multiscale reinforcement of polymers combining cellulosic wood pulp with layered silicate...

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Multiscale Reinforcement ofPolymers Combining CellulosicWood Pulp with Layered SilicateNanoplateletsLloyd A. Goettler a , Michal Benes b & Meagan Hill Marko ca Institute of Polymer Engineering, The University ofAkron, Akron, OH 44325-0301, USAb Institute of Polymer Engineering, The University ofAkron, Akron, OH 44325-0301, USAc Institute of Polymer Engineering, The University ofAkron, Akron, OH 44325-0301, USAPublished online: 02 Apr 2012.

To cite this article: Lloyd A. Goettler , Michal Benes & Meagan Hill Marko (2009)Multiscale Reinforcement of Polymers Combining Cellulosic Wood Pulp withLayered Silicate Nanoplatelets, Composite Interfaces, 16:7-9, 599-618, DOI:10.1163/092764409X12477406466426

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Composite Interfaces 16 (2009) 599–618www.brill.nl/ci

Multiscale Reinforcement of Polymers Combining CellulosicWood Pulp with Layered Silicate Nanoplatelets

Lloyd A. Goettler ∗, Michal Benes and Meagan Hill Marko

Institute of Polymer Engineering, The University of Akron, Akron, OH 44325-0301, USA

Received 7 June 2008; accepted 6 August 2008

AbstractSynergies resulting from the combination of discontinuous reinforcing elements at two different size scalesare examined in two polymer types — rubbery matrices comprising acrylonitrile-co-butadiene (NBR) elas-tomer and a high density polyethylene plastic. The latter is derived from a recycled post-consumer wastestream that is upgraded by the reinforcement to compensate for any degradation experienced during prioruse or impurities introduced during recycling. The two reinforcements are wood pulp at the microscale andexfoliating layered silicate clays at the nanoscale. Appropriate compatibilizing agents are employed to allowwetting of the reinforcement with the polymer matrix, promote dispersion and provide a strong interface.In general, the microscale elements provide mechanical strengthening in tension, while the nanoscale rein-forcements enhance stiffening and reduce failure propagation by tearing. The use of natural reinforcementsand recycled feedstocks imparts environmental acceptability to such formulations.© Koninklijke Brill NV, Leiden, 2009

KeywordsCellulose, montmorillonite, composite, multiscale, reinforcement, polyethylene, butadiene-co-acrylonitrile

1. Introduction

A considerable literature now exists on the reinforcement of polymers withnanoscale elements in many forms and dimensionalities. The most common andcommercially viable of these nanoscale reinforcements at the current time is thesmectic layered silicate, usually used in the form of montmorillonite clay [1]. Itis well known that these can be incorporated into polymer matrices via variousprocesses, including in situ polymerization, solution intercalation and melt interca-lation. This work utilizes the intercalation of neat polymer in the form of a melt orrubber.

It is further well known to reinforce polymers amongst other materials with dis-continuous inclusions of high aspect ratio, of which the most common are fibrous

* To whom correspondence should be addressed. E-mail: [email protected]

© Koninklijke Brill NV, Leiden, 2009 DOI:10.1163/092764409X12477406466426

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600 L. A. Goettler et al. / Composite Interfaces 16 (2009) 599–618

in nature. An advantage of their utilization in comparison to continuous reinforce-ment (yarn, cord, or fabric) lies in more economical processing, for example forthe manufacture of rubber articles such as hose and belting. Chopped glass fibersare commonly used in rigid plastics, but their brittle nature causes easy fracture tolow aspect ratio during processing, especially in polymer matrices of high viscosity.Natural cellulose fibers have become an attractive reinforcement because they areabundant, inexpensive, renewable, recyclable and biodegradable [2]. Other benefitsof cellulose fibers include low density and abrasiveness, high specific strength andadequate thermal stability [3]. While they confer enhanced tensile strength, theirrelatively low modulus limits their stiffening capabilities.

During processing of short fiber composites, the fibers tend to align under thecombined actions of shear and elongational deformations [4–8]. The resultinganisotropic mechanical properties are dependent on the direction and extent of fiberalignment [9]. For example, tears in rubbery matrices and cracks in hard plasticcomposites propagate easily parallel to the orientation direction of the reinforcingfibers, causing premature failure under transverse stresses [10].

Nanoplatelets derived from montmorillonite, a smectic clay with a three-layer crystalline structure that can be swollen and delaminated to provide two-dimensional platelet reinforcements of high aspect ratio with a micrometer planarsize and nanometer thickness, would impart distinctly different reinforcement andtear mechanics compared to traditional fillers [11] such as silica or reinforcementssuch as short fibers. For example, they confer less anisotropy than a fibrous geom-etry and furthermore fit easily within the polymer interstices between the alignedmicrofibers of a multiscale hybrid composite. The small nanometer size of theclay particles would allow them to function independently of the presence of thelarger discontinuous fiber reinforcement. Thus, they can serve to reduce crack ortear propagation between the fibers, increasing strength and elongation to fail overthe purely microscale fibrous composite. The high aspect ratio of several hundredfor the nanoplatelets deriving from montmorillonite clay along with their high ten-sile modulus of about 150–170 GPa allow them to provide substantial stiffening atlow volume concentration. Their lower strengthening capacity in plastic and car-bon black-reinforced rubbers can be compensated by the presence of microscalecellulose fibers, thus providing an optimal combination of reinforcing elements foroverall composite performance. An idealized schematic of such a multiscale hybridcomposite is shown in Fig. 1.

The naturally hydrophilic nature of the cellulose fibers and the montmorilloniteclay makes them both incompatible with the non-polar polymer matrices. Sur-face treatment and compatibilization are both required to produce high dispersion,interfacial strength and mechanical properties. Layered silicates are essentially hy-drophilic in their pristine state, but the space (gallery) between the platelet surfacesof layered silicate clays can be rendered partially hydrophobic through substitutionof ammonium cations with long hydrocarbon chains for the natural sodium ions:this makes the polar clay layers more compatible with polymer matrices and in-

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L. A. Goettler et al. / Composite Interfaces 16 (2009) 599–618 601

Figure 1. Schematic of multiscale hybrid composite.

creases the spacing between the clay layers in order to facilitate the subsequentintercalation of polymer molecules during compounding. Various types of claytreatment have been identified in patents by Usuki and Kato et al. [12, 13]. A widevariety of ammonium groups can be selected for clay treatment to generate the high-est enthalpic interaction with the degree of polarity present in the polymer matrixmolecule. A relatively polar rubber matrix such as poly(acrylonitrile-co-butadiene)(NBR) can provide a good environment for dispersion of clay with a polar treat-ment and, on the other hand, a non-polar matrix such as polyethylene (PE) mightprovide improved dispersion and compatibility for clay that has alkyl ammoniumswith long aliphatic molecules decorating the gallery.

The layered structure of treated montmorillonite clay allows separation of theplatelet particles and dispersion throughout the matrix when favorable thermody-namic and shear conditions are met during mixing. The degree of delamination anddispersion of the reinforcement is critical for mechanical property development. In-tercalation results in the polymer being present between swollen clay layers, withthe clay layers maintaining their registry in relation to the initial tactoid assembly.In an exfoliated structure, the polymer has separated the clay layers to the extentthat they can be considered as dispersed in the polymer matrix.

Upgrading of high density polyethylene (HDPE) post-consumer waste recyclestreams through reinforcement has been addressed by Hill [14, 15]. Recycled HDPE(RHDPE), the highest volume plastic in the recycling market, is generally reclaimedvia a mechanical process in which the plastic is ground, melted and reconverted intoproducts usually having less demanding performance requirements, such as drums,pails, flowerpots and plastic lumber used in decking, railing, fencing, doors, win-dow frames, siding and roofing. The cellulose fiber could itself also be a recycledmaterial. Wood–plastic composites usually comprising wood flour rather than highaspect ratio cellulose fibers are growing significantly at about 20% annually in theautomotive and plastic lumber markets, and their use with a layered silicate clayco-filler has been addressed by Yeh et al. [16] and by Matuana and Faruk [17],in polypropylene and poly(vinyl chloride), respectively. Maleic anhydride grafted

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polyolefins can serve as a common compatibilizer for both reinforcements in poly-olefinic matrices [11–20].

The part of this paper addressing the effect of organically treated layered sili-cates on the reinforcement of NBR rubber focuses on increasing the tear resistanceof aligned short fiber composites. More comprehensive data including tensile prop-erties of these and similar styrene–butadiene rubber composites are given by Benes[21] and Benes and Goettler [22]. The synthesis and characterization of rubber–claynanocomposites has been discussed in a review by Akelah and Salahuddin [23].Several methods have been utilized for rubber nanocomposite preparation, utilizinglatex, solvent and melt. For example, Zhang et al. [24] report a successful approachfor preparing styrene-co-butadiene rubber nanocomposites by mixing rubber latexwith a clay/water dispersion and coagulating the mixture. Okada et al. [25] treatmontmorillonite clay with an organic solution of amino-terminated butadiene acry-lonitrile liquid rubber. An alternative latex method for NBR is presented by Wuand Zhang [26]. Mülhaupt and Ganter [27] prepare rubber compounds based onbutadiene rubber and containing organophilic layered silicates during swelling. Ex-cellent dispersion of organoclay nanofillers in rubber matrixes was demonstrated.Al-Yamani and Goettler [28] built on the earlier work of Benes and Goettler [22]to produce rubber/clay nanocomposites with high degree of exfoliation. Nah [11]prepared NBR nanocomposite via melt compounding in a Banbury type of internalmixer. Here we utilize the melt compounding approach for clay incorporation.

2. Experimental

2.1. Rubber Composites

Acrylonitrile-butadiene copolymer (Nysyn 35-5 NBR; acrylonitrile content 35%;Zeon Chemicals Co.) was the matrix used to produce rubbery hybrid compos-ites. Wood cellulose microscale fibers treated with proprietary coatings for easeof dispersion (TCF) and having aspect ratios in the range 100–150 (SantowebW, Flexsys America, L.P.) were combined with a bonding system comprisinghexamethoxymethylmelamine, or HMMM, (Cyrez® 964, Cytec Industries Inc.) andresorcinol-formaldehyde (RF) novolac resin (P-170 resin, Akrochem Corp.) to pro-mote adhesion of the cellulose to the matrix. The nanoclay filler (Cloisite® 30B,Southern Clay Products Co.), hereafter referred to as type 30B organoclay, wasmontmorillonite treated with methyl, tallow, bis-2-hydroxyethyl, quaternary am-monium chloride, which is best suited for the fairly polar NBR polymer matrix.

Non-black NBR compounds were reinforced with various ratios of clay to fiber.The experimental plan incorporated four groups of compounds containing differentconcentrations of total reinforcement, while the composition within each group wasvaried on a substitutional basis between cellulose fiber and nanoclay so that the totalloading of reinforcement remained constant (Table 1).

To prepare NBR rubber composites, various amounts of clay and treated cellu-lose fiber were mixed in a Brabender Plasticorder internal mixer (C. W. Brabender

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Table 1.NBR formulations with varying clay/cellulose fiber ratio

Ingredient Mixa

A B C D E F G H I J K L M N O P Q

Clay 30B (vol%) 0 5 5 0 3 5 10 3 5 10 0 3 5 10 5 10 10TCF (vol%) 0 0 10 20 17 15 10 22 20 15 30 27 25 20 30 25 30∑

Reinforcement 0 5 15 20 20 20 20 25 25 25 30 30 30 30 35 35 40

a Basic recipe (phr): NBR 100, ZnO 5, Stearic acid 1.5, MBTS 1.5, TMTM 0.15, Antioxidant 1.5,HMMM 3.5, Resin P-170 2.5. Fill factor 70%. (MBTS = Benzothiazyl Disulfide, TMTM = Tetram-ethylthiuram monosulfide, HMMM = hexamethylmethoxymelamine).

Co.) at a rotor speed of 65 rev/min and a set temperature of 90◦C for 9 min. Thecure system was added to the mixer in the last minute of mixing. Orientation of thefibers was imposed by multiple passes of the composite sheet through the nip ofa two-roll mill. Composites were cured at 160◦C for 7–10 min, depending on thecure curve characteristics.

Die C and trouser tear specimens were die cut from vulcanized NBR compositesheets of 2 mm thickness and tested on an Instron 6021 testing machine at roomtemperature and crosshead speeds of 500 mm/min or 50 mm/min, respectively, ac-cordingly to ASTM D 624-00, such that the tear propagated in the weak directionparallel to the fiber orientation, which the inclusion of nanoparticles was intendedto strengthen. Values of tear strength were calculated from the measured force, F ,according to Ts = F/d , where d is the thickness of the tear [29].

2.2. Plastic Composites

2.2.1. MaterialsRHDPE recovered from detergent bottles and milk cartons, washed and regroundwithout modification was obtained from CAPCO Recycling, Inc. Virgin HDPEused as a control was Alathon M 6210® medium molecular weight high-densityhomopolymer from Equistar Chemicals. Maleic anhydride grafted PE (MA-g-PE)compatibilizer having a functionality of 1.0 wt% and a melt flow rate of 5 g/10 minat 190◦C was Polybond 3009® from the Crompton Corporation.

Hardwood (TC 2500) and softwood (TC 1004) CreaTech TC cellulose fibers ob-tained from CreaFill Fibers Corp. are highly pure with a cellulose content of at least99.5% in a dry state and average length × width of 900 × 20 and 3030 × 130 µm,respectively [30]. They are available either unlubricated or containing 2 wt% (soft-wood) or 3 wt% (hardwood) of a low molecular weight PE.

Treated and untreated montmorillonite was obtained from Southern Clay Prod-ucts. Cloisite® Na+ purified natural montmorillonite clay (hereafter referred toas clay type NA) is untreated, thus having only Na+ cations present in the claygalleries. Cloisite® 20A organoclay (to be called clay type 20A) comprises mont-

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morillonite that has been ion exchanged with dimethyl, dehydrogenated tallowquaternary ammonium ion.

All melt compounding of the composites was done in a Brabender Plasticordermixer at 160◦C and 60 rpm for a total mixing time of 15 min using roller-style mix-ing rotors. For the cellulose and clay homocomposites, the recycled polyethylene(RHDPE) was first added and allowed to melt for 2 min. Next, the MA-g-PE wasadded all at once and mixed with the RHDPE for 1 min. Finally, the reinforcementwas added over a two-minute interval and the compound was allowed to mix foran additional 10 min. In this scheme, the MA-g-PE already well dispersed in thepolymer matrix could promote the dispersion of the reinforcement as it was intro-duced. In the case of the multiscale hybrid composites, the nanoclay reinforcementwas introduced immediately prior to the cellulose fibers.

2.2.2. Composite FormulationsIn order to determine optimum loadings of montmorillonite clay along with theMA-g-PE compatibilizer, a central composite design was employed to vary theconcentrations of both clay and MA-g-PE over five levels from 0–3.0 vol% and0–8.0 wt%, respectively. Cellulose fiber microcomposites were prepared to opti-mize the selection of cellulose type, cellulose loading and MA-g-PE concentrationon the composite properties. Cellulose fibers were tested as softwood and hardwoodfibers with and without lubricant in a two-level factorial experimental design thatvaried the cellulose content from 11 to 22 vol% and the MA-g-PE from 0–4 wt%.

Clay/cellulose multiscale hybrid composites were prepared with 20A nanoclay,non-lubricated soft wood cellulose fibers and MA-g-PE compatibilizer in order todetermine the interactions of variables in this complex system, including reinforce-ment loading and MA-g-PE concentration. The various formulations prepared areshown in the data table of the Results section.

2.2.3. Characterization and TestingComposites compounded in the Brabender Plasticorder mixer were granulated us-ing a Weima granulator, followed by either injection molding into tensile bars orcompression molding into thin films. The injection molded tensile bars with a gaugelength of 7.6 mm were made using a DSM Research Micro-Injection Molding Ma-chine at a cylinder temperature of 220◦C, mold temperature of 55◦C and injectionpressure of 6 bar. The compression molded films were pressed to less than 1 mmusing a Carver press at a temperature of 160◦C.

Tensile testing was performed on ASTM D 638 Type II specimens usingan Instron 4204 mechanical testing machine at room temperature and a rate of12.5 mm/min. The elastic modulus was calculated using the initial tangent to thestress–strain curve and the tensile strength from the maximum load. Design Expertsoftware from Stat-Ease Inc. was used to analyze the experimental data and producecontour plots depicting the mechanical properties of the various composites.

Square specimens cut from compression molded samples less than 1 mm thickwere mounted for Wide Angle X-ray Diffraction (WAXD) such that the X-ray beam

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was directed at the compression molded surface in a reflection mode. Measurementswere conducted using the Rigaku X-ray generator over a range of 1.5 to 10 2θ val-ues. Intensity vs 2θ data were obtained and used along with Bragg’s law to calculatethe interlayer spacing of the clay in the hybrid composites.

Scanning Electron Microscopy (SEM) images were obtained using a Hitachi S-2150 scanning microscope on fractured composite surfaces generated by flexingtensile bars that had been submerged in liquid nitrogen for 5 min to failure. Prior tomounting, all samples were coated in silver using an Emitech K575x Turbo SputterCoater.

Differential Scanning Calorimetry (DSC) was performed for neat RHDPE, anRHDPE/clay composite containing 3 vol% silicate made with 20A organoclay anda hybrid composite with 3 vol% silicate and 11 vol% cellulose. Each compositewas made with 10 vol% MA-g-PE compatibilizer. The DSC was conducted on aTA Instruments Differential Scanning Calorimeter on a 12 mg sample with temper-ature ramping from room temperature to 190◦C at a rate of 10◦C/min. The heat ofcrystallinity of 100% crystallized HDPE used for calculating percent crystallinitywas 293 J/g [31].

3. Results

3.1. Rubber Composites

Tear strengths measured on trouser and ASTM Die C specimens in the directionparallel to the fiber orientation, which was more or less identical with the millingdirection, are presented in Fig. 2 for the NBR multiscale hybrid nanocomposites.While the trouser tear results were more reproducible, the two methods correlatequite well, providing a quantitative validation of the measurements.

The trouser tear data in Fig. 3 show a strong drop in tear strength followed bya plateau as the overall reinforcement loading is raised above 20% by volume at aconstant clay/cellulose ratio, due to greater stress concentrations that persist at thehigher concentration levels. Tear strength can be enhanced in all cases by increasingthe clay fraction of the total reinforcement. For example, tear strength almost dou-bles as 10 vol% nanoclay is added to a 10 vol% cellulose fiber composite. It should

Figure 2. Correlation between ASTM Die C and trouser tear tests.

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Figure 3. Trouser tear of nitrile rubber composite at various levels of reinforcement.

Figure 4. Contours of trouser tear strength as a function of cellulose fiber and clay concentrations.

be noted that a similar increase in tear strength with clay addition also obtains inblack NBR compounds, even though tensile strength is not similarly enhanced [21].The regression contours depicted in Fig. 4 show how the tear strength varies di-rectly with all combinations of cellulose fiber and nanoclay concentration. Theseare calculated from a polynomial model fit to the data points at cellulose fiber con-centrations �15 vol% presented in Table 1. It is linear in clay concentration with afirst order interaction between the clay and the cellulose fiber concentrations alongwith a quadratic term in the cellulose concentration. The fit with thirteen degreesof freedom generates an F -value of 14.7 indicating a statistical significance greaterthan 99.9% and r2 = 0.81. The standard error is 1.5 kN/m. A larger increase en-sues when the comparison is made on a substitution basis, as seen for example byfollowing the diagonal tie line representing 20 vol% total reinforcement in the fig-ure, which is nearly normal to the tear strength contours. Other parallel tie linesin the figure serve to depict total reinforcement concentrations at 25 and 30 vol%.Either mode of comparison proves the effectiveness of nanoclay addition as a teartoughener for short fiber reinforced rubber composites.

Fiber filled composites tested via a trouser tear specimen had a typical knotty tearmechanism. The SEM of a brittle fracture surface in Fig. 5 reveals cellulose fibers

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Figure 5. SEM of fiber embedded in rubber matrix (10 vol% clay, 10 vol% fibers).

Figure 6. SEM of torn surface of ASTM Die C specimen (10 vol% clay, 10 vol% fibers).

embedded in the rubber matrix and polymer adhesion to the fibers that is importantfor stress development in the composite. The torn surface depicted in the SEM ofFig. 6 shows breakage of some fibers along with bonded matrix.

3.2. Plastic Composites

3.2.1. Clay NanocompositesAfter incorporation into the RHDPE matrix, the average WAXD d-spacing of the20A organoclay is observed in Fig. 7 to be similar to that of the treated clay, whilethe degeneration of the clay peak into a broad shoulder suggests highly disorganizedclay structures with a range of d-spacings.

Experimental mechanical property data for the clay nanocomposites are givenin Table 2. In this and subsequent data tables and graphs, the term vol% silicaterefers to the volume occupied by the mineral portion of the treated organoclay. Thetensile strength increases slightly with the loading of clay in the composite andthe concentration of MA-g-PE compatibilizer, though it is independent of the latterat low clay concentration (<1.5 vol% silicate). Similar to the trends observed fortensile strength, at low clay levels (below 1 vol% silicate) the elastic modulus is

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Figure 7. WAXD spectrum for the nanoclay in the hybrid composite.

Table 2.Experimental data for clay RHDPE nanocomposites

Run vol% vol% Elastic modulus Tensile strength % elongationsilicate MA-g-PE (MPa) (MPa)

1 2.6 6.8 1540 32.7 3012 1.5 4.0 1462 31.6 4103 1.5 4.0 1218 31.5 5114 0.4 1.2 1117 29.7 1925 2.6 1.2 1156 31.9 3086 0.4 6.8 1247 29.7 2167 3.0 4.0 1122 33.3 1658 0.0 4.0 832 28.5 1699 1.5 0.0 1158 30.7 291

10 1.5 4.0 1226 31.6 51211 1.5 4.0 1198 31.4 46012 1.5 8.0 1347 32.0 534

also independent of the concentration of MA-g-PE compatibilizer and is increasedwith the further addition of clay and MA-g-PE at higher clay loadings. However,the build in elastic modulus begins to slow or even decrease at higher clay levels.

3.2.2. Cellulose MicrocompositesTensile data for the cellulose microcomposites are given in Table 3. Softwood pulpfibers were selected as the optimal cellulosic reinforcement because they are moreeconomical than hardwood fibers, they are effective in reinforcing without the use of

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Table 3.Experimental data for cellulose fiber RHDPE microcomposites

Run vol% vol% Lubricant Cellulose Elastic modulus Tensile strength % elongationcellulose MA-g-PE level type (MPa) (MPa)

1 11 2 No lubricant TC2500 1569 35.9 22.12 22 2 Lubricant TC2500 2682 44.7 4.83 22 2 No lubricant TC1004 3068 52.4 4.04 11 4 Lubricant TC2500 1509 36.6 15.95 11 4 No lubricant TC1004 1770 38.6 21.86 22 4 Lubricant TC1004 2054 35.1 5.57 22 4 No lubricant TC2500 2472 45.1 6.78 11 2 Lubricant TC1004 1305 30.9 12.89 22 0 No lubricant TC2500 1364 36.0 4.4

10 22 0 No lubricant TC1004 1328 37.3 3.411 22 0 Lubricant TC2500 1645 36.4 4.012 22 0 Lubricant TC1004 1894 27.7 4.1

lubricants and they can be derived from recycled materials, thereby providing an en-vironmental incentive. The addition of MA-g-PE lowers the tensile strength of soft-wood cellulose composites at fiber loadings below 19 vol% where reinforcement islow, while increasing tensile strength at higher cellulose loadings. As expected, at aconstant level of MA-g-PE, the addition of more cellulose reinforcement results inincreased tensile strength. The elastic modulus increases somewhat with addition ofcellulose fiber to the composite over the data range of 18.5–22 vol%. The presenceof MA-g-PE does not have any significant effect on the low strain modulus prop-erty for which interfacial bonding is not critical. As in the case of tensile strength,the addition of the MA-g-PE compatibilizer at low cellulose fiber loadings (below14 vol%), where reinforcement is not substantial, also has a deleterious effect onelastic modulus, presumably due to matrix dilution with lower weight polymer. Theelongation to fail was found to be solely dependent upon the cellulose loading forthis composite, with which it varies inversely. It is significantly lower in general forthe more highly reinforced cellulose composites compared to the clay composites.

The substantial increase in tensile strength and modulus upon introduction of theMA-g-PE compatibilizer implies a weak interaction between the cellulose fiber andthe uncompatibilized RHDPE matrix, causing weak interfaces and poor dispersion.The highest properties are observed at high levels of both MA-g-PE and cellulosewhere both sufficient MA-g-PE is present to compatibilize the system along withsufficient cellulose to provide effective reinforcement. The generic stress–straincurve presented in Fig. 8 for those conditions is characteristic of a well-bondedreinforced polymer.

3.2.3. Multiscale Hybrid CompositesMultiscale hybrid composites were made using 20A organoclay as the nanoscalereinforcement and TC1004 softwood cellulose fibers without lubricant as the mi-

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Figure 8. Shape of stress–strain curve for cellulose/RHDPE microcomposite.

Table 4.Experimental data for the multiscale hybrid RHDPE composites

Run vol% vol% vol% Elastic modulus Tensile strength % elongationMA-g-PE silicate cellulose (MPa) (MPa)

1 5.3 2 4.4 1615 32.6 372 10.5 2 4.4 1590 32.3 273 10.5 4 7.8 2330 33.6 64 5.3 4 4.4 1923 32.4 115 5.3 2 7.8 1966 33.6 116 10.5 2 7.8 2006 34.0 127 10.5 4 4.4 1745 32.7 108 5.3 4 7.8 2307 31.9 6

croscale reinforcement. The shape of the stress–strain curve for the hybrid compos-ite is similar to that pictured earlier for the cellulose fiber composite.

The data for the multiscale hybrid composite are shown in Table 4. Of course, thedata from the previous tables also constitute limiting cases for the hybrid composite,as well as the data point for pure matrix having a tensile strength of 28 MPa, tensilemodulus of 800 MPa and elongation to fail of 400%. Correlations of the tensileproperty data for the multiscale hybrid composites are covered in the Discussionsection.

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(a) (b)

Figure 9. SEM for the hybrid composite: (a) cellulose fiber; (b) clay.

The SEM of Fig. 9(a) shows incorporated cellulose fibers coated with RHDPE,suggesting effective compatibilization by the MA-g-PE. Figure 9(b) shows the claydispersion. The specks of clay dispersed evenly throughout the polymer provideevidence of adequate clay dispersion in the multiscale hybrid composite.

The crystallinity for the HDPE in the clay and hybrid composites determined byDSC is 55.5% with 3 vol% silicate (from 20A organoclay) and 10 vol% MA-g-PEpresent and 53.8% with 3 vol% silicate (from 20A organoclay), 10 vol% MA-g-PEand 11% cellulose fiber present, respectively. Both the crystallinity and the crys-tallization temperature are reduced slightly with addition of reinforcements in bothcomposites. Crystallinity is thus not a significant factor in the mechanical propertiesof the polyethylene composites.

4. Discussion of Multiscale Hybrid Composites

Regression models for the tensile strength and modulus linear in all of the com-position variables and their first order interaction terms can be fitted with an r2

of >0.9 to the composite data [14]. Contour plots are presented in the successivegraphs to show these compositional effects on the mechanical properties. For ten-sile strength, a correlation taken over the combined composite data incorporatingclay and/or softwood cellulose fiber without lubricant having 23 degrees of freedomgenerates an F -value of 15 that indicates a high statistical significance greater than99%. The standard error of this correlation is 2 MPa. A similar correlation of theelastic modulus data produces a standard error of 300 MPa with an F -value of 10also corresponding to a high statistical significance of greater than 99%.

Figure 10 correlates tensile strength against various levels of cellulose fiber andclay with 7.50 vol% MA-g-PE compatibilizer present. At all cellulose fiber lev-els, strength continues to increase by adding cellulose, although only marginallyso when the clay loading is also high, since clay presence at high concentrationis detrimental to strengthening. Overall, the contribution of cellulose to the ten-sile strength is stronger than the contribution of the clay. According to Fig. 11

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Figure 10. Tensile strength dependence on cellulose and clay loading with 7.5 vol% MA-g-PE.

Figure 11. Elastic modulus dependence on cellulose and clay loading with 7.5 vol% MA-g-PE.

depicting the elastic modulus over the same levels of cellulose and clay in themultiscale hybrid composite, the addition of either reinforcement increases theelastic modulus. However, on a volume basis, the effects of the mineral reinforce-ment on stiffening are greater than those imparted by the lower modulus cellulosefibers.

Variation of the tensile properties with the MA-g-PE compatibilizer concen-tration provides insight into the state of the clay and cellulose dispersion in the

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Figure 12. Tensile strength dependence on MA-g-PE and clay loading at 7.5 vol% cellulose.

composite and the interactions that take place at the polymer/reinforcement inter-face. For the levels of compatibilization employed in this study, the MA-g-PE itselfis found to have no significant effects on the mechanical properties of polymer ma-trix.

Figure 12 shows that higher concentrations of the MA-g-PE compatibilizer inthe multiscale hybrid composite result in increased tensile strength, though the rateof increase of tensile strength is greater at lower clay levels and decreases at higherclay levels. The concentration of MA-g-PE is seen to have a positive effect on ten-sile strength at all cellulose levels in Fig. 13. Figures 14 and 15 similarly show theeffects of MA-g-PE concentration on the elastic modulus. The addition of MA-g-PE is found to generally increase the elastic modulus, though the rate of increase isgreater at lower clay levels and at higher cellulose loadings, up to a limiting value.

At lower levels of MA-g-PE, the tensile strength and modulus both show an in-crease as the clay proportion is augmented in the multiscale hybrid composites. Athigher levels of MA-g-PE, the trend is reversed showing higher mechanical proper-ties with the presence of more cellulose in the composite. The crossover where thetrend is reversed, occurs at approximately 6 wt% MA-g-PE for the elastic modulus(Fig. 16) and over a wider range of MA-g-PE values (approximately 4–8 wt%) forthe tensile strength. This change in effect of the MA-g-PE is attributed to a differ-ent bonding activity of the MA-g-PE with the clay and cellulose components. Itis expected that the clay, being highly polar though swollen with non-polar tallowchains on the cationic surfactant, is more attracted to and even potentially reactivewith the maleic anhydride groups on the compatibilizer, thereby requiring a lowerconcentration to produce effective properties. At higher MA-g-PE concentrations,

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Figure 13. Tensile strength dependence on MA-g-PE and cellulose loading at 1.0 vol% clay silicate.

Figure 14. Elastic modulus dependence on MA-g-PE and clay loading at 7.5 vol% cellulose.

the dilution of the polymer matrix with this lower molecular weight functionalizedcompatibilizer would reduce properties were it not for a stronger interface producedagainst the cellulose fiber reinforcements.

Table 5 compares representative moduli and tensile strengths for the varioustypes of RHDPE composites presented above. Data selected from Tables 2–4 atan MA-g-PE compatibilizer level of 4–5 vol% are listed next to the respective com-

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Figure 15. Elastic modulus dependence on MA-g-PE and cellulose loading at 1.0 vol% clay silicate.

Figure 16. Elastic modulus dependence on MA-g-PE for the RHDPE multiscale hybrid composite.

Table 5.Illustration of mechanical reinforcement effects on RHDPE by clay silicate and/or cellulose fiber at4–5 vol% MA-g-PE compatibilizer

System description Elastic modulus Increase Tensile strength Increase(MPa) (%) (MPa) (%)

RHDPE 830 0 28.5 0RHDPE + 2.6 vol% nanoclay 1350 63 32.3 13RHDPE + 11 vol% cellulose fiber 1770 113 38.6 35RHDPE + 4% clay and 8% fiber 2310 178 31.9 12

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posite description. The data entry for the clay nanocomposite represents a mean oftwo data points taken at higher and lower compatibilizer concentrations. At only2.6 vol% loading, the clay nanocomposite provides a significant increase in elas-tic modulus due to its efficient stiffening, with a small accompanying increase intensile strength. The cellulose fiber microcomposite appears to offer more totalstiffening, but only at a higher loading of 11 vol% typical for that type of com-posite. At a lower volume fraction, it would provide a smaller increase in elasticmodulus as compared to the clay-only composite. However, the cellulose fiber iscapable of generating more than twice the tensile strength of the clay-only compos-ite, again at their typical loading levels. As mentioned, the tensile properties of theclay nanocomposites level off and even decrease if their loading is taken too high.The 4 vol% clay/7.8 vol% cellulose fiber presented in the last row of the table dis-plays a further increase in modulus due to the higher clay content despite reductionin cellulose concentration, highlighting the efficient stiffening imparted by the clayand the low stiffening contribution of the cellulose fiber. The concomitant reduc-tion in tensile strength accompanying the reduced cellulose concentration despitethe higher clay concentration reflects the importance of cellulose fiber for strengthdevelopment. Thus, the increase in modulus for the multiscale hybrid composite isoccurring primarily through reinforcement by the clay while increased strength re-sults from the cellulose reinforcement. The hybrid composition effectively providesa balance of modulus with strength, whereas the individual reinforcements canbest each provide only one of the properties. Thus the multiscale hybrid compositecould not only offer a viable approach for improving the mechanical properties ofRHDPE, but one which is also cost effective.

5. Summary

The synthesis and characterization of various hybrid composites with celluloseand/or clay multiscale reinforcements have elucidated the effects of reinforcementsize and type, reinforcement loading and compatibilizer concentration on mechan-ical properties. The tear strength of nitrile rubber composites is increased signifi-cantly. The multiscale hybrid composite is capable of producing RHDPE compos-ites with tensile strengths approaching 40 MPa and elastic moduli of over 2 GPa,which are competitive with virgin HDPE both economically and in performance.

6. Conclusions

Cellulose fibers serve as effective microscale reinforcement for polyethylene by in-creasing tensile strength and, to a lesser extent, elastic modulus at low loading. Type20A organoclay is a viable nanoscale reinforcement for this polymer by providinggood stiffening with lesser tensile strengthening. The clay–cellulose multiscale hy-brid composite represents an effective approach to achieving high modulus andhigh strength simultaneously. Through its combination with microscale cellulosefiber reinforcement in a hybrid rubber composite, the characteristically low tear

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strength imparted by the latter can be largely overcome. The MA-g-PE and RFresin/HMMM bonding system/compatibilizers effectively interact with both thecellulose and clay components of the multiscale hybrid RHDPE and NBR com-posites, respectively, to produce meaningful mechanical properties.

Acknowledgements

The authors gratefully acknowledge Dr. Otto Maender of Flexsys America L.P. andMr. Ed Schut of CreaFill Fibers Corporation for their helpful discussions and supplyof cellulose fibers.

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multiscale reinforcement of polymers combining cellulosic wood pulp with layered silicate...

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