polymers Article
Morphology Development and Mechanical Properties Variation during Cold-Drawing of Polyethylene-Clay Nanocomposite Fibers Bartolomeo Coppola *, Paola Scarfato, Loredana Incarnato and Luciano Di Maio Department of Industrial Engineering, University of Salerno, Via Giovanni Paolo II n. 132, 84084 Fisciano (SA), Italy;
[email protected] (P.S.);
[email protected] (L.I.);
[email protected] (L.D.M.) * Correspondence:
[email protected]; Tel.: +39-089-963-186 Academic Editors: Francesco Paolo La Mantia and Maria Chiara Mistretta Received: 17 May 2017; Accepted: 15 June 2017; Published: 20 June 2017
Abstract: In this work, the influence of composition and cold-drawing on nano- and micro-scale morphology and tensile mechanical properties of PE/organoclay nanocomposite fibers was investigated. Nanocomposites were prepared by melt compounding in a twin-screw extruder, using a maleic anhydride grafted linear low density polyethylene (LLDPE–g–MA) and an organomodified montmorillonite (Dellite 67G) at three different loadings (3, 5 and 10 wt %). Fibers were produced by a single-screw extruder and drawn at five draw ratios (DRs): 7.25, 10, 13.5, 16 and 19. All nanocomposites, characterized by XRD, SEM, TEM, and FT-IR techniques, showed an intercalated/exfoliated morphology. The study evidenced that the nanoclay presence significantly increases both elastic modulus (up to +115% for fibers containing 10 wt % of D67G) and drawability of as-spun nanocomposite fibers. Moreover, at fixed nanocomposite composition, the cold-drawing process increases fibers elastic modulus and tensile strength at increasing DRs. However, at high DRs, “face-to-edge” rearrangement phenomena of clay layers (i.e., clay layers tend to rotate and touch each other) arise in fibers at high nanoclay loadings. Finally, nanocomposite fibers show a lower diameter reduction during drawing, with respect to the plain system, and surface feature of adjustable roughness by controlling the composition and the drawing conditions. Keywords: fibers; nanocomposites; polyethylene (PE); drawing
1. Introduction The use of synthetic fibers is of widespread technological importance in different fields (textiles, sport goods, military and transportation sectors, concrete reinforcement, 3D printing by Fused Deposition Modeling (FDM), etc.) and attracts extensive research efforts aimed to continuously improving their properties and processability to produce high performance fibers [1–3]. Among the many classes of synthetic polymer fibers, the market is dominated by polyamides, polyesters, acrylic and polyolefins. In particular, polyethylene (PE) fibers have great interest for several applications due to their easy processability, low cost, good mechanical properties, chemical resistance, lightness and resistance to abrasion. Disadvantages include the need of long chains, to provide sufficient load transfer between the weakly interacting macromolecules, low dimensional stability and low tolerance for high temperatures, which tend to cause swelling. One possible strategy to overcome these issues, and extend the use of PE fibers to new processing technologies such as the additive manufacturing by FDM, can be the inclusion in the PE matrix of small amounts of nanosized reinforcing fillers (e.g., carbon nanotubes, organomodified layered clays, and nano-silicon dioxide). Among them, organomodified layered clays resulted particularly attractive from a technological point of view for the production of polymeric nanocomposites with unique Polymers 2017, 9, 235; doi:10.3390/polym9060235
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multifunctional properties. In fact, they are widespread available and cheap, can be delaminated into nanometric clay layers with high aspect-ratio into several matrices and provide the possibility of producing nanocomposite systems with tailored and balanced strength/stiffness/toughness and improved dimensional and thermal stability, fire retardancy and gas barrier characteristics, without detrimental effects on recyclability [4–16]. In addition, nanoparticles can be also used as synthetic fibers coating, in order to improve mechanical properties of interfaces in composite materials [17–21]. Compared to the traditional filler-reinforced systems, the improved properties of polymer nanocomposites are mainly due to the stronger interfacial interaction that occurs between the polymer matrix and the silicate layers. Depending on nanocomposite morphology and clay loading, different effects on mechanical performances are reported in literature: low content and intercalated/exfoliated morphology lead to properties increase [9,22,23], while higher content and/or hybrid morphology (i.e., the presence of both tactoids and intercalated zones) produce no increase or a decrease of mechanical properties [24,25]. Moreover, the nano-platelets affect also rheological properties that are of widespread importance for process stability of polymeric systems. Generally when an intercalated/exfoliated morphology is achieved, a network structure is formed and the interactions between clay layers and polymer chains result in the improvement of both processability and drawability, due to the increase of extensional viscosity [24,26–29]. Therefore, unlike what happens for micro-composites, the use of nanofillers allows performing processes in the molten state such as fiber spinning, producing nanocomposite fibers with enhanced performances and virtually helpful in improving additive manufacturing processes. However, remarkably for polyolefins, clay exfoliation is not easy to achieve due to the intrinsic incompatibility between hydrophilic layered silicates and hydrophobic polymers. The use of a compatibilizer, generally maleic anhydride, is thus necessary to overcome the poor compatibility between polyolefins and organoclays [12,30–33]. Even if several authors investigated the possibility of producing nanocomposite polyolefin fibers focusing on layered silicates as reinforcing filler, the combined effects of nanoclays addition and cold-drawing on the resulting morphology and fiber performances are rarely considered, particularly in compatibilized systems [24,25,34]. Within this context, we performed a study designed to investigate the effects of fiber composition and cold-drawing ratio on the evolution of the fiber nano- and micro-scale morphology (clay intercalation/exfoliation level, polymer and clay orientation, polymer crystallinity, and fiber surface defects), using a maleated PE as matrix. To this aim, PE/clay nanocomposite fibers at three different loadings of a layered organoclay (3, 5 and 10 wt %) were produced by melt compounding and extrusion, and subsequently cold-drawn at five draw ratios (7.25, 10, 13.5, 16 and 19). Morphological and mechanical properties of the nanocomposite systems were then analyzed by several techniques, i.e., X-ray diffraction (XRD), infra-red spectroscopy (FT-IR), electron and optical microscopy, and tensile mechanical testing, to evaluate the level of achieved benefits and fibers properties. 2. Materials and Methods 2.1. Materials A linear low density polyethylene grafted with maleic anhydride, Bondyram 4108 (Polyram Industries Ltd., Ram-On, Israel), was used as matrix for the nanocomposite production. The density and the melt flow index are, respectively, 0.92 g/cm3 and 6.5 g/10 min (190 ◦ C/10 kg) while the maleic anhydride (MA) content is 1 wt %. An organomodified montmorillonite, Dellite 67G (Laviosa Chimica Mineraria S.p.A, Livorno LI, Italy), functionalized with 48 wt % of di (hydrogenated tallow)–dimethylammonium (2HT) and 115 meq/100 g of cation exchange capacity, was used as filler for composites preparation.
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2.2. Preparation of the Nanocomposites Prior to processing, the matrix pellets (denoted PE) and the organoclay powder (denoted D67G) were dried at 70 ◦ C for 18 h under vacuum to remove moisture. The melt compounding was conducted using a twin-screw extruder (Collin GmbH-ZK 25, Ebersberg, Germany) with co-rotating intermeshing screws (Dscrew = 25 mm, L/D = 42). A temperature profile of 180–200–200–200–200–200–200–120 ◦ C from hopper to die was used and a screw speed of 350 rpm was set. Nanocomposite pellets (at 3, 5 and 10 wt % of silicate content) were produced by a laboratory pelletizer from the extruded strand which was quenched using a cold water bath. The samples will be called PE/xD67G where x stays for the clay percentage content. Fibers, from both matrix and nanocomposite pellets, were produced by a single screw extruder (Brabender Do-Corder E330, Dscrew = 20 mm, L/D = 20, Duisburg, Germany) operating at 180, 200, and 170 ◦ C, from hopper to die, and at 3 rpm. A capillary die (D = 0.5 mm, L = 10 mm) and a take up system with air cooling were used in order to produce undrawn filaments, with an average diameter ranging from 0.5 to 0.7 mm, which were collected on a take-up roll at 9.5 m/min. 2.3. Fibers Drawing Fibers were drawn starting from spun yarns by a universal testing machine (Sans CMT6000 series, Shanghai, China) at a crosshead speed of 4 mm/min and 30 mm of gauge length. Five different draw ratios (DRs) were investigated, considering the ratio between the final length and the initial one (i.e., DR = Lf /Lo ): 7.25, 10, 13.50, 16 and 19. Fibers drawing was performed at 25 ◦ C. 2.4. Characterization Small angle X-ray diffraction (SAXD) measurements were performed with a Bruker D8 Advance (Bruker, Billerica, MA, USA) diffractometer in the 2θ range 2–10◦ at a scanning rate of 0.2◦ /min. The interlayer distances were calculated using Bragg’s formula (Equation (1)): d = nλ/(2 sinθ)
(1)
where n is an integer and is equal to 1 for the first-order 001 X-ray reflection. Wide angle X-ray diffraction (WAXD) measurements were performed using a single crystal Bruker D8 Quest diffractometer (Bruker, Billerica, MA, USA), in the 2θ range 5–55◦ . A single fiber was mounted on a custom-made sample holder. Two-dimensional WAXD patterns were collected using a Photon II detector (Bruker, Billerica, MA, USA) and evaluated by APEX3 software (Bruker, Billerica, MA, USA). Exposure time for each pattern was 60 s. Both diffractometers were used operating with a Ni-filtered CuKα (λ = 0.15418 nm) radiation generated at a voltage of 35 kV and current of 40 mA. FT-IR (Fourier Transform Infrared Spectroscopy) measurements were carried out on as-spun and drawn fibers in the range of 4000–600 cm−1 , using a Nexus ThermoNicolet spectrometer (Waltham, MA, USA) equipped with a SmartPerformer accessory for attenuated total reflectance (ATR) analysis. Transmission Electron Microscopy (TEM) investigations were conducted in bright field on a FEI (Hillsboro, OH, USA) Tecnai G12 Spirit-Twin (120 kV, LaB6 ) equipped with a FEI Eagle 4K CCD camera at different magnifications. Prior to observations, samples were embedded in epoxy resin and oven cured at 50 ◦ C. Samples for TEM analysis were cut longitudinally (considering fiber axis) using an ultramicrotome Leica EM UC7 (Leica Microsystems, Wetzlar, Germany) under cryogenic conditions (nominal thickness of 100 nm), and then placed into copper grids. As-spun and drawn fibers were observed by a Scanning Electron Microscope (SEM, mod. EVO HD 15, Carl Zeiss AG, Jena, Germany) operating at 12 kV. To observe the obtained micro-scale morphology of drawn fibers, samples were chemically etched according the procedure described by Bassett et al. [35] and widely used to verify molecular orientation and defect formation [29,36]. Fibers were immersed into a solution at 1 w/v % of potassium permanganate in a mixture of concentrated sulfuric acid, orthophosphoric acid and distilled water (10:4:1 by volume) for 2 h at room temperature.
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room temperature. Then fibers were gold coated using an automatic sputter coater (mod. B7341, Then fibers were gold coated using an automatic sputter coater (mod. B7341, Agar Scientific, Stansted, Agar Scientific, Stansted, UK) for SEM observations. Samples were cut at three different positions of UK) for SEM observations. Samples were cut at three different positions of as-spun and drawn fibers as-spun and drawn fibers on three different samples, in order to ensure observations reproducibility. on three different samples, in order to ensure observations reproducibility. Tensile properties were determined using the same machine used for cold drawing (Sans Tensile properties were determined using the same machine used for cold drawing (Sans CMT6000 CMT6000 series, Shanghai, China) equipped with a load cell of 1 kN, according to ASTM C 1557-03. series, Shanghai, China) equipped with a load cell of 1 kN, according to ASTM C 1557-03. Fibers Fibers diameter was determined by an optical microscope (Zeiss Axioskop 40, Carl Zeiss AG, Jena, diameter was determined by an optical microscope (Zeiss Axioskop 40, Carl Zeiss AG, Jena, Germany) Germany) before the test. Tensile tests were performed at two crosshead speeds (4 and 40 mm/min) before the test. Tensile tests were performed at two crosshead speeds (4 and 40 mm/min) and fixed and fixed gauge length (30 mm) for determining elastic modulus (E) and properties at failure (tensile gauge length (30 mm) for determining elastic modulus (E) and properties at failure (tensile strength, strength, σb, and strain at failure, εb), respectively. σb , and strain at failure, εb ), respectively. 3. Results Results 3. 3.1. Morphology of As-Spun Fibers 3.1. investigated by XRD, TEM and ATR-FTIR ATR-FTIR to to investigate investigate Morphology of as-spun fibers was investigated polymer/clay interactions. Figure 1 shows XRD patterns of as received D67G organoclay and polymer/clay interactions. Figure 1 shows XRD patterns of as received D67G organoclay and as-spun ◦ ◦ as-spun fibers PE/clay at clay different clay loadings. D67G pattern shows three distinct PE/clay at fibers different loadings. The D67GThe pattern shows three distinct peaks at 2.67peaks , 5.04at ◦ (Figure 2.67°,7.35 5.04° and 7.35° (Figure 1). to According to Bragg’s formula(1)), (Equation (1)), three 2θ values and 1). According Bragg’s formula (Equation the three 2θ the values reported in reported in Figure at 1 the correspond the following interlayer distances: 1.8 and The 1.2 first nm, Figure 1 correspond followingat interlayer distances: 3.3, 1.8 and 1.2 nm, 3.3, respectively. respectively. The firstatpeak, corresponding the (001) reflection, is representative the regularly peak, corresponding the (001) reflection, isatrepresentative of the regularly stacked of layers; the high stacked layers; the high order reflections, corresponding to the (002) and (003) reflections, derive order reflections, corresponding to the (002) and (003) reflections, derive from the regular arrangement from regular arrangement of the PE/clay intercalant [37].itConsidering it canthe be three seen that in of thethe intercalant [37]. Considering fibers, can be seenPE/clay that in fibers, every case sharp every ofcase thedisappeared, three sharp peaks of disappeared, indicating the achievement of peaks D67G indicating the D67G achievement of intercalated/exfoliated structure with intercalated/exfoliated structure with interlayer higher interlayer higher than 4.4 nm (corresponding at 2θ valuethan < 2◦4.4 ). nm (corresponding at 2θ value < 2°).
Figure 1. 1. SAXD SAXD patterns patternsof oforiginal originalDellite Dellite67G 67G(D67G) (D67G)and andas-spun as-spun PE-based nanocomposite fibers Figure PE-based nanocomposite fibers at at different clay content (3, 5 and 10 wt %, respectively). different clay content (3, 5 and 10 wt %, respectively).
The organoclay organoclay delamination delamination was was verified verified also also by by ATR-FTIR ATR-FTIR that that allows allows relating relating the the increase increase of of The the dispersion level of the clay to the change in its Si–O bond absorption band [38]. As demonstrated the dispersion level of the clay to the change in its Si–O bond absorption band [38]. As demonstrated for several several natural natural and and organomodified organomodified clays clays (montmorillonite, (montmorillonite, bentonite bentonite and and hectorite), hectorite), Si–O Si–O bond bond for appears as as one one broad broad bond bond when when the the material material is is agglomerated agglomerated and and is is resolved resolved in in two two peaks peaks (in-plane (in-plane appears and out-of-plane) after delamination. The ATR-FTIR spectra for D67G, PE and PE/5D67G as-spun and out-of-plane) after delamination. The ATR-FTIR spectra for D67G, PE and PE/5D67G as-spun fibers are are compared compared in in Figure Figure 2. 2. The The main main peak peak assignments assignments are are listed listed in in Table Table 11 [39]. [39]. The The PE PE sample sample fibers shows all the characteristic peaks of polyethylene grafted with maleic anhydride (MA). The shows all the characteristic peaks of polyethylene grafted with maleic anhydride (MA). The appearance appearance of the carbonyl stretching bands and the ring stretching vibrations of saturated cyclic of the carbonyl stretching bands and the ring stretching vibrations of saturated cyclic five-membered five-membered anhydride confirms that MA grafted The D67G sample showsSi–O one broad anhydride confirms that MA is grafted ontoisPE. The onto D67GPE. sample shows one broad band Si–O band in the 1150–950 cm−1 region, where clays typically absorb IR radiation, and two low peaks in the 2950–2800 cm−1 region, corresponding to CH2 stretching due to the organic modifier. The
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in the 1150–950 cm−1 region, where clays typically absorb IR radiation, and two low peaks in the Polymers 2017, 9, x FOR PEER REVIEW 5 of 17 2950–2800 cm−1 region, corresponding to CH2 stretching due to the organic modifier. The PE/5D67G sample shows all the signals two components but in its spectrum. Si–O signal is resolved PE/5D67G sample shows all of theitssignals of its two components but in itsThe spectrum. The Si–O signal in is −1 and the in-plane band at 1045 and 1021 cm−1 , as two absorption bands: the out-of-plane at 1070 cm −1 resolved in two absorption bands: the out-of-plane at 1070 cm and the in-plane band at 1045 and a1021 consequence the clay layer delamination occurred during melt compounding. cm−1, as a of consequence of the clay layer delamination occurred during melt compounding.
Figure 2. spectra of Dellite 67G,67G, PE and as-spunas-spun fibers infibers the range 600–4000 Figure 2. ATR-FTIR ATR-FTIR spectra of Dellite PE PE/5D67G and PE/5D67G in the range −1 −1 −1 cm . Detail of the regions 1650–1850 cm and 800–1200 cm are also reported. − 1 − 1 − 1 600–4000 cm . Detail of the regions 1650–1850 cm and 800–1200 cm are also reported. Table 1. Absorption Absorption bands bands and and their their peak peak assignments. assignments. Table 1.
Bands (cm−1) Bands (cm−1 ) 3600–3200 2913 3600–3200 2847 2913 2847 1790 1790 1710 1710 1470 1470 1070 1070 1045, 1021 1045, 1021 919 919 729 729
Assignments –OH stretching vibration –OH vibration –CHstretching 2 asymmetric stretching –CH2 asymmetric stretching –CH2 symmetric stretching –CH2 symmetric stretching C=O symmetric stretching C=O symmetric stretching C=O symmetric stretching C=O symmetric stretching –CH bending deformation –CH 2 bending deformation 2 Si–O outout of plane bending Si–O of plane bending Si–O in plane bending Si–O in plane bending Ring stretching vibration of saturated cyclic anhydride Ring stretching vibration of saturated cyclic anhydride –CH 2 rocking vibration –CH2 rocking vibration Assignments
To To better better understand understand the obtained morphology in terms of clay distribution and orientation, TEM investigations were carried out on on undrawn undrawn fibers, fibers, on on sections sections parallel parallel to to the the fiber fiber axis. axis. As representative case, a micrograph as-spun fiber is reported in micrograph of of longitudinal longitudinal section section of of PE/5D67G PE/5D67G as-spun Figure 3. In addition, TEM confirms level of the layered silicate in confirms the the high high intercalation/exfoliation intercalation/exfoliation level the PE matrix and evidences that as-spun fibers display silicate layers partially oriented along flow direction, as a consequence of the shear deformation associated to the flow inside the extruder, as reported in literature literature on on other other nanocomposite nanocomposite systems systems [28,40–42]. [28,40–42].
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Figure 3. TEM image of as-spun PE/5D67G fiber (section parallel to fiber axis). Figure 3. TEM image of as-spun PE/5D67G fiber (section parallel to fiber axis). Figure 3. TEM image of as-spun PE/5D67G fiber (section parallel to fiber axis).
Wide angle X-ray diffraction (WAXD) patterns report the presence of the typical crystalline Wideofangle X-ray (WAXD) patterns report of typical structure PE with twodiffraction dominant peaks at 2θ angles of 21.6° the andpresence 23.9°, which correspond to the 110 Wide angle X-ray diffraction (WAXD) patterns report the presence of the the typical crystalline crystalline ◦ and 23.9◦ , which correspond to the 110 structure of PE with two dominant peaks at 2θ angles of 21.6 plane andofthe plane of an orthorhombic (Figure 4),which respectively. At increasing structure PE200 with two dominant peaks at 2θcrystal anglesstructure of 21.6° and 23.9°, correspond to the 110 plane and the 200 plane of an orthorhombic crystal structure (Figure 4), respectively. At increasing nanoclay content, a considerable decrease in the intensity and a slight broadening and shift towards plane and the 200 plane of an orthorhombic crystal structure (Figure 4), respectively. At increasing nanoclay content, a considerable decrease in the intensity and a slight broadening and shift towards lower 2θ of the 110 reflection is observed while a lower influence was registered for the 200 nanoclay content, a considerable decrease in the intensity and a slight broadening and shift towards lower 2θ of the 110 reflection is observed while a lower influence was registered for the 200 reflection. reflection. to the literature, the while presence of aninfluence intercalated/exfoliated lower 2θ ofAccording the 110 reflection is observed a lower was registeredmorphology, for the 200 According to the literature, the presence of an intercalated/exfoliated morphology, combined to the combined to the high interfacial area and adhesion between the polymer and the clay, hinders PE reflection. According to the literature, the presence of an intercalated/exfoliated morphology, high interfacial area and adhesion between the polymer and the clay, hinders PE chain mobility, having chain mobility, two effects: reduction ofbetween the crystallinity degree the peak combined to the having high interfacial area aand adhesion the polymer and (related the clay,tohinders PE two effects: a reduction the crystallinity degree (related thebroadening peak intensity) and dimension the intensity) and dimension of the crystallites (related to the the to peak and(related shift) [23,25,43]. chain mobility, having oftwo effects: a reduction of crystallinity degree to the of peak crystallitesand (related to the of peak shift) intensity) dimension thebroadening crystallites and (related to[23,25,43]. the peak broadening and shift) [23,25,43].
Figure 4. WAXD patterns of nanocomposite fibers. Figure 4. WAXD patterns of nanocomposite fibers.
Figure 4. Fibers WAXD patterns of nanocomposite fibers. 3.2. Mechanical Properties of As-Spun
3.2. Mechanical Properties of As-Spun Fibers To investigate the influence of nanoclay addition on mechanical performances of as-spun fibers, 3.2. Mechanical Properties of As-Spun Fibers tensile tests were performed. The effect of clay amountononmechanical the tensileperformances stress–strain curves of as-spun To investigate the influence of nanoclay addition of as-spun fibers, To investigate the influence of nanoclay addition on mechanical performances of as-spun fibers, fibers istests shown in performed. Figure 5. The effect of clay amount on the tensile stress–strain curves of as-spun tensile were tensile tests were performed. The effect of clay amount on the tensile stress–strain curves of as-spun fibers is shown in Figure 5. fibers is shown in Figure 5.
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Figure 5. Tensile stress (σ)/strain (ε) curves of as-spun PE, PE/3D67G, PE/5D67G and PE/10D67G
Figure 5. Tensile stress (σ)/strain (ε) curves of as-spun PE, PE/3D67G, PE/5D67G and fibers. Figure 5.fibers. Tensile stress (σ)/strain (ε) curves of as-spun PE, PE/3D67G, PE/5D67G and PE/10D67G PE/10D67G fibers. Nanocomposite fibers exhibit not only a considerable increase of the elastic modulus but also a
significant increasefibers of ductility. In fact, PE fibersincrease show a of lower strain at failure (about Nanocomposite exhibit not onlyas-spun a considerable the elastic modulus but also a Nanocomposite fibers exhibit not only a considerable increase of the elastic modulus but also a 1500%) compared to nanocomposite fibers. Even if, unfortunately, the maximum strain of 1500%) the significant increase of ductility. In fact, as-spun PE fibers show a lower strain at failure (about significant increase of ductility. In fact, as-spun PE fibers show a lower strain at failure (about apparatusnanocomposite (i.e., 1800%, corresponding 540 mm of elongation) did not lead failure of(i.e., compared fibers. Evenfibers. if,tounfortunately, the maximum of theto apparatus 1500%)tocompared to nanocomposite Even if, unfortunately, the strain maximum strain of the nanocomposite fibers, hence a conclusive analysis of their cold-drawability is not possible, and 1800%, corresponding to 540corresponding mm of elongation) not of lead to failure of fibers, of hence apparatus (i.e., 1800%, to 540didmm elongation) didnanocomposite not lead to failure nanoclay addition significantly postpones fiber failure. Similar toughening effect was described in a conclusive analysis of their cold-drawability is not possible, and nanoclay addition significantly nanocomposite fibers, hence a conclusive analysis of their cold-drawability is not possible, and literature by other authors for several polymer/clay nanocomposite systems with postpones fiber failure. Similar toughening effect was Similar described in literature other authors nanoclay addition significantly postpones fiber failure. toughening effectby was described in for intercalated/exfoliated morphology and was explained as a consequence of temporary polymer literature by other authors forsystems severalwith polymer/clay nanocomposite systems and withwas several polymer/clay nanocomposite intercalated/exfoliated morphology chains physical crosslinking created by the nanoclay and resulting in local regions of enhanced intercalated/exfoliated morphology and was explained as a consequence of created temporary polymer explained as asoconsequence of temporary polymer chains[29]. physical crosslinking by the nanoclay strength, that the growth of microdefects is retarded chains physical crosslinking created by the nanoclay and resulting in local regions of enhanced and resulting in local regions strength, sonanocomposite that the growth of microdefects is retarded Moreover, with respectoftoenhanced the neat PE fibers, the ones show also a lower slope of[29]. strength, so that growth of microdefects is retarded [29]. Moreover, withthe respect the neat PE fibers, nanocomposite ones show a lower slope of their stress/strain curves toafter yielding, whichthe implicates a reduction of thealso polymer chains Moreover, with respect to the neat PE fibers, the nanocomposite ones show also a lower slope of during the after tensile test due which to hindering effectaofreduction nanoclaysof[23–25,43]. theirorientation stress/strain curves yielding, implicates the polymer chains orientation their stress/strain curves after yielding, which implicates a reduction of the polymer chains Elastic modulus of to undrawn fibers at different nanoclay contents is reported in Figure 6. The during the tensile test due hindering effect of nanoclays [23–25,43]. orientation during the tensile test due to hindering effect of nanoclays [23–25,43]. influence of clay layers on mechanical properties was significant, corresponding an increase the 6. Elastic modulus of undrawn fibers at different nanoclay contents is to reported in of Figure Elastic modulus of undrawn fibers at different nanoclay contents is reported in Figure 6. The elastic modulus of 42%, 56% and 115% for 3, 5 and 10 wt % of nanoclay, respectively. Moreover, the The influence influenceofofclay claylayers layers mechanical properties significant, corresponding to an increase onon mechanical properties was was significant, corresponding to an increase of the of presence of a linear relationship between elastic modulus and clay content is clearly recognizable. the elastic modulusofof42%, 42%, 56% and 115% 5 and of nanoclay, respectively. Moreover, elastic modulus 56% and 115% for for 3, 5 3, and 10 wt10%wt of % nanoclay, respectively. Moreover, the The same behavior was observed also for stress at yield (Figure 6). the presence of a linear relationship between elastic modulus and clay content is clearly recognizable. presence of a linear relationship between elastic modulus and clay content is clearly recognizable. The The same behavior was observed also (Figure6). 6). same behavior was observed alsofor forstress stressat at yield yield (Figure
Figure 6. As-spun fibers elastic modulus and stress at yield correlated to the nanoclay wt % (0, 3, 5 and 10, respectively). Figure 6. As-spun fibers elastic modulus and stress yieldcorrelated correlatedtotothe thenanoclay nanoclaywt wt % % (0, (0, 3, 3, 55 and Figure 6. As-spun fibers elastic modulus and stress atat yield and 10, respectively). 10, respectively).
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3.3. Structuraland andMorphological MorphologicalCharacterization Characterization of Drawn Fibers 3.3. Structural Fibers The structure and morphologyofofnanocomposite nanocompositefibers fibersafter aftercold coldcold-drawing cold-drawingwere werestudied studiedby The structure and morphology by XRD, and microscopy investigations. WAXD patterns of asand spun and drawn fibers, at XRD, FT-IR FT-IR and microscopy investigations. WAXD patterns of as spun drawn fibers, at different different DRs, are reported in Figure 7. The fiber axis, indicated by an arrow in Figure 7, corresponds DRs, are reported in Figure 7. The fiber axis, indicated by an arrow in Figure 7, corresponds to the to the meridian The diffraction as spun samples showrings twocorresponding sharp rings meridian direction.direction. The diffraction patterns ofpatterns as spun of samples show two sharp to orthorhombic planes and (110) of PE, whichinare in all to corresponding orthorhombic lattice planes (200)lattice and (110) of (200) PE, which are non-oriented all non-oriented cases, superimposed cases, superimposed on the amorphous halo. on the amorphous halo. Noclay clayreflections reflectionscan canbe beobserved, observed, as as expected, expected, because because the No the nanocomposite nanocompositelow lowdiffraction diffraction angles are out of the measurement range (5°–55°). On the contrary, after drawing, strong angles are out of the measurement range (5◦ –55◦ ). On the contrary, after drawing, strongreflections reflections appear, at small angles (near the beam stopper) along equatorial lines for the nanocomposite fibers. appear, at small angles (near the beam stopper) along equatorial lines for the nanocomposite fibers. These reflections, more intense at increasing clay content, suggest the morphology modification These reflections, more intense at increasing clay content, suggest the morphology modification related related to orientation and to distance reduction of the of silicate layers during the cold drawing to orientation and to distance reduction of the of silicate layers during the cold drawing process, as also process, as also found in other polymer/layered clay nanocomposite fibers [28,42,44,45]. With found in other polymer/layered clay nanocomposite fibers [28,42,44,45]. With concern to the matrix concern to the matrix morphology, even at the lowest DR = 7.25, XRD patterns show the presence of morphology, even at the lowest DR = 7.25, XRD patterns show the presence of well oriented crystallites well oriented crystallites and poorly oriented material, related to the formation of oriented and poorly oriented material, related to the formation oriented semi-crystalline fibers (Figure 7). semi-crystalline fibers (Figure 7). At higher DRs, only of a slight increase of orientation is observed, as aAt higher DRs, only a slight increase of orientation is observed, as a consequence of a two-slip process consequence of a two-slip process (interlamellar and intralamellar deformations) as described in the (interlamellar and intralamellar deformations) as described in the literature [38,46,47]. literature [38,46,47].
Figure 7. WAXD patterns for the investigated composite fibers at the different DRs (as-spun, 7.25, 16 Figure 7. WAXD patterns for the investigated composite fibers at the different DRs (as-spun, 7.25, 16 and 19, respectively). and 19, respectively).
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To analyze in more depth the crystallographic deformation with strain, full width at To analyze in more depth the crystallographic deformation with strain, full width at half-maximum (FWHM) values along the equatorial line (Figure 8) for at the To analyze in more depthwere the determined crystallographic withline strain, full 8)width half-maximum (FWHM) values were determined alongdeformation the equatorial (Figure for the orthorhombic lattice planes values (110) and (200). Accordingalong to Sherrer’s formula, line at increasing FWHM, the half-maximum (FWHM) determined the equatorial (Figure 8) for the orthorhombic lattice planes (110) were and (200). According to Sherrer’s formula, at increasing FWHM, average thickness of crystallites of polyethylene decreases. As evident, the increase of FWHM in (110) orthorhombic lattice planes (110) and (200). According to Sherrer’s formula, increasing FWHM, the average thickness of crystallites of polyethylene decreases. As evident, the at increase of FWHM in reflections is prevalent than that of (200) reflections. Such behavior can be explained by the chain slip the average thickness of crystallites of polyethylene decreases. As evident, the increase of FWHM in (110) reflections is prevalent than that of (200) reflections. Such behavior can be explained by the in (110) polyethylene crystals, resulting fibrils formation, that occurs preferably inbe the (110) that reflections is prevalent thaninthat of (200) reflections. Such behavior can explained the chain slip in polyethylene crystals, resulting in fibrils formation, that occurs preferably inplanes theby(110) present weak van der Waals interactions rather than in the (200) planes that have more strong covalent chain polyethylene crystals, resulting in fibrilsrather formation, that the more (110) planesslip thatinpresent weak van der Waals interactions than in theoccurs (200) preferably planes thatin have bonds [47,48]. planes that present weak van der Waals interactions rather than in the (200) planes that have more strong covalent bonds [47–48]. Coherently previously WAXD patterns patterns (Figure (Figure7), 7),FWHM FWHMvalues valuesofof strong covalent to bonds [47–48]. Coherently towhat what previously observed observed from from WAXD Coherently to what previously observed from WAXD patterns (Figure 7), FWHM values ofa a PE/5D67G and PE/10D67G show a maximum for a DR of 16. On the contrary, PE/3D67G shows PE/5D67G and PE/10D67G show a maximum for a DR of 16. On the contrary, PE/3D67G shows PE/5D67G and PE/10D67G show a maximum for a DR of 16. On the contrary, PE/3D67G shows a progressive increase lower thickness thicknessof ofcrystallites. crystallites. progressive increaseofofFWHM FWHMatatincreasing increasingDRs, DRs,corresponding corresponding to a lower progressive increase of FWHM at increasing DRs, corresponding to a lower thickness of crystallites.
Figure8. 8.Changes Changes of of FWHM FWHM in in equatorial equatorial (110) Figure (110) and and (200) (200) reflection reflectionagainst againstdraw drawratio ratioofofthe the Figure 8. Changes of FWHM in equatorial (110) and (200) reflection against draw ratio of the investigated fibers. investigated fibers. investigated fibers.
The nanoclay morphology evolution related to cold drawing process was investigated by The morphology evolution related to cold cold10, drawing process was investigated by Thenanoclay nanoclay drawing process investigated ATR-FTIR and TEMmorphology analyses, asevolution shown in related Figuresto9 and respectively, forwas PE/5D67G fibers by at ATR-FTIR and TEM analyses, as shown in Figures 9 and 10, respectively, for PE/5D67G fibers ATR-FTIR and TEM as spectra shown ininFigures 10, respectively, for PE/5D67G atat different DRs. The analyses, ATR-FTIR Figure9 9andindicate the occurrence of clayfibers layers different DRs. The ATR-FTIR spectra in Figure indicate the occurrence of level, clay layers rearrangement different DRs. The ATR-FTIR inof9Figure indicate the occurrence clay layers rearrangement phenomena with a spectra reduction the clay9layer delamination as itof can be inferred phenomena with a reduction of the clay layer delamination level, as it can be inferred by the intensity rearrangement phenomena with a reduction of the clay layer delamination level, as it can be inferred by the intensity decrease and the tendency to the merging of the Si–O signals, at increasing DRs, decrease and the tendency to thethe merging of the Si–O signals, increasing DRs, by the intensity decrease and tendency to the merging the Si–O1021 signals, DRs, particularly relevant for the in-plane absorption bands (i.e., atofat 1045 and cm−1particularly ).at increasingrelevant − 1 −1 relevant for the in-plane bands forparticularly the in-plane absorption bands (i.e.,absorption at 1045 and 1021(i.e., cm at).1045 and 1021 cm ).
Figure 9. ATR-FTIR of PE/5D67G fibers at different draw ratios (i.e., as-spun, 7.25 and 19) in the −1. of Figure 9. ATR-FTIR PE/5D67G fibers different draw ratios as-spun, 19)the in range the Figure ATR-FTIR PE/5D67G fibers at at different draw ratios (i.e.,(i.e., as-spun, 7.257.25 andand 19) in range9.1200–800 cmof −1. − 1 range 1200–800 cm 1200–800 cm .
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The TEM TEM images images in in Figure Figure 10 10 demonstrate demonstrate that, that, with with respect respect to to the the as-spun as-spun system, system, at at low low DRs, DRs, The The TEM images in of Figure 10 demonstrate that, with respect tothe theextrusion as-spun system, at low DRs, there is an improvement clay orientation along the fiber axis (i.e., direction), while, at there is an improvement of clay orientation along the fiber axis (i.e., the extrusion direction), while, there is an improvement of clay orientation along the fiber axis (i.e., the extrusion direction), while, high DRs, there is the occurrence of self-assembling phenomena of theofclay whichwhich tend totend rotate at high DRs, there is the occurrence of self-assembling phenomena thelayers, clay layers, to at high DRs, there is the occurrence of self-assembling phenomena of the clay layers, which tend to and touch each other, due to electrostatic attractions between positively charged edges and negatively rotate and touch each other, due to electrostatic attractions between positively charged edges and rotate and touch other,indue electrostatic attractions positively charged edges and charged faces, as each described theto literature Such between a “face-to-edge” morphology is made negatively charged faces, as described in the [28,42,49]. literature [28,42,50]. Such a “face-to-edge” morphology negatively charged faces, as described in the gradually literature [28,42,50]. Such a “face-to-edge” morphology possible during cold-drawing process, which reduces fibers diameter, by the good is made possible during cold-drawing process, which gradually reduces fibers diameter, by adhesion the good is made possible during cold-drawing process, which gradually reduces fibers diameter, by the good between clay layers and polymer matrix achieved using a compatibilized PE matrix. adhesion between clay layers and polymer matrix achieved using a compatibilized PE matrix. adhesion between clay layers and polymer matrix achieved using a compatibilized PE matrix.
Figure 10. TEM images of longitudinal sections of PE/5D67G fibers: as-spun (Left); drawn at DR 7.25 Figure fibers: as-spun as-spun(Left); (left); drawn Figure 10. 10. TEM TEM images images of of longitudinal longitudinal sections sections of of PE/5D67G PE/5D67G fibers: drawn at at DR DR 7.25 7.25 (Center); and at DR 19 (Right), respectively. (center); and at DR 19 (right), respectively. (Center); and at DR 19 (Right), respectively.
Optical microscopy allowed to investigate the changes, at DRs and clay wt %, in the Optical at increasing increasing Optical microscopy microscopy allowed allowed to to investigate investigate the the changes, changes, at increasing DRs DRs and and clay clay wt wt %, %, in in the the fiber macroscale morphology in terms of diameter reduction. The relative diameter variations (Δd/d 0, fiber fiber macroscale macroscale morphology morphology in in terms terms of of diameter diameterreduction. reduction.The Therelative relativediameter diametervariations variations(∆d/d (Δd/d00,, %) for for all the the investigated fibers fibers at each each DR are are compared in in Figure 11. 11. The graph graph shows aa general general %) %) for all all the investigated investigated fibers at at each DR DR are compared compared in Figure Figure 11. The The graph shows shows a general reduction of the fiber diameter during cold-drawing: the decrease is continuous for PE and reduction fiber diameter during cold-drawing: the decrease is continuous for PE andfor PE/3D67G reduction ofofthe the fiber diameter during cold-drawing: the decrease is continuous PE and PE/3D67G systems, whereas reaches a plateau at DR 13.5 for PE/5D67G and PE/10D67G ones. systems, whereas reaches a plateau at DR 13.5 for at PE/5D67G ones. Moreover, ones. with PE/3D67G systems, whereas reaches a plateau DR 13.5 and for PE/10D67G PE/5D67G and PE/10D67G Moreover, with respect to the neat PE fibers, the relative diameter variation is progressively lower respect to the neat PE fibers, theneat relative diameter is progressively lower for nanocomposite Moreover, with respect to the PE fibers, the variation relative diameter variation is progressively lower for nanocomposite fibers at loading. increasing nanoclay loading. Again, the phenomenon is related fibers at increasing nanoclay Again, the phenomenon is related to the clay particles thatto actthe as for nanocomposite fibers at increasing nanoclay loading. Again, the phenomenon is related to the clay particles that act as a rigid body inside the PE matrix hindering the polymer chains mobility and aclay rigid body inside the PE matrix hindering the polymer chains mobility and reducing their alignment particles that act as a rigid body inside the PE matrix hindering the polymer chains mobility and reducing their alignment during stretching. during stretching. reducing their alignment during stretching.
Figure 11. Relative diameter variation after drawing for the investigated fibers (i.e., PE, PE/3D67G, Figure 11. Relative diameter variation after drawing for the investigated fibers (i.e., PE, PE/3D67G, PE/5D67G and PE/10D67G). Figure 11. Relative diameter variation after drawing for the investigated fibers (i.e., PE, PE/3D67G, PE/5D67G and PE/10D67G). PE/5D67G and PE/10D67G).
SEM investigations were performed to detect the effects of cold-drawing and nanoclay on the SEM investigations investigations were performed to detect the the effects of of cold-drawing cold-drawing and and nanoclay nanoclay on on the the SEM werefiber performed detect development of microscopic surface to defects. The effects fiber surface patterning, in fact, plays a key development of microscopic fiber surface defects. The fiber surface patterning, in fact, plays a key development of microscopic fiberofsurface defects. The fiber surface patterning, in fact, plays a key role on mechanical performance fiber reinforced composite systems since it strongly affects the role on on mechanical mechanical performance performance of of fiber fiber reinforced reinforced composite composite systems systems since since it it strongly strongly affects affects the the role fiber wettability and the fiber/matrix interfacial adhesion, which generally increase at increasing fiber wettabilityand andthe thefiber/matrix fiber/matrixinterfacial interfacial adhesion, which generally increase at increasing generally increase at increasing fiber fiber wettability roughness for many polymer systems, evenadhesion, if also thewhich surface chemistry contributes to regulate fiber roughness for many polymer systems, even if also the surface chemistry contributes to regulate roughness many polymer systems, even if also the surface contributes to regulate the the whole for behavior [51]. Since all the investigated systemschemistry at increasing DR showed similar the whole behavior [51]. all Since all the investigated systems at DR increasing DR showed similar whole behavior [50]. Since the investigated systems at increasing showed similar occurrence of occurrence of defect formation, the PE/3D67G was chosen as example. Figure 12 shows that the fiber occurrence of defect formation, the PE/3D67G was chosen as example. Figure 12 shows that the fiber surface is quite smooth in undrawn system and increases its roughness during drawing. In surface is quite smooth in undrawn system and increases its roughness during drawing. In particular, at intermediate DRs (Figure 12b,c), transversal band appearance and growing are particular, at intermediate DRs (Figure 12b,c), transversal band appearance and growing are
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defect formation, the PE/3D67G was chosen as example. Figure 12 shows that the fiber surface is quite smooth in undrawn system and increases its roughness during drawing. In particular, at intermediate DRs (Figure band appearance and growing are recognizable; at the highest Polymers12b,c), 2017, 9, xtransversal FOR PEER REVIEW 11 of 17DR, the transversal rows extend across the whole width of the fiber (Figure 12d). Such defects occur in systems Polymers 2017, 9, xat FOR PEER REVIEW of 17 recognizable; the highest DR, the transversal rows extend across the whole width of the11fiber having low polymer chains mobility and thus low capability of the macromolecules to be extended (Figure 12d). Such defects occur in systems having low polymer chains mobility and thus low and aligned duringatdrawing [29,35,51–54]. recognizable; the highest DR, to thebetransversal rows extendduring acrossdrawing the whole width of the fiber capability of the macromolecules extended and aligned [29,35,49,52–54]. (Figure 12d). Such defects occur in systems having low polymer chains mobility and thus low capability of the macromolecules to be extended and aligned during drawing [29,35,49,52–54].
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Figure 12. SEM micrographs (2000×) of PE/3D67G fibers surface: (a) as-spun fiber; (b) DR = 7.25; (c)
(c) Figure 12. SEM micrographs (2000 ×) of PE/3D67G fibers surface:(d) (a) as-spun fiber; (b) DR = 7.25; DR = 13.50; and (d) DR = 19. (c) DR = 13.50; and (d) DR = 19. Figure 12. SEM micrographs (2000×) of PE/3D67G fibers surface: (a) as-spun fiber; (b) DR = 7.25; (c) DR = 13.50; and DR = 19.of cavities on fibers surface can be recognized (Figure 13c,d). This kind Moreover, the (d) presence
Moreover, thedefect presence ofwell-known cavities onforfibers surface can be recognized (Figure 13c,d). This kind of superficial is also traditional micro-composites and has to be related to the Moreover, the presence of cavities on fibers surface can be recognized (Figure 13c,d). This kind of superficial defect is also well-known for traditional micro-composites and has to be related to formation, during fiber drawing, of clay agglomerates that act as stress concentrators promoting the of superficial defect is also well-known for traditional micro-composites and has to be related to the the formation, during fiber drawing, agglomerates act as concentrators promoting voids setup [48,55–56]. In addition,ofatclay increasing nanoclay that content, thestress surface roughness is more formation, fiber drawing, of clay agglomerates that act as stress concentrators promoting the pronounced. the voids setup during [48,55,56]. In addition, at increasing nanoclay content, the surface roughness is voids setup [48,55–56]. In addition, at increasing nanoclay content, the surface roughness is more more pronounced. pronounced.
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Figure 13. SEM micrographs (c)(2000×) of fibers surface at their respective (d)highest DR: (a) PE at DR 16; (b)PE/3D67G; (c) PE/5D67G; and (d) PE/10D67G, at DR 19. Figure 13. SEM micrographs (2000×) of fibers surface at their respective highest DR: (a) PE at DR 16; Figure 13. SEM micrographs (2000×) of fibers surface at their respective highest DR: (a) PE at DR 16; (b)PE/3D67G; (c) PE/5D67G; and (d) PE/10D67G, at DR 19.
(b) PE/3D67G; (c) PE/5D67G; and (d) PE/10D67G, at DR 19.
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To To analyze analyze more more in in depth depth fibers fibers surface surface topology, topology, after after fiber fiber etching etching SEM SEM images images were were taken taken to to recognize in the samples zones having different density of the polymer matrix. Figure 14 compares recognize in the samples zones having different density of the polymer matrix. Figure 14 compares the nanocomposite fibers fibers drawn drawn at at the the maximum maximum investigated investigated DR DR (i.e., (i.e., 19). 19). the micrographs micrographs obtained obtained on on nanocomposite The PE/3D67G fiber shows a morphology denoted in the literature as “Pisa structure” [35] and due The PE/3D67G fiber shows a morphology denoted in the literature as “Pisa structure” [35] and due to to voids voids opening opening due due to to the the etchant etchant and and occurring occurring in in low low density density zones. zones. The The density density deficiency deficiency of of these regions results from the macromolecular rearrangements related to the material transformation these regions results from the macromolecular rearrangements related to the material into a fibrillar structure. This morphology not morphology present in theisPE/5D67G andinPE/10D67G systems, transformation into a fibrillar structure. is This not present the PE/5D67G and which indicates a more homogenous distribution and thusdistribution a delayed transition a PE/10D67G systems, which indicates adensity more homogenous density and thus towards a delayed fibrillar structure, as a consequence of the higher amount of nanoclay in these systems that reduces the transition towards a fibrillar structure, as a consequence of the higher amount of nanoclay in these polymer chains mobility. These observations suggest that, at the considered with respect to the systems that reduces the polymer chains mobility. These observations suggestDR, that, at the considered PE/3D67G systems, the PE/3D67G PE/5D67Gsystems, and PE/10D67G fibersand are more far from the fiber rupture limit DR, with respect to the the PE/5D67G PE/10D67G fibers are more far from resulting in their improved drawability. the fiber rupture limit resulting in their improved drawability.
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(c) Figure 14. 14. SEM (b) (b) PE/5D67G; andand (c) Figure SEM micrographs micrographs (10000×) (10000×)ofofetched etchedfibers: fibers:(a)(a)PE/3D67G; PE/3D67G; PE/5D67G; PE/10D67G, at DR 19 (white arrow indicates the fiber axis). (c) PE/10D67G, at DR 19 (white arrow indicates the fiber axis).
3.4. Mechanical of Drawn Drawn Fibers Fibers 3.4. Mechanical Properties Properties of The effects effectsofof cold-drawing nanoclay content on mechanical fibers mechanical performance were The cold-drawing andand nanoclay content on fibers performance were assessed assessed carrying outtests. tensile tests. Nanoclays increases fibers ductility and nanocomposite carrying out tensile Nanoclays additionaddition increases fibers ductility and nanocomposite fibers fibers reached the maximum achievable value of crosshead displacement. For all nanocomposite reached the maximum achievable value of crosshead displacement. For all nanocomposite systems, systems, at increasing DR, higher tensile strengths are while achieved while areduction ductility reduction was at increasing DR, higher tensile strengths are achieved a ductility was observed, observed, as Figure shown 15 in for Figure for the PE/3D67G as a consequence defect formation that as shown in the 15 PE/3D67G fibers, as fibers, a consequence of defect of formation that initiates initiates premature tensile failure. Moreover, the fibers drawn at DR equal to 16 and 19 have premature tensile failure. Moreover, the fibers drawn at DR equal to 16 and 19 have approximately approximately thefailure, same asstrain at failure, as expected that they haveafter alsodrawing similar the same strain at expected considering that theyconsidering have also similar diameter diameter after drawing (Figure 11). (Figure 11). Fiber cold-drawing cold-drawing produces produces an an overall overall increase increase of of mechanical mechanical properties properties at at increasing increasing DRs, DRs, as as Fiber shown in Figure 16. The elastic modulus of neat PE increases more than that of nanocomposites shown in Figure 16. The elastic modulus of neat PE increases more than that of nanocomposites fibers. fibers. This behavior was expected since nanocomposite fibers at higher DRs show higher diameters This behavior was expected since nanocomposite fibers at higher DRs show higher diameters with with respect to neat PE fibers, confirming the nanoclay hindering effect. respect to neat PE fibers, confirming the nanoclay hindering effect.
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Figure Stress/straincurve curvefor forPE/3D67G PE/3D67G drawn 1616 and 19). Figure 15.15. Stress/strain drawnfibers fibersatatdifferent differentDRs DRs(7.25, (7.25,10, 10,13.5, 13.5, and 19). Figure 15. Stress/strain curve for PE/3D67G drawn fibers at different DRs (7.25, 10, 13.5, 16 and 19).
Figure 16. Increase of the elastic modulus at increasing draw ratios (DRs) of PE, PE/3D67G, Figure 16. Increase of the elastic modulus at increasing draw ratios (DRs) of PE, PE/3D67G, PE/5D67G and PE/10D67G fibers, standard deviation ±10%. Figure 16. Increase of the elastic modulus at increasing draw ratios (DRs) of PE, PE/3D67G, PE/5D67G PE/5D67G and PE/10D67G fibers, standard deviation ±10%. and PE/10D67G fibers, standard deviation ±10%.
However, in fact, nanocomposite fibers always have better performances, in terms of tensile However, in fact, nanocomposite fibers always have better performances, in terms of tensile strength, compared to PE fibers and significant differences between the blends are recognizable for However, in fact,tonanocomposite fibers always have between better performances, terms of tensile strength, compared PE fibers and significant differences the blends areinrecognizable for draw ratios higher than 10 (Figure 17). PE and PE/3D67G fibers exhibit a progressive increase of strength, compared PE 10 fibers and significant the blends are recognizable draw ratios higherto than (Figure 17). PE anddifferences PE/3D67G between fibers exhibit a progressive increase offor tensile strength with DR; on the contrary, PE/5D67G and PE/10D67G show a maximum in the tensile draw ratios higherwith thanDR; 10 (Figure 17). PE and PE/3D67G exhibitshow a progressive increase tensile tensile strength on the contrary, PE/5D67G and fibers PE/10D67G a maximum in theof tensile strength values at DRs 16 and 13.50, respectively. This behavior is coherent with the structural strength with DR; on the contrary, PE/5D67G and PE/10D67G show a maximum in the tensile strength strength values at DRs 16 and 13.50, respectively. This behavior is coherent with the structural variations previously observed, in terms of crystallite thickness and orientation and defect values at DRs 16 and 13.50, respectively. This coherent with the structural variations previously observed, in terms of behavior crystalliteis thickness and orientation andvariations defect formation. formation.observed, in terms of crystallite thickness and orientation and defect formation. previously
Figure 17. Tensile strength (σb) of PE, PE/3D67G, PE/5D67G and PE/10D67G drawn fibers at the Figure17.17.Tensile Tensile strength(σ(σ)b)ofofPE, PE, PE/3D67G, PE/5D67G PE/5D67G and fibers at at the Figure andPE/10D67G PE/10D67Gdrawn drawn fibers the investigated DRs strength (7.25, 10, 13.5, 19), standard deviation ±5%. b 16 andPE/3D67G, investigated DRs (7.25, 10, 13.5, 16 and 19), standard deviation ±5%. investigated DRs (7.25, 10, 13.5, 16 and 19), standard deviation ±5%.
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4. Conclusions The results obtained in this work allowed analyzing the relationships among composition, cold-drawing extent, morphology evolution and mechanical performance improvements in PE/clay nanocomposite fibers. All the as-spun fibers showed intercalated/exfoliated morphology and partial orientation of the nanoclay along the fiber axis even before drawing. This resulted in beneficial effects on the mechanical behavior of the as-spun fibers in terms of elastic modulus, which increases up to 115% for the more loaded system, and drawability, thanks to the efficient level of stress transfer between the polymer matrix and the nano-dispersed clay platelets. Morphological XRD and microscopy investigations evidenced that, in general, the cold-drawing process promotes the polymer chain orientation along the drawing direction. For neat PE and PE/3D67G fibers, the matrix orientation increases continuously with the DR. On the contrary, for PE/5D67G and PE/10D67G, a maximum level of orientation is obtained at intermediate DR due to hindering effect of the nanoclay on polymer chain mobility. Regarding the nanoclay, after an initial increase of the silicate layer orientation at low DRs, the cold-drawing determines the occurrence of silicate rearrangement phenomena with the development of a “face-to-edge” morphology, which was related to the good adhesion between clay layers and polymer matrix. These clay rearrangements reduce the polymer chains mobility and are responsible of the development of the surface micro-defects and of the delayed transition towards a fibrillar structure in more loaded systems (i.e., PE/5D67G and PE/10D67G), as inferred by SEM investigations. Because of the matrix orientation and clay rearrangements, a relevant increase of the tensile strength (up to ca. 600% for PE/3D67G and PE/5D67G), drawability and dimensional stability were obtained. The straightforward approach to combining organomodified layered silicate with maleated PE showed interesting performance benefits helpful for enhancement of additive manufacturing technologies such as Fused Deposition Modeling and other rapid prototyping methods. Acknowledgments: Authors would like to thank Gaetano Guerra for the valuable discussions and Mario Maggio for the help in the use of the X-ray diffractometer of the Department of Chemistry and Biology of University of Salerno. Author Contributions: Bartolomeo Coppola, Luciano Di Maio and Loredana Incarnato designed the experiments; Bartolomeo Coppola and Paola Scarfato performed the experiments; Bartolomeo Coppola, Paola Scarfato and Luciano Di Maio analyzed the data; and Bartolomeo Coppola, Paola Scarfato and Loredana Incarnato wrote the paper. Conflicts of Interest: The authors declare no conflict of interest.
References 1. 2. 3.
4. 5. 6.
Hearle, J.W. High-Performance Fibres; Woodhead Publishing: Cambridge, UK, 2001. Coppola, B.; Scarfato, P.; Incarnato, L.; Di Maio, L. Durability and mechanical properties of nanocomposite fiber reinforced concrete. City Saf. Energy 2014, 2, 127–136. [CrossRef] De Leon, A.C.; Chen, Q.; Palaganas, N.B.; Palaganas, J.O.; Manapat, J.; Advincula, R.C. High performance polymer nanocomposites for additive manufacturing applications. React. Funct. Polym. 2016, 103, 141–155. [CrossRef] Alateyah, A.I.; Dhakal, N.H.; Zhang, Z.Y. Processing, properties, and applications of polymer nanocomposites based on layer silicates: A review. Adv. Polym. Technol. 2013, 32, 21368. [CrossRef] Garofalo, E.; Fariello, M.L.; Di Maio, L.; Incarnato, L. Effect of biaxial drawing on morphology and properties of copolyamide nanocomposites produced by film blowing. Eur. Polym. J. 2013, 49, 80–89. [CrossRef] Di Maio, L.; Scarfato, P.; Milana, M.R.; Feliciani, R.; Denaro, M.; Padula, G.; Incarnato, L. Bionanocomposite polylactic acid/organoclay films: Functional properties and measurement of total and lactic acid specific migration. Packag. Technol. Sci. 2014, 27, 535–547. [CrossRef]
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7. 8. 9. 10.
11. 12.
13. 14.
15.
16.
17. 18. 19.
20.
21.
22. 23. 24. 25. 26. 27.
15 of 17
Russo, G.; Simon, G.; Incarnato, L. Correlation between rheological, mechanical, and barrier properties in new copolyamide-based nanocomposite films. Macromolecules 2006, 39, 3855–3864. [CrossRef] Ray, S.S.; Okamoto, M. Polymer/layered silicate nanocomposites: A review from preparation to processing. Prog. Polym. Sci. 2003, 28, 1539–1641. [CrossRef] Alexandre, M.; Dubois, P. Polymer-layered silicate nanocomposites: Preparation, properties and uses of a new class of materials. Mater. Sci. Eng. 2000, 28, 1–63. [CrossRef] Alexandre, M.; Dubois, P.; Sun, T.; Garces, J.M.; Jérôme, R. Polyethylene-layered silicate nanocomposites prepared by the polymerization-filling technique: Synthesis and mechanical properties. Polymer 2002, 43, 2123–2132. [CrossRef] D’Amato, M.; Dorigato, A.; Fambri, L.; Pegoretti, A. High performance polyethylene nanocomposite fibers. Express Polym. Lett. 2012, 6, 954. [CrossRef] Scarfato, P.; Incarnato, L.; Di Maio, L.; Dittrich, B.; Schartel, B. Influence of a novel organo-silylated clay on the morphology, thermal and burning behavior of low density polyethylene composites. Compos. B Eng. 2016, 98, 444–452. [CrossRef] Mistretta, M.; Morreale, M.; La Mantia, F. Thermomechanical degradation of polyethylene/polyamide 6 blend-clay nanocomposites. Polym. Degrad. Stab. 2014, 99, 61–67. [CrossRef] Russo, G.; Nicolais, V.; Di Maio, L.; Montesano, S.; Incarnato, L. Rheological and mechanical properties of nylon 6 nanocomposites submitted to reprocessing with single and twin screw extruders. Polym. Degrad. Stab. 2007, 92, 1925–1933. [CrossRef] Scarfato, P.; Di Maio, L.; Milana, M.R.; Giamberardini, S.; Denaro, M.; Incarnato, L. Study of performance properties, lactic acid specific migration and swelling by simulant of biodegradable poly(lactic acid)/nanoclay multilayer films for food packaging. Food Addit. Contam. A 2017. [CrossRef] [PubMed] Garofalo, E.; Scarfato, P.; Di Maio, L.; Incarnato, L. Tuning of co-extrusion processing conditions and film layout to optimize the performances of PA/PE multilayer nanocomposite films for food packaging. Polym. Compos. 2017. [CrossRef] Karger-Kocsis, J.; Mahmood, H.; Pegoretti, A. Recent advances in fiber/matrix interphase engineering for polymer composites. Prog. Mater. Sci. 2015, 73, 1–43. [CrossRef] Tsirka, K.; Foteinidis, G.; Dimos, K.; Tzounis, L.; Gournis, D.; Paipetis, A.S. Production of hierarchical all graphitic structures: A systematic study. J. Colloid Interface Sci. 2017, 487, 444–457. [CrossRef] [PubMed] Felisberto, M.; Tzounis, L.; Sacco, L.; Stamm, M.; Candal, R.; Rubiolo, G.H.; Goyanes, S. Carbon nanotubes grown on carbon fiber yarns by a low temperature CVD method: A significant enhancement of the interfacial adhesion between carbon fiber/epoxy matrix hierarchical composites. Compos. Commun. 2017, 3, 33–37. [CrossRef] Tzounis, L.; Liebscher, M.; Tzounis, A.; Petinakis, E.; Paipetis, A.S.; Mäder, E.; Stamm, M. CNT-grafted glass fibers as a smart tool for epoxy cure monitoring, UV-sensing and thermal energy harvesting in model composites. RSC Adv. 2016, 6, 55514–55525. [CrossRef] Coppola, B.; Di Maio, L.; Scarfato, P.; Incarnato, L. Use of polypropylene fibers coated with nano-silica particles into a cementitious mortar. In Polymer Processing with Resulting Morphology and Properties: Feet in the Present and Eyes at the Future, Proceedings of the GT70 International Conference, Salerno, Italy, 15–17 October 2015; Pantani, R., De Santis, F., Speranza, V., Eds.; AIP Publishing: Melville, NY, USA, 2015. Cricrì, G.; Garofalo, E.; Naddeo, F.; Incarnato, L. Stiffness constants prediction of nanocomposites using a periodic 3D-FEM model. J. Polym. Sci. 2012, 50, 207–220. [CrossRef] Joshi, M.; Shaw, M.; Butola, B. Studies on composite filaments from nanoclay reinforced polypropylene. Fibers Polym. 2004, 5, 59–67. [CrossRef] Lee, S.H.; Youn, J.R. Properties of polypropylene/layered-silicate nanocomposites and melt-spun fibers. J. Appl. Polym. Sci. 2008, 109, 1221–1231. [CrossRef] Rangasamy, L.; Shim, E.; Pourdeyhimi, B. Structure and tensile properties of nanoclay–polypropylene fibers produced by melt spinning. J. Appl. Polym. Sci. 2011, 121, 410–419. [CrossRef] La Mantia, F.P.; Arrigo, R.; Morreale, M. Effect of the orientation and rheological behaviour of biodegradable polymer nanocomposites. Eur. Polym. J. 2014, 54, 11–17. [CrossRef] Incarnato, L.; Scarfato, P.; Scatteia, L.; Acierno, D. Rheological behavior of new melt compounded copolyamide nanocomposites. Polymer 2004, 45, 3487–3496. [CrossRef]
Polymers 2017, 9, 235
28. 29. 30. 31. 32.
33. 34.
35. 36.
37. 38. 39. 40. 41. 42. 43. 44. 45. 46.
47. 48.
49. 50. 51.
16 of 17
Garofalo, E.; Russo, G.M.; Di Maio, L.; Incarnato, L. Study on the effect of uniaxial elongational flow on polyamide based nanocomposites. Macromol. Symp. 2007, 247, 110–119. [CrossRef] Chantrasakul, S.; Amornsakchai, T. High strength polyethylene fibers from high density polyethylene/organoclay composites. Polym. Eng. Sci. 2007, 47, 943–950. [CrossRef] Wang, K.H.; Choi, M.H.; Koo, C.M.; Choi, Y.S.; Chung, I.J. Synthesis and characterization of maleated polyethylene/clay nanocomposites. Polymer 2001, 42, 9819–9826. [CrossRef] Zhai, H.; Xu, W.; Guo, H.; Zhou, Z.; Shen, S.; Song, Q. Preparation and characterization of pe and PE–g–mah/montmorillonite nanocomposites. Eur. Polym. J. 2004, 40, 2539–2545. [CrossRef] Morawiec, J.; Pawlak, A.; Slouf, M.; Galeski, A.; Piorkowska, E.; Krasnikowa, N. Preparation and properties of compatibilized ldpe/organo-modified montmorillonite nanocomposites. Eur. Polym. J. 2005, 41, 1115–1122. [CrossRef] Pegoretti, A.; Dorigato, A.; Penati, A. Tensile mechanical response of polyethylene–clay nanocomposites. Express Polym. Lett. 2007, 1, 123–131. [CrossRef] Coppola, B.; Di Maio, L.; Scarfato, P.; Incarnato, L. Production and Characterization of Polyethylene/Organoclay Oriented Fiber. In Proceedings of the VIII International Conference on Times of Polymers and Composites, Ischia, Italy, 19–23 June 2016; D’Amore, A., Acierno, D., Grassia, L., Eds.; AIP Publishing: Melville, NY, USA, 2016. Amornsakchai, T.; Olley, R.; Bassett, D.; Al-Hussein, M.; Unwin, A.; Ward, I. On the influence of initial morphology on the internal structure of highly drawn polyethylene. Polymer 2000, 41, 8291–8298. [CrossRef] Liparoti, S.; Sorrentino, A.; Guzman, G.; Cakmak, M.; Titomanlio, G. Fast mold surface temperature evolution: Relevance of asymmetric surface heating for morphology of ipp molded samples. RSC Adv. 2015, 5, 36434–36448. [CrossRef] Galimberti, M. Rubber-Clay Nanocomposites: Science, technoLogy, and Applications; John Wiley & Sons: New York, NY, USA, 2011. IJdo, W.L.; Kemnetz, S.; Benderly, D. An infrared method to assess organoclay delamination and orientation in organoclay polymer nanocomposites. Polym. Eng. Sci. 2006, 46, 1031–1039. [CrossRef] Silverstein, R.M.; Webster, F.X.; Kiemle, D.J.; Bryce, D.L. Spectrometric Identification of Organic Compounds; John Wiley & Sons: New York, NY, USA, 2014. La Mantia, F.P.; Dintcheva, N.T.; Scaffaro, R.; Marino, R. Morphology and properties of polyethylene/clay nanocomposite drawn fibers. Macromol. Mater. Eng. 2008, 293, 83–91. [CrossRef] Guo, Z.; Hagström, B. Preparation of polypropylene/nanoclay composite fibers. Polym. Eng. Sci. 2013, 53, 2035–2044. [CrossRef] Pavliková, S.; Thomann, R.; Reichert, P.; Mülhaupt, R.; Marcinˇcin, A.; Borsig, E. Fiber spinning from poly (propylene)–organoclay nanocomposite. J. Appl. Polym. Sci. 2003, 89, 604–611. [CrossRef] Gopakumar, T.; Lee, J.; Kontopoulou, M.; Parent, J. Influence of clay exfoliation on the physical properties of montmorillonite/polyethylene composites. Polymer 2002, 43, 5483–5491. [CrossRef] Garofalo, E.; Russo, G.; Scarfato, P.; Incarnato, L. Nanostructural modifications of polyamide/mmt hybrids under isothermal and nonisothermal elongational flow. J. Polym. Sci. B 2009, 47, 981–993. [CrossRef] Di Maio, L.; Garofalo, E.; Scarfato, P.; Incarnato, L. Effect of polymer/organoclay composition on morphology and rheological properties of polylactide nanocomposites. Polym. Compos. 2015, 36, 1135–1144. [CrossRef] Hiss, R.; Hobeika, S.; Lynn, C.; Strobl, G. Network stretching, slip processes, and fragmentation of crystallites during uniaxial drawing of polyethylene and related copolymers. A comparative study. Macromolecules 1999, 32, 4390–4403. [CrossRef] Peterlin, A. Crystalline character in polymers. J. Polym. Sci. C 1965, 9, 61–89. [CrossRef] Wang, K.H.; Chung, I.J.; Jang, M.C.; Keum, J.K.; Song, H.H. Deformation behavior of polyethylene/silicate nanocomposites as studied by real-time wide-angle X-ray scattering. Macromolecules 2002, 35, 5529–5535. [CrossRef] Okamoto, M.; Nam, P.H.; Maiti, P.; Kotaka, T.; Hasegawa, N.; Usuki, A. A house of cards structure in polypropylene/clay nanocomposites under elongational flow. Nano Lett. 2001, 1, 295–298. [CrossRef] Mittal, K.L. Polymer Surface Modification: Relevance to Adhesion; CRC Press: Leiden, The Netherlands, 2004. Alonso, Y.; Martini, R.E.; Iannoni, A.; Terenzi, A.; Kenny, J.M.; Barbosa, S.E. Polyethylene/sepiolite fibers. Influence of drawing and nanofiller content on the crystal morphology and mechanical properties. Polym. Eng. Sci. 2015, 55, 1096–1103. [CrossRef]
Polymers 2017, 9, 235
52. 53. 54. 55. 56.
17 of 17
Amornsakchai, T.; Songtipya, P. On the influence of molecular weight and crystallisation condition on the development of defect in highly drawn polyethylene. Polymer 2002, 43, 4231–4236. [CrossRef] Amornsakchai, T.; Bassett, D.; Olley, R.; Unwin, A.; Ward, I. Remnant morphologies in highly-drawn polyethylene after annealing. Polymer 2001, 42, 4117–4126. [CrossRef] El Maaty, M.A.; Bassett, D.; Olley, R.; Dobb, M.; Tomka, J.; Wang, I.-C. On the formation of defects in drawn polypropylene fibres. Polymer 1996, 37, 213–218. [CrossRef] Zuiderduin, W.; Westzaan, C.; Huetink, J.; Gaymans, R. Toughening of polypropylene with calcium carbonate particles. Polymer 2003, 44, 261–275. [CrossRef] La Mantia, F.P.; Morreale, M.; Scaffaro, R.; Tulone, S. Rheological and mechanical behavior of ldpe/calcium carbonate nanocomposites and microcomposites. J. Appl. Polym. Sci. 2013, 127, 2544–2552. [CrossRef] © 2017 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).
