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Enhancing toughness of low-density polyethylene filaments through infusion of multiwalled carbon nanotubes and ultrahigh molecular weight polyethylene.

INTRODUCTION

Traditional applications of polymeric fibers have been in textiles and furnishings. However, in recent years, fiber applications have expanded markedly to industries such as geo-textile, composites, aerospace, automotive, biomedical products, protective clothing, plastic bags, containers, and packaging where absorption of energy under low or high strain loading would be important. Any increase in elastic energy [1] of the fiber will no doubt be beneficial for these industries. Low-density polyethylene (LDPE) with its excellent chemical resistance, flexibility, and electrical insulation is a prime candidate for such applications. To satisfy the broad spectrum of performance demanded by these industries, LDPE must have adequate tensile strength, high modulus, high fracture strain, and long-term thermal and structural stability. Especially, if strength, modulus, and fracture strain of LDPE can be enhanced, it will significantly increase its elastic energy. LDPE have been used for many years to extract useful fibers, and it is attempted in this article to enhance its mechanical and thermal properties further through nanoscale inclusion and blending of polymers.

LDPE has branched and linear chain structure with a molecular weight of typically <50,000 g/mol [2, 3]. It is considered as an important thermoplastic, especially for its low density, good processability, and easier moldability. There have been considerable efforts in recent years to increase the strength and modulus of polymeric fibers, including LDPE, through nanoscale reinforcement. Over the past several years, researchers have attempted to infuse carbon nanotubes (CNTs) into LDPE and into other thermoplastics [4-11]. These studies have demonstrated that infusion of CNTs or carbon nanofibers into textile polymers through extrusion or spinning have consistently improved both mechanical and thermal properties. With the addition of 2 wt.% of carbon nanoparticles in LLDPE, the tensile strength and modulus of the composite have been shown to increase by about 17% [12]. Composites using multiwalled carbon nanotubes (MWCNTs) were found to possess higher tensile strength and modulus by about 34-38% compared with neat polyethylene control samples [13]. Tensile properties of nylon 6 infused with MWCNTs have also been reported. It has been shown that with 1.0 wt% loading, the strength and modulus could be enhanced almost by a factor of two [14]. To use the extraordinary strength and stiffness of carbon nanotubes in bulk materials, several other researchers [9-11, 15] have dispersed CNTs into nylon and attempted to align CNTs along the length of the drawn filament. Infusion was performed through either a liquid route using sonication or a dry route followed by melt mixing in an extruder. Alignment of CNTs in the filament was enforced by extrusion and spinning. The resulting composites either in consolidated or in filament form have demonstrated significant improvement in mechanical and thermal properties. Although it was encouraging to see phenomenal increase in strength and stiffness [14], it was also observed that the improvement was at the cost of sacrificing a considerable amount of fracture strain. In an attempt to improve on the fracture strain, spherical silica nanoparticles in pristine and functionalized forms were used [16]. The result was somewhat mixed. Introduction of spherical silica particles allowed modest improvement in strength and modulus of nylon by about 30%, but fracture strain remained high at 80% similar to that of unreinforced nylon. This allowed strain energy to increase by approximately 11%. When particles were functionalized, fracture strain of nylon however dropped to 30%, but strength and modulus increased by 76 and 55%, respectively. This resulted in 58% reduction in strain energy in case of functionalized particles. All comparisons were made with respect to neat (unreinforced) nylon.

In recent years, ultrahigh molecular weight polyethylene (UHMWPE) has also been reinforced with CNTs to enhance both mechanical and thermal properties [17-22], Fibers have been mostly produced through a gel-spinning process. It was reported that [23] at intermediate strains, CNTs may act as slippage sites, allowing the matrix to deform without significant bond stretching. However, these sites may also act as pseudo taut tie molecules at large strains to produce strain hardening effects that are absent from the highly anisotropic pure UHMWPE filaments. Because of such tying effects, substantial enhancement in strain hardening behavior of UHMWPE filaments [24] has been reported. However, increase in toughness and associated elastic energy remained in the modest range.

Brief overview presented above gives a good insight as to how the toughness of LDPE fiber can be enhanced further. One effective way to deal with this issue is to add a second (minor)-phase polymer with the CNT-reinforced LDPE matrix. This can induce a sliding behavior between polymer-polymer interfaces during loading, resulting in a large tensile elongation. The phenomenon is much similar to what is seen when a polar and apolar polymer are blended in the presence of a compatibilizer [25]. The minor phase would get dispersed as droplets and elongate under the stress, allowing large deformation of the fiber. We have considered UHMWPE as a minor phase in this investigation. UHMWPE is a linear homopolymer, which can induce complex polymer architecture with a large amount of interfaces. The molecular chain of UHMWPE consists of as many as 200,000 ethylene repeat units with an average molecular weight of 3 X [10.sup.6] g/mol [26, 27]. Its longer chain configuration compared with LDPE helps form a large phase-separated domain when integrated into the LDPE matrix. At molecular level, the carbon backbone of UHMWPE can twist, rotate, and fold into ordered crystalline regions and change crystallite size of the combined polymer. The large molecular structure of UHMWPE would also interact with embedded nanoparticles and facilitate nucleation during the cross-linking phase. The idea in this investigation is to use this concept of dual inclusion that is a second-phase polymer and nanotubes to collectively harness their individual benefits. Second-phase polymer would allow improvement of fracture strain, whereas nanoparticles would enhance strength and modulus.

MATERIALS SYNTHESIS AND FABRICATION

Commercial grade LDPE and UHMWPE were procured from Sigma Aldrich Co (Milwaukee, WI) The average density and melting point of LDPE were 0.925 g/[cm.sup.3] and 116[degrees]C, whereas that of UHMWPE were 0.94 g/[cm.sup.3] and 138[degrees]C, respectively. The MWCNTs were obtained form Nanocraft (NanoCraft, Renton, WA). The nanotubes were 10-20 nm in diameter and approximately 2 [micro]m long. Materials synthesis and fabrication of filaments were performed for three categories of samples: (1) Neat LDPE; (2) LDPE with MWCNT; and (3) LDPE with MWCNT and UHMWPE. A mechanical crusher was used to produce micron size powders from procured LDPE. Then LDPE powders were poured through a feed hopper in a single-screw laboratory mixing extruder (LME). After the feed hopper, the materials passed through an annular zone of the extruder where actual melting took place. The annular zone was formed by a turning rotor placed inside the extruder. The temperature in the annular zone was controlled by heat conduction to the rotor. As rotation continued, the molten mass experienced shear force enabling it to flow to the outlet dies and exit through the orifice. Orifice diameter was 3 mm. To draw the material from LME into filaments, LME take-up system was used. LME take-up system is a filament winder attached to LME to collect spun filaments. It is a variable drive machine which can draw and wind small diameter extrudate fibers from the LME onto a spindle. The speed of the system was adjusted to match the rate of extrusion and provide the desired fiber diameter. As the filaments were stretched, diameter reduced substantially. Filaments were then wrapped around a spindle and drawn at 76 m/min. Draw ratio was more or less same for all categories of samples. The rotor and header temperature was set at 125 and 150[degrees]C, respectively. For second category of filaments, MWCNTs were dry mixed with LDPE. The dry mixing was performed in a mechanical blender for about 2 hr. Filament extrusion process was identical. In the third category, both MWCNTs and UHMWPE were dry mixed with LDPE, and the manufacturing process followed as before except readjusting the rotor and header temperatures. The temperatures were set at 200 and 220[degrees]C, respectively. Higher extrusion temperature was necessary to increase melt flow when UHMWPE was added.

X-RAY DIFFRACTION

Powder x-ray diffraction (XRD) measurements were performed using a Siemens D5000 powder diffractometer operating at 45 kV and 40 mA with CuK[alpha] radiation and a diffracted beam monochromator. Data were collected in the 2[theta] range of 8-90[degrees] with a step size of 0.02[degrees] and a counting time of 2 s at each step. The data bank from the International Center for Diffraction Data was used in a search/match program for phase identification. XRD traces of (a) neat LDPE; (b) LDPE-MWCNTs; and (c) LDPE-MWCNTs-UHMWPE are shown in Fig. 1. The polyethylene phase [([CH.sub.2]).sub.x] (PDF # 40-1995) was identified in the three types of samples, as it was expected because of the small weight percentage of the additive MWCNTs and UHMWPE. Small additional peaks in samples (b) indicate presence of a secondary phase, but below the experimental detection limits (~1% wt). The Miller indices 110 and 200 of the main Bragg peaks and 200 reflections of polyethylene are marked in the figure. The diffraction traces show that all three samples consist of a mixture of well-crystallized phase with amorphous regions. The crystallite size in the samples was estimated from the full width at half maximum (FWHM) of the 110 diffraction peaks by using the Scherer equation [28],

[tau] = K[lambda]/[beta]cos([theta]),

where [tau] is the crystallite size, [beta] is the FWHM for the peak in radians, [lambda] is the wavelength of the incident x-ray radiation, and [theta] is the diffraction. K is an experimental constant equal to 0.9 [29]. The calculated values of the crystallite size t are listed in Table 1. Value of [tau] seems to be slightly higher with the MWCNT infusion, whereas a small reduction in crystallite size appears when UHMWPE is added but still slightly higher compared with that of the neat LDPE. It is reasonable to expect an increase in crystallite size with the inclusion of MWCNTs because they can serve as nucleation agents for growth of crystallites during polymerization.

[FIGURE 1 OMITTED]
TABLE 1. XRD results.

Type of sample     FWHM (degrees)  Crystallite size [tau] (nm)

Neat LDPE               0.50                   16
LDPE-MWCNT              0.46                   18
LDPE-UHMWPE-MWCNT       0.47                   17


DIFFERENTIAL SCANNING CALORIMETRY TESTS

The differential scanning calorimetry (DSC) ramp tests were performed under inert atmosphere with a heating rate of 10[degrees]C/min. Tests were performed in a TA Q10 DSC apparatus with about 5 mg samples. DSC curves provided two key pieces of information: (1) melting temperature and (2) degree of crystallinity. Melting temperature was distinct as shown by DSC curves in Fig. 2. Degree of crystallinity was determined from the area under the melting endotherm and heat of fusion for polyethylene [30, 31]. Because the amount of heat required is proportional to the fraction of crystal present, the melting endotherm can be used to measure degree of crystallinity. The melting point and percent crystallinity are shown in Table 2. A moderate increase in crystallinity and a reduction in melting temperature by about 5[degrees]C are observed when MWCNTs were added. With the inclusion of UHMWPE, melting point and crystallinity did not change. Embedded MWCNTs inside the polymer matrix can act as nucleation agents for crystal growth around the nanotubes. This may lead to coating of nanotubes with a layer of crystalline polymer and thus an interfacial region can develop [32]. Bulk morphological changes can also take place because of MWCNT-nucleated crystallization. Reduction in melting temperature by nanoscaled materials has been reported earlier [33, 34]. The melting point of a solid or crystal is reached when the order of the lattice is beginning to be destroyed [35]. This transition from solid to liquid starts at the surface. Because nanotubes have large number of atoms at the surface, they can be rearranged or dislocated more easily during heating, causing the melting process to start early. Crystalline layer at the nanotube surface will therefore have a lower melting point than the surrounding polymer. As a result, the melting point of the nanotube-reinforced polymer system will be lower than their neat counterpart as seen in Table 2.
TABLE 2. DSC results.

Type of sample     Melting point, [T.sub.m]  Crystallinity (%)
                         ([degrees]C)

Neat LDPE                   116.7                  29.0
LDPE-MWCNT                  111.6                  39.5
LDPE-MWCNT-UHMWPE           111.4                  39.1


[FIGURE 2 OMITTED]

TENSILE TESTS

Individual filament from each category was tested under tension according to ASTM D3379-75 using a Zwick/Roell tension testing machine with a 20-N load cell. Tests were performed at a crosshead speed of 2.5 mm/rain with a gauge length of 30 mm. The diameter of the filament was measured using a scanning electron microscope (SEM). Tensile test data are presented in Table 3. Representative stress--strain diagrams for three categories of filaments are shown in Fig. 3. It is observed that the gain in tensile strength and Young's modulus are around 23 and 57%, respectively, with the inclusion of 2 wt% MWCNTs. It is also noticed that, the fracture strain of the filament, in this case, is reduced by 45%. When 8 wt% UHMWPE is added, modulus and fracture strain increased by 39 and 44%, respectively. There is however a slight reduction in ultimate strength from 26 to 25 MPa. It is obvious from Fig. 3 that inclusion of MWCNTs increases modulus and strength, whereas it reduces the fracture strain. On the other hand, when UHMWPE is added, fracture strain increases substantially from 99 to 143%. Although the strength and modulus in the LDPE-MWCNT-UHMWPE system are lower than that of the two-phase (LDPE-MWCNT) system, they are still higher in case of modulus, and almost equal in case of strength, compared with neat LDPE samples.
TABLE 3. Data from tensile tests.

Type of sample         Tensile strength,     Tensile modulus, [E.sub.y]
                   [[sigma].sub.ymax] (MPa)            (MPa)

Neat LDPE                26 [+ or -] 1            274 [+ or -] 10
LDPE-MWCNT               32 [+ or -] 4            430 [+ or -] 50
LDPE-MWCNT-UHMWPE        25 [+ or -] 2            381 [+ or -] 21

Type of sample     Ultimate fracture strain,  Modulus of toughness,
                   [[epsilon].sub.ymax] (%)         [u.sub.t]
                                                  (MJ/[m.sup.3])

Neat LDPE                 99 [+ or -] 8            22 [+ or -] 2
LDPE-MWCNT                54 [+ or -] 9            14 [+ or -] 4
LDPE-MWCNT-UHMWPE        143 [+ or -] 6            34 [+ or -] 4

Type of sample     Normalizing velocity [[cube root of ([OMEGA])]
                                         (m/s)

Neat LDPE                                 197
LDPE-MWCNT                                184
LDPE-MWCNT-UHMWPE                         230


[FIGURE 3 OMITTED]

It is to be mentioned here that, when we increased the weight percent of UHMWPE beyond 8%, there was significant reduction in strength but with an increase in fracture strain [18]. On the other hand, as we reduced the concentration of UHMWPE to 3-5 wt%, there was a modest increase in fracture strain but with a slight reduction in strength. For example, at 10 wt%, strength, modulus, and fracture strain values were 20 MPa, 332 MPa, and 184%, respectively. On the other hand, when concentration was reduced to 5 wt%, these values were 25 MPa, 398 MPa, and 109%, respectively. We have tried several such concentrations of UHMWPE and came up with 8 wt% concentration, which provided the best result.

As mentioned earlier, crystallinity increased from 29 to 39% when MWCNTs were added. A slight increase in crystallite size was also evident from the XRD tests. Because these changes in crystal structure were small, their effects on the tensile properties were modest. Improvement in strength and modulus was mainly due to interaction of nanotubes with the polymer chain. During the extrusion process, free radicals are generated by the C--C bond cleavage due to the extreme shear deformation of the polymer [36]. On the other hand, defects in the lattice of the nanotubes can cause charge localization on the lattice structure. It is therefore possible to form strong interaction between nanotubes and polyethylene chain using these free radicals. Significant reduction in fracture strain was also an indirect evidence of this phenomenon. Polymer chain extension was highly restricted because of the interaction of the backbone chain with highly stiff nanotubes.

The modulus of toughness ([u.sup.t]) was calculated from the area under the stress-strain curve as shown in Fig. 3 and Table 3. Trapezoidal rule was used with large number of data points to calculate the area under the curves. A reduction in modulus of toughness by 36% was observed in MWCNT-reinforced samples. Although the modulus and strength increased significantly, loss in fracture strain resulted in such reduction of [u.sub.t]. With the addition of UHMWPE, [u.sub.t], increased by about 55%.

Morphological changes in undeformed and deformed filaments were investigated using SEM. High-resolution images of filament cross-sections are shown in Fig. 4a and b. Evidence of microscopic minor phase in an extruded filament is shown in Fig. 4a. Figure 4a is a SEM picture of the cross-section of extruded filament before any deformation. Presence of droplets as minor phase is indicated by arrows. It is observed that droplets are of various sizes ranging from 1 to 5 [micro]m and distributed more or less uniformly over the entire cross-section of the filament. Each droplet is surrounded by LDPE matrix and forms an interface with the matrix as seen in Fig. 4a. Once the filament was deformed during tensile tests, presence of minor phase and interface was still evident as seen in Fig. 4b. Figure 4b shows fractured surface of a filament after conducting the tensile test. Minor phases seem to have elongated along the direction of the tensile load. Prolonged interfaces along the minor phase suggest that sliding between the minor phase and LDPE matrix might have occurred. During the extrusion process, as the molten materials came out through the orifice, the low viscous LDPE deformed more easily compared with the highly viscous UHMWPE. As both polymers solidified almost simultaneously, LDPE tended to encapsulate UHMWPE and interfacial regions were formed as seen in Fig. 4a and 4b. These interfacial regions can cause sliding behavior between the polymers as tensile load is applied, resulting in a large fracture strain.

[FIGURE 4 OMITTED]

Elastic Energy Storage Capacity ([cube root of([OMEGA])])

To assess fiber's capacity to withstand high strain rate loading, two properties are important: elastic energy storage capacity per unit mass and the tensile wave speed through the fiber. A parameter known as "normalizing" velocity is usually used to express these two quantities. Normalizing velocity is defined as [1]:

[cube root of ([ohm])] = [([[[[sigma].sub.ymax][[epsilon].sub.ymax]]/[2[[rho].sub.y]]][square root of ([[E.sub.y]/[[rho].sub.y]])).sup.1/3] (1)

where [[sigma].sub.ymax] is the ultimate tensile strength, [[epsilon].sub.ymax] is the fracture strain, [E.sub.y] is the modulus, and [p.sub.y] is the density of the fiber.([cube root of([OMEGA])]) is regarded as a performance predictor for the fiber to absorb impact energy. Normalizing velocities for various categories of filaments were calculated and are shown in Table 3. To determine ([cube root of([OMEGA])]), the mechanical properties such as [[sigma].sub.ymax], [[epsilon].sub.ymax], and [E.sub.y] were used from Table 3. As mentioned earlier, the densities of LDPE, UHMWPE, and MWCNTs were 0.925, 0.94, and 2.6 g/[cm.sup.3] respectively. The densities ([p.sub.y]) of LDPE-MWCNT and LDPE-MWCNT-UFIMWPE systems were determined using the rule of mixture and were 0.937 and 0.938 g/[cm.sup.3], respectively. Because the loading of nanotubes was only 21) wt%, the change in [p.sub.y] was small. As we have seen in case of [u.sub.t], normalizing velocity also reduced by about 7% because of inclusion of MWCNTs. However, with the addition of UHMWPE, normalizing velocity increased by 17%. It is therefore apparent that both fracture toughness and normalizing velocity are controlled by fracture strain and hence by the addition of UHMWPE.

STRAIN HARDENING EXPERIMENTS

For semi-crystalline polymers, rearrangements of amorphous portion of polymer chains can take place through strain hardening process. The ability of fibrils to undergo strain hardening can enhance the strength of the polymer. To visualize hysteresis effect on fiber, samples were strain hardened through a series of loading and unloading cycles. Strain hardening tests were simple tension tests in which samples were first loaded up to a point slightly beyond the yield point and then unloaded to zero load level. This completed the first loop. In the next step, samples were loaded again to a higher point, incremented by both load and deflection and followed by unloading. This formed the second loop. Such loading and unloading continued until the fiber failed. A set of hysteresis curves for each category of fibers were generated in this manner as shown in Fig. 5a. Next, true strains were calculated [37], and maximum true-stress vs. true-strain plots were drawn in logarithmic scale for each category of samples as shown in Fig. 5b. Each line in Fig. 5b depicts [sigma] = K[~.[epsilon].sup.n], where sigma] is the true stress, [~.[epsilon] is true strain, n is strain hardening exponent, and K is strength coefficient. Values of n and K are shown in Table 4. Value of n with LDPE-MWCNT samples was 0.74 compared with that of 0.58 with neat LDPE. Higher value of n in case of LDPE-MWCNT sample indicates that strengthening of filament is possible by strain hardening when MWCNTs are infused. However, that was not the case with LDPE-MWCNT-UHMWPE filaments. The value of n decreased to 0.49 even lower than that of the neat LDPE. As seen in Table 4, values of K fluctuated in the same manner as it was expected. Values of strength, modulus, and fracture strain obtained from the final loading cycle are also shown in Table 4. Normalizing velocity was then calculated using these values. Table 4 indicates a different trend in normalizing velocity than what we noticed with n and K. It was observed that normalizing velocity decreased with LDPE-MWCNT but increased with LDPE-MWCNT-UHMWPE filaments. The reason for such difference in behavior is because both fracture strain and modulus increased substantially (Table 4) for LDPE-MWCNT-UHMWPE samples. Now, comparing with Table 3, it is found that normalizing velocity basically did not change after strain hardening. Although tensile strength for each category of filaments increased after strain hardening, fracture strain reduced significantly causing normalizing velocity to remain unchanged.

[FIGURE 5 OMITTED]
TABLE 4. Data from strain hardening tests.

Type of sample     [[sigma].sub.ymax] (MPa)  [E.sub.y] (MPa)

Neat LDPE               70 [+ or -] 3        333 [+ or -] 13
LDPE-MWCNT              85 [+ or -] 3        624 [+ or -] 40
LDPE-MWCNT-UHMWPE       71 [+ or -] 4        423 [+ or -] 24

Type of sample     [[epsilon].sub.ymax] (%)  [u.sub.t] (MJ/[m.sup.3])

Neat LDPE              28 [+ or -] 4              10 [+ or -] 1
LDPE-MWCNT             16 [+ or -] 3               7 [+ or -] 1
LDPE-MWCNT-UHMWPE      37 [+ or -] 6              16 [+ or -] 2

Type of sample      n    K (MPa)  [[cube root of ([OMEGA])] (m/s)

Neat LDPE          0.58    74                    185
LDPE-MWCNT         0.74   137                    181
LDPE-MWCNT-UHMWPE  0.49    64                    212

[+ or -] for standard deviation.


In an attempt to examine the morphology of the strain-hardened samples, DSC tests were run on strain-hardened samples and curves are shown in Fig. 6. Percent crystallinity and melting points evaluated from the DSC curves are also shown in Fig. 6. It was observed in Fig. 6 that percent crystallinity of LDPE increased from 31 to 40% with LDPE-MWCNT and LDPE-MWCNT-UHMWPE samples, respectively. These values are similar to what were shown in Table 2, suggesting that strain hardening did not have much of an effect on crystallinity. Identical arguments could be made with melting temperatures before and after strain hardening. During strain hardening, the backbone chain of LDPE can be stretched and aligned along the loading direction. However, the irregular presences of short-chain branches in LDPE disrupt the sequence of ethylene mers ([CH.sub.2]); therefore, crystallinity of the filament did not increase even after the strain hardening process. On the other hand, a small second melting endotherm was observed in LDPE-MWCNT-UHMWPE sample as seen in Fig. 6. It was mentioned earlier that a polymer-polymer interface grew when UHMWPE were added in LDPE. It is believed that in this interface region, a semi-interpenetrating polymer network can be constituted by the entangled combination of two polymers largely different in their molecular weights [2], It is possible that molecular rearrangement in this polymer network can take place during the strain hardening cycles and can exhibit a small secondary crystal fraction as evidenced by the second melting endotherm (Fig. 6).

[FIGURE 6 OMITTED]

FRACTURE STUDIES

SEM micrographs of fractured filaments are shown in Fig. 7. The fractured samples were collected after tensile tests of individual filament in each category. SEM analysis was performed in a D8303 Quanta 200 SEM. It is reported in the literature [16, 38-40] that fracture process in textile fibers such as nylon 6 has three distinct features. These three features are formation of a V-notch, a region of slow crack growth, followed by a fast crack growth region characterized by a large and relatively rough surface. In Fig. 7a-c, we notice existence of only two such stages rather than three. Multiple V-notches are seen to have formed on the fracture surface of neat LDPE filament as shown in Fig. 7a. These V-notches have originated from surface crazes present on the outer surface of the filament. We also notice that the number of V-notches is proportional to the number of surface craze. These multiple V-notches basically define the first stage of fracture and seem to terminate in a region (Fig. 7a) where the surface seems to be rough and molten. This is the second and final fracture stage of neat LDPE filaments. The area for the final stage fracture in Fig. 7a is relatively small compared with the total cross-section of the filament. It is also featured by elongated and drawn polymer ligaments, suggesting that this area indeed formed during the final failure event. It is also noticed that all V-notches traversed the fiber cross-section in a transverse manner meaning that they were short lived as opposed to the prolonged ones that run along the length of the filament. When LDPE is reinforced with MWCNTs, the fracture process is still seen to be of two stages (Fig. 7b). However, the formation of V-notch is different--it is prolonged along the filament length, and there is only one dominant V-notch. The V-notch culminates to a rough region as seen in Fig. 7b. The area of this second-stage fracture in this case is much larger than what we have seen in Fig. 7a and seems to be almost equal to that of the cross-sectional area of the filament. A larger area would require a larger force to break and that is why we see highest strength with LDPE-MWCNT filaments. Now, looking at Fig. 7c, that is for LDPE-MWCNT-UHWMWPE filament, we see an identical fracture process as we have seen in Fig. 7b. The area of the second-stage fracture is smaller and that explains the lower strength of the third category of filament. But the question remains as to why the fracture strain of LDPE-MWCNT-UHMWPE filament (~143%) is 2.65 times higher than that of LDPE-MWCNT filaments (54%)--both having identical fracture pattern. The fracture strain during the tensile test can be visualized as the sum of two elongations--one is before the initiation of fracture and the other during the fracture process. What we have observed in Fig. 7 are features from initiation of fracture to final failure. An assessment of total elongation can however be made from the reduced diameter of the filament at the time of failure. Filament diameter in Fig. 7c is seen to be around 46 [micro]m, whereas that in 7b is around 83 [micro]m. Although the diameter of both categories of filament was around 120 [micro]m at the beginning of the test, LDPE-MWCNT-UHMWPE filaments have undergone more reduction in diameter than that of the LDPE-MWCNT filaments. More reduction in diameter means more elongation of the filament. This suggests that most of the elongation in case of UHMWPE-reinforced samples took place before the initiation of fracture. This explains the large fracture strain of LDPE-MWCNT-UHMWPE filaments, although having identical fracture process as that of LDPE-MWCNT. This also implies that addition of UHMWPE does not change the fracture process but instead allows larger elongation because of the sliding of the polymer interfaces. It is noticed in three categories of samples that smoothness of fiber surface improves as one move from (a) to (c). In a smoother surface, number of surface crazes is less and hence the chance of developing V-notches is less, therefore, delaying the initiation of failure process.

[FIGURE 7 OMITTED]

At multiple stages of strain hardening experiments, filament surfaces were examined by SEM. This was necessary to measure fiber diameter and calculate true stresses. During those observations we did not notice any initiation of fracture at the fiber surface even though some of those observations were made just before the final failure. Final failure event was quick and occurred within the last few milliseconds of the loading. This precluded any observation of two fracture stages as they progressed.

CONCLUSION

The study has demonstrated that a dual inclusion strategy of dispersing a second-phase polymer and an inorganic phase can be used to significantly enhance fracture toughness and elastic energy storage capacity of a textile polymer (LDPE) without any loss of strength or modulus. It is shown that addition of 8 wt% of UHMWPE and 2 wt% of MWCNTs can increase modulus of toughness and elastic energy storage capacity of LDPE by 55 and 17%, respectively. Such improvement in energy absorption features occurred with 39 and 44% increase in modulus and fracture strain without any loss of strength. XRD, DSC, and SEM characterizations have revealed that increase in crystallinity, crystallite size, and sliding between the minor and major phases were responsible for such improvement in energy absorption criteria. Investigation of fracture surface has illustrated that the fracture process is made of two distinct stages: first is the formation of V-notches, followed by a stage of rapid crack growth characterized by a rough and seemingly melted surface. It was also noticed that during first stage fracture, multiple V-notches grew in case of neat LDPE, whereas there was only one dominant V-notch with the other two cases. The second-stage fracture is identical in all three cases. Examination of fiber surfaces at various stages of loading revealed that final failure event is rapid and takes place in the last few milliseconds of the loading. Strain hardening experiments have demonstrated that modulus of toughness is reduced after repeated loading and unloading cycles, suggesting that higher draw ratio in solid state may not be beneficial for these categories of filaments.

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Mujibur R. Khan, (1) Hassan Mahfuz, (1) Theodora Leventouri, (2) Vijaya K. Rangari, (3) Andreas Kyriacou (2)

(1) Nanocomposites Laboratory, Ocean & Mechanical Engineering Department, Florida Atlantic University, Boca Raton, FL 33431

(2) Physics Department, Florida Atlantic University, Boca Raton, FL 33431

(3) Center for Advanced Materials, Tuskegee University, Tuskegee, AL 36088

Correspondence to: Hassan Mahfuz; e-mail: hmahfuz@fau.edu

Contract grant sponsor: National Science Foundation; contract grant number: HRD-976871.

DOI 10.1002/pen.21873

Published online in Wiley Online Library (wileyonlinelibrary.com).

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Author:Khan, Mujibur R.; Mahfuz, Hassan; Leventouri, Theodora; Rangari, Vijaya K.; Kyriacou, Andreas
Publication:Polymer Engineering and Science
Article Type:Report
Geographic Code:1USA
Date:Apr 1, 2011
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