease with increasing polymerization time, which is likely to happen after nucleation of chain

3. RESULTS AND DISCUSSION The results of the ethylene polymerization with varying polymerization time at fixed monomer pressure (1.1 atm) are summarized in Table 1. For comparison, two entangled Table 1. Molecular Characteristics of the Polymers Synthesized at 1.1 atm Monomer Pressure run 1 2 3 4b 5b 6 7c 8c

polym time (min)

Mwa (×106 g/mol)

2 10 30 60 60 90

0.7 2.3 5.3 8.0 9.9 14.0 7.0 5.0

PDIa 1.4 2.3 2.6 3.1 2.6 4.0 14.3

τ1 (min)

τ2 (min)

τ2/τ1

6.6 2.1 1.0 0.8 0.8 0.7 2.1 2.2

7.2 6.9 7.8 8.9 10.3 10.6 4.0 3.9

1.1 3.3 7.6 10.8 13.0 16.1 1.9 1.8

a

As determined by rheology. bMw and PDI data taken from previous reported findings.32 cCommercial entangled UHMWPE samples.

samples, runs 7 and 8, synthesized using a Ziegler−Natta catalyst, obtained from commercial sources are used. All samples, including the commercial samples, in the first heating run showed a melting peak close to 138 °C and crystallinity between 75 and 78%. For the estimation of crystallinity a reference of 293 J/g for 100% crystallinity was used.31 From Table 1 it is evident that Mw increases with polymerization time at a fixed ethylene pressure. These observations are in accordance with our previous findings where the same catalytic system is used.17,32 Crystallinity of the

Figure 1. Graphical representation on the evolution of entanglements formation during synthesis. The figure is partially reproduced from ref 16. The figure depicts the chains growing away from the active site, and the “heads” of the chains become entangled at the very initial stages of polymerization. Nucleation preceded by crystallization, favored by the supercooled state, freezes the initial entangled state. With the suppression of the nucleation barrier, the ethylene units growing at the end of the chain crystallize and retrospectively reduce the entanglement formation. From this model it is anticipated that number of entanglements per chain length decreases with the increasing molar mass. Such a possibility is represented by a decrease in the initial storage modulus with increasing molar mass or polymerization time for the catalyst used (see figure on the right). D


Article

Macromolecules

Figure 2. Characteristic melting times (τ1, a; τ2, b) as a function of Mw for the polymers of runs 1−6 (black filled squares) and the commercial entangled samples (red unfilled squares) for comparison.

Figure 3. Characteristic melting time of the nascent powder normalized by the characteristic melting time of the melt-crystallized as a function of Mw for the polymers of runs 1−6 (black filled squares) and the commercial entangled samples (red unfilled square) for comparison. The arrow in the figure refers to the value of entangled Mw at τ2/τ1 = 1.

crystallinity of 75% measured by NMR13) during conventional DSC scan. The difference in the rate of melting becomes apparent on changing the heating rate or performing the annealing experiments close to the onset of the melting temperature.36 On second heating, independent of the polymerization conditions, in the melt-crystallized samples both melting temperature and crystallinity reduce to 135 °C and below 50%, respectively. Compared to the respective nascent samples, the melt-crystallized samples show considerable increase in the transfer of polarization from noncrystalline to crystalline region of the semicrystalline polymer. The increase in the polarization time is attributed to the entropic barrier that methylene segments have to overcome for chain diffusion between the mobile noncrystalline and the rigid crystalline domains of the melt-crystallized samples.34 Consid-

segments synthesized in the initial stages of polymerization. Such a possibility has been schematically depicted in Figure 1. The comparative commercial samples, runs 7 and 8, in spite of high molar mass, but a broad molar mass distribution shows relatively less departure between the characteristic times of the nascent and the melt-crystallized sample. The difference in the segmental mobility in the noncrystalline region of the commercial samples, in its nascent and melt-crystallized state, has been also addressed by solid-state NMR.35 To recall, Yao et al. showed faster transfer of polarization from noncrystalline to crystalline region in the disentangled crystals, synthesized using the same catalytic system, compared to the entangled crystals synthesized using a Ziegler−Natta catalyst. However, independent of the entangled states, the nascent crystals of both polymers show high melting temperature of 141 °C (same E


Article

Macromolecules

nascent sample. The conclusion is in agreement with the higher entangled state formed during polymerization using a Ziegler− Natta catalyst under commercial conditions. From Figure 3, the value of molecular weight corresponding at τ2/τ1 = 1 is indicative of nascent crystals having maximum number of entanglements. The value of entangled Mw, at τ2/τ1 = 1, is found to be 0.50 × 106 g/mol. This molar mass is representative of the minimum chain length that during the adopted synthesis conditions results into a maximum entangled statepossibly a minimum required chain length to overcome the nucleation barrier for further crystallization in disentangled state. Thus, making use of the entangled Mw, it is possible to determine the entangled fraction for the polymers, shown in Table 2.

ering the similar polymerization conditions of the samples (runs 1−6), the influence of the entangled state in the melting response of the samples, having similar crystallinity of 75%, becomes apparent from Figures 2−5. Figure 2a shows an exponential decay in the characteristic melting time of the nascent powder as a function of the molar mass. Considering similar crystal thickness of these samples, evident from electron microcopy and SAXS,13 we attribute the decay in the characteristic melting time of the nascent polymer to reduction in the entanglement density in the noncrystalline region. The comparative commercial samples because of different polymerization conditions combined with its molecular characteristics and entanglement density fall out of the decay curve. These observations on decay in characteristic melting time with increasing reduced number of entanglements with molar mass are in agreement with the melting kinetics.18 To recall, the time needed in melting of crystals, under isothermal condition, is reported to be strongly dependent on the detachment of methylene segments from crystal surface and their reeling-in into melt. The detachment time is found to be dependent on the constraints that the chain segment, in the crystalline zone, experiences from the topological constraints present in the noncrystalline region. Considering that the characteristic melting time characterizes the kinetics of the melting transition of the crystallites, recorded by modulation over linear ramp, it is independent of both lateral and thickness of crystals.24 Thus, in this respect modulated DSC provides clear perspective of linking melting kinetics with topological constraints. The influence of topological constraints on melting kinetics, or characteristic time, becomes more evident when the same nascent sample is crystallized from its melt, ramps C−D and D−E in Scheme 1. Figure 2b shows changes in the characteristic time of the melt crystallized sample (τ2) as a function of molar mass, recorded during ramp E−F in Scheme 1. Unlike the decay observed in Figure 2a, in Figure 2b the characteristic time, τ2, increases with increasing molar mass, suggesting increasing number of topological constraints with increasing chain length. In Figure 3, the ratio between the two characteristic times, τ2/τ1, is depicted as a function of the molecular weight. A cause for the remarkable linear relationship with increasing polymerization time or molar mass for the same polymerization conditions can be attributed to the gradual decrease in the entanglement density with polymerization time and subsequent increase in the entangled state of the same sample crystallized from its melt state. From Table 1, for the same polymer, the values of τ2 are always found to be higher than the values of τ1. The τ2 reflects the characteristic melting time of the melt-crystallized sample having higher number of entanglements compared to the nascent state of the same sample; therefore, the values of τ2/τ1 will be higher than or equal to 1. Runs 1−5 follow a linear relationship between the characteristic melting time ratio τ2/τ1 and the Mw. The ratio of the run 6 is found to be out of the linear relationship, probably because at longer polymerization time chain termination processes followed by formation of new chains are likely to occur. The very low value of τ2/τ1 in the commercial samples is representative of smaller differences in the topological constraints of nascent and melt-crystallized samples, though crystallinity of the melt sample reduces below 50% compared to the 70% (approximately) crystallinity of the

Table 2. Entangled Fraction of the Polymers run

Mw (×106 g/mol)

entangled fraction

1 2 3 4 5

0.7 2.3 5.3 8.0 9.9

0.73 0.22 0.10 0.06 0.05

Figure 4 shows the evolution of the entanglement fraction with increasing molar mass. The adopted approach holds some limitations; for example, polymer of run 6 cannot be placed in the figure because it does not fit in the linear regression of Figure 3, and thus the determined entangled Mw is not representative for this sample. Considering that both the melting kinetics and processing are influenced by entangled state in the noncrystalline region of the semicrystalline polymer raises the question about a possible relationship between polymer processing and kinetics in melting. However, it is to be realized that melting kinetics is more of a local process depending on connectivity between the adjacent chains (intracrystalline topology) compared to the ease in processing, which is dependent on not only the intracrystalline topology but also connectivity between crystals (intercrystalline topology). At low draw ratios, the intercrystalline topology is desired for crystal orientation, whereas subsequent mechanically induced reeling of chains from crystalline domains are feasible when chains in the intracrystalline region are less entangled. 3.2. Polymer Processing and Mechanical Properties. Implementing the processing data of our previous study17 with some of the polymers synthesized in Table 1, it is possible to determine a correlation between the molecular weight and the total drawing tension. The drawing tension is defined as the force experienced by the rollers to stretch the sample normalized by the cross-sectional area of the tape at the given draw ratio. All the polymers synthesized can be uniaxially oriented in the solid state to high draw ratios (Figure 5). According to previous study this is related with a lower amount of entanglements in the noncrystalline region, established between the crystals. For all the polymers tested, the total drawing tension increases almost linearly as a function of draw ratio. The increase in the drawing tension, as a function of draw ratio, is attributed to the continuous chain alignment that increases the interchain van der Waals interaction. The force required to uniaxially draw the material is lower for lower molecular weights (or lower polymerization time). It is possible to determine the ease in F


Article

Macromolecules

Figure 4. Entangled fraction as a function of Mw for the polymers of runs 1−5 of Table 2.

Figure 5. Total drawing tension as a function of draw ratio for the polymers synthesized in Table 1. Since run 1 cannot be measured due to low yield in the given polymerization time, the data cannot be included.

processing of the compressed polymer into tapes. In particular, by fitting the values of total drawing tension as a function of draw ratio, we can define a processability index (PI) as the resulting slope of the linear fitting. The cause for the increase in PI with the molar mass is attributed to the increase in the number of entanglements in spite of the decrease in entanglement density. Moreover, the increase in molar mass also increases the force required for slippage between the crystalline planes. Table 3 summarizes the processability index and the mechanical properties for the polymers of Table 1. It is apparent that all samples, independent of their high molar mass, are processable in solid state and provide good mechanical properties. From Figure 6 it is apparent that the processability index, a quantitative measure for resistance in processing, increases with increasing molar mass. To have more information on the trend, the processing indexes of the polymers synthesized at different

Table 3. Processability Index (PI) and the Main Mechanical Properties at a Draw Ratio of ∼200 PI run (N/mm2) 2 3 5 6

1.8 3.2 4.5 4.8

Mw (×106 g/mol)

breaking tenacity (N/tex)

tensile modulus (N/tex)

energy at break (J/g)

2.3 5.3 9.9 14.0

3.71 4.05 4.19 3.93

185 191 194 192

46 57 59 51

ethylene pressures 2.1 and 4.1 atm are added in the figure. More details on the synthesis conditions of these polymers can be found in Table 4. At first instance the increase in processability index with subsequent decrease in the entanglement fraction (Figure 4) may seem to be in contradiction. It is important to understand the methodology adopted in estimation of the entanglement fraction and the origin of the processability index. The entanglement fraction has been G


Article

Macromolecules

knowing the characteristic time ratio and molar mass of the synthesized polymer a good indication on the ease in processing can be established. This route opens the possibility of estimating mechanical properties by thermal analysis and linking them to synthesis conditions. Such a possibility has been feasible because of the common determining denominator, entangled state, in the thermal and mechanical response of the semicrystalline polymer.

4. CONCLUSIONS In this study we have reported the synthesis and characterization of some disentangled UHMWPE samples having different molecular weight and/or entanglement density synthesized at various polymerization time and/or monomer pressure. The disentangled state is characterized using temperature-modulated DSC, dependence on the characteristic melting time, and ease in processing (processability index) has been established. The processability index refers to force experienced by the rollers during solid-state deformation of the nascent compressed powder into tapes. With the increasing draw ratio tensile strength and tensile modulus are found to increase. The characteristic melting time of the nascent polymers synthesized at 1.1 atm decays with molar mass. The decay is related with the decrease in entanglement density for longer polymerization time (and molecular weight) that lowers the time required for the consecutive detachment of chain stems during slow melting. The characteristic melting time of the melt-crystallized sample when normalized by the characteristic melting time of its nascent state (τ2/τ1) shows linear increase with the increasing molecular weight. The minimum entangled molar mass, created during synthesis, can be determined by extrapolating the linear fitting to τ2/τ1 = 1 for the 1.1 atm series. This value is used to trace the entangled fraction present in the synthesized polymer series. Important to notice is that for the single-site catalyst investigated in this study, under same polymerization conditions, entangled fraction decreases with the increasing polymerization time. The processability index of the samples synthesized show dependence on the molecular weight. In particular, PI increases with increasing Mw. It is concluded that independent of the polymerization conditions, a qualitative indication of the outcome of the mechanical properties can be given by the characteristic melting time of the polymer. In particular, the value of τ 2 /τ 1 independent of the polymerization conditions for all catalytic systems can be used to predict the mechanical properties of the polymer under investigation. Hence TM-DSC, combined with

Figure 6. Processability index (PI) as a function of Mw for the polymers of Tables 3 and 4. The dotted lines are fitted on the samples synthesized using the monomer pressure. Independent of the monomer pressure, the black line has been fitted over all the samples.

determined by considering the ratio of the characteristic melting time, which is dependent on the intracrystalline topology. On the other hand, force required for uniaxial deformation is influenced by chain topology in the intercrystalline as well as intracrystalline regions. With increasing molar mass the density of the topological constraints in the intracrystalline region decreases, while the constraints in the intercrystalline region increase. Thus, in spite of the decrease in the entanglement fraction with molar mass, considering the requirement of coherent deformation of intercrystalline regions during uniaxial deformation, the required force increases with increasing molar mass. Moreover, the increase in secondary forces between the crystalline planes, with the increasing molar mass, will also contribute to the force required for uniaxial deformation. Figure 7 shows dependence on the characteristic time ratio over molar mass for the samples in Tables 3 and 4. The polymers that fall on the linear line show higher breaking energy (>56 J/g) and were easier to process. The breaking energy is considered to be the measure for processing of the set of the samples as it is the determining factor for many composite applications, which depends on the tensile strength, tensile modulus, and elongation to break. Whereas the polymers, including the commercial samples (runs 7 and 8), which fall out of the linear line were difficult to process under the processing conditions established for the processing of disentangled UHMWPE. In this respect, from Figure 7, only by

Table 4. Molecular Characteristics for the Polymers Synthesized at Different Reaction Conditions

a

run

polym time (min)

monomer pressure (atm)

Mwa (×106 g/mol)

PDIa

PI (N/mm2)

τ1 (min)

τ2 (min)

τ2/τ1

breaking tenacity (N/tex)

tensile modulus (N/tex)

breaking energy (J/g)

9 10 11b 12 13 14b 15c

10 30 60 10 30 60 60

2.1 2.1 2.1 4.1 4.1 4.1 2.0

4.3 10.8 15.3 9.6 14.8 34.0 7.2

3.7 4.5 3.1 3.7 5.8 7.2 5.5

3.3 5.2 5.8 5.3 6.0 3.3 1.5

1.1 0.7 0.6 0.7 0.6 0.7 1.0

8.4 9.3 7.2 8.1 7.7 7.1 6.2

7.8 13.9 11.5 11.2 12.1 9.7 6.1

4.32 4.18 4.10 4.17 3.62 3.74 3.92

190 194 210 195 201 207 204

61 57 56 59 45 46 48

As obtained by rheology. bData taken from previous reported findings.17 cUsing the catalytic system described in ref 14. H


Article

Macromolecules

Figure 7. Ratio τ2/τ1 as a function of Mw. The dashed line is chosen as reference for the mechanical properties. Polymers below the reference line go in the direction of the lower mechanical properties in contrast with the polymers above the reference line. The values in the figure refer to breaking energy in J/g for individual samples synthesized at different polymerization conditions. The breaking energy is determined for the tapes having the draw ratio of ∼200. The sample synthesized in run 11, in spite of it not falling on the test line, shows higher breaking energy (56 J/g) as it was processed at the higher temperatures than the remaining samples. (4) Smith, P.; Lemstra, P. J.; Booij, H. C. Ultradrawing of highmolecular-weight polyethylene cast from solution. II. Influence of initial polymer concentration. J. Polym. Sci., Polym. Phys. Ed. 1981, 19, 877−888. (5) Smith, P.; Lemstra, P. J. Ultra-high-strength polyethylene filaments by solution spinning/drawing. J. Mater. Sci. 1980, 15, 505−514. (6) Smith, P.; Lemstra, P. J. Ultrahigh-strength polyethylene filaments by solution spinning/drawing, 2. Influence of solvent on the drawability. Makromol. Chem. 1979, 180, 2983−2986. (7) Smith, P.; Lemstra, P. J. Ultra-high strength polyethylene filaments by solution spinning/drawing. 3. Influence of drawing temperature. Polymer 1980, 21, 1341−1343. (8) Smith, P.; Lemstra, P. J. Tensile strength of highly oriented polyethylene. J. Polym. Sci., Polym. Phys. Ed. 1981, 19, 1007−1009. (9) Pennings, A. J.; Van der Hooft, R. J.; Postema, A. R.; Hoogsteen, W.; Ten Brinke, G. High Speed gel spinning of Ultra High Molecular Weight Polyethyelen. Polym. Bull. 1986, 16, 167−174. (10) Schaller, R.; Feldman, K.; Smith, P.; Tervoort, T. A. Highperformance polyethylene fibers “al dente”: Improved gel-spinning of ultrahigh molecular weight polyethylene using vegetable oils. Macromolecules 2015, 48, 8877−8884. (11) Kanamoto, T.; Ohama, T.; Tanaka, K.; Takeda, M.; Porter, R. S. Two-stage drawing of ultra-high molecular weight polyethylene reactor powder. Polymer 1987, 28, 1517−1520. (12) Smith, P.; Chanzy, H. D.; Rotzinger, B. P. Drawing of virgin ultrahigh molecular weight polyethylene: An alternative route to high strength/high modulus materials. Part 2 Influence of polymerization temperature. J. Mater. Sci. 1987, 22, 523−531. (13) Rastogi, S.; Yao, Y.; Ronca, S.; Bos, J.; Van der Eem, J. Unprecedented high-modulus high-strength tapes and films of ultrahigh molecular weight oolyethylene via solvent-free route. Macromolecules 2011, 44, 5558−5568. (14) Romano, D.; Ronca, S.; Rastogi, S. A hemi-metallocene chromium catalyst with trimethylaluminum-free methylaluminoxane for the synthesis of disentangled ultra-high molecular weight polyethylene. Macromol. Rapid Commun. 2015, 36, 327−331. (15) Rastogi, S.; Lippits, D. R.; Peters, G. W. M.; Graf, R.; Yao, Y.; Spiess, H. W. Heterogeneity in polymer melts from melting of polymer crystals. Nat. Mater. 2005, 4, 635−641.

the data given in Figure 7, can be used as a fast tool to predict the resultant mechanical properties and ease in processing of the uniaxially deformed UHMWPE in its solid state. This approach has been feasible because of the common factor, entangled state, involved in the thermal and mechanical response of the semicrystalline polymer. In this paper we have made an attempt to correlate the melting kinetics with the processability of nascent UHMWPE reactor powder in the solid state, viz. ultradrawing. TM-DSC is a fast tool to measure the ratio (τ2/τ1) which provides a qualitative indication regarding processability. The authors realize that TM-DSC measurements focus on intracrystalline topology and arrangement of stems whereas during the total draw process other factors play a role such as intercrystalline entanglements and intermolecular van der Waals forces during the final orientation of chains, and consequently, TM-DSC measurements can predict the ease and feasibility of solid state processing influencing the mechanical properties.

■

AUTHOR INFORMATION

Corresponding Author

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

Sanjay Rastogi: 0000-0002-7804-7349 Notes

The authors declare no competing financial interest.

■

REFERENCES

(1) Abiru, S.; Iizuka, K. Bulletproof fabric and process for its production. U.S. Patent 5,935,881, 1999. (2) Bovenschen, S. J.; Van der Eem, J.; Rastogi, S.; Harings, J. A. W.; Schaap, A. A. Process for making a curved product. EP 2650111A1, 2011. (3) Kurtz, S. M. The UHMWPE Handbook; Ultra High Molecular Weight Polyethylene in Total Joint Replacement; Elsevier Academic Press: New York, 2004. I


Article

Macromolecules (16) Pandey, A.; Champouret, Y.; Rastogi, S. Heterogeneity in the distribution of entanglement density during polymerization in disentangled ultrahigh molecular weight polyethylene. Macromolecules 2011, 44, 4952−4960. (17) Romano, D.; Tops, N.; Andablo-Reyes, E. A.; Ronca, S.; Rastogi, S. Influence of polymerization conditions on melting kinetics of low entangled UHMWPE and its implications on mechanical properties. Macromolecules 2014, 47, 4750−4760. (18) Pandey, A.; Toda, A.; Rastogi, S. Influence of amorphous component on melting of semicrystalline polymers. Macromolecules 2011, 44, 8042−8055. (19) Yao, Y.; Jiang, S.; Rastogi, S. 13C solid state NMR characterization of structure and orientation development in the narrow and broad molar mass disentangled UHWMPE. Macromolecules 2014, 47, 1371−1382. (20) Keller, A. A note on single crystals in polymers: Evidence for a folded chain configuration. Philos. Mag. 1957, 2, 1171−1175. (21) Flory, P. J.; Yoon, D. Y.; Dill, K. A. The interphase in lamellar semicrystalline polymers. Macromolecules 1984, 17, 862−868. (22) Toda, A.; Hikosaka, A.; Yamada, K. Superheating of the melting kinetics in polymer crystals: A possible nucleation mechanism. Polymer 2002, 43, 1667−1679. (23) Toda, A.; Arita, T.; Hikosaka, M. Transition kinetics of a Ti−Ni alloy examined by temperature-modulated DSC. Thermochim. Acta 2005, 431, 98−105. (24) Toda, A.; Kojima, I.; Hikosaka, M. Melting kinetics of polymer crystals with an entropic barrier. Macromolecules 2008, 41, 120−127. (25) Pandey, A. V. Nonlinear viscoelastic response of a thermodynamic metastable polymer melt. PhD Thesis, Loughborough University, 2011. (26) Andablo-Reyes, E. A.; De Boer, E. L.; Romano, D.; Rastogi, S. Stress relaxation in the nonequilibrium state of a polymer melt. J. Rheol. 2014, 58, 1981−1991. (27) Romano, D.; Andablo-Reyes, E. A.; Ronca, S.; Rastogi, S. Aluminoxane co-catalysts for the activation of a bis phenoxyimine titanium (IV) catalyst in the synthesis of disentangled ultra-high molecular weight polyethylene. Polymer 2015, 74, 76−85. (28) Rastogi, S.; Odell, J. A. Stress induced stabilization of hexagonal phase in gel spun fibre of polyethylene. Polymer 1993, 34, 1523−1526. (29) Van Aerle, N. A. J. M.; Lemstra, P. J. Chain-Extended Polyethylene in Composites Melting and Relaxation Behaviour. Polym. J. 1988, 20, 131−141. (30) Van Aerle, N. A. J. M.; Lemstra, P. J.; Braam, A. W. M. A realtime X-ray melting study of partly drawn UHMW-polyethylene. Polym. Commun. 1989, 30, 7−10. (31) Wunderlich, B. Crystal Melting, Macromolecular Physics; Academic Press: New York, 1980; Vol. 3. (32) Romano, D.; Andablo-Reyes, E. A.; Ronca, S.; Rastogi, S. Effect of a cocatalyst modifier in the synthesis of ultrahigh molecular weight polyethylene having reduced number of entanglements. J. Polym. Sci., Part A: Polym. Chem. 2013, 51, 1630−1635. (33) Ronca, S.; Romano, D.; Forte, G.; Andablo-Reyes, E. A.; Rastogi, S. Improving the performance of a catalytic system for the synthesis of ultra high molecular weight polyethylene with a reduced number of entanglements. Adv. Polym. Technol. 2012, 31, 193−204. (34) Yao, Y.; Graf, R.; Spiess, H. W.; Rastogi, S. Influence of crystal thickness and topological constraints on chain diffusion in linear polyethylene. Macromol. Rapid Commun. 2009, 30, 1123−1127. (35) Yao, Y.; Rastogi, S.; Xue, H. J.; Chen, Q.; Graf, R.; Verhoef, R. Segmental mobility in the non-crystalline regions of nascent polyethylene synthesized using two different catalytic systems with implications on solid state deformation. Polymer 2013, 54, 411−422. (36) Lippits, D. R.; Höhne, G. W. H.; Rastogi, S. Melting Kinetics in Polymers. Phys. Rev. Lett. 2006, 96, 218303−218307.

J


Correlation between Thermal and Mechanical ... - ACS Publications

Recommend Documents