characterisation of fibre–polymer interactions and transcrystallinity in short...

Characterisation of fibre–polymer interactionsand transcrystallinity in short keratinfibre–polypropylene composites

J. R. Barone* and N. T. Gregoire

Short fibre reinforced composites were made from keratin fibres obtained from poultry feathers.

The matrix material was either polypropylene or a blend of polypropylene and maleic anhydride

modified polypropylene (MaPP). In general, the addition of MaPP to polypropylene (PP) did not

alter the tensile properties of the blend in a significant way. When not using MaPP, composites

had lower breaking stresses than samples without fibres. However, MaPP at concentrations of

.4 wt-% enhanced the breaking stress of the composites to above the value without fibres.

Concurrent thermal analysis using differential scanning calorimetry (DSC) showed a distinct

increase in the amount of transcrystallinity in the composites at MaPP concentrations greater than

4 wt-%. Scanning electron microscopy (SEM) revealed increased interactions between the

protein fibres and the PP/MaPP matrix. Annealing showed that more large polymer crystals near

the fibres were not enough to increase stress transfer across the interface.

Keywords: Composites, Polypropylene, Maleic anhydride, DSC, Properties

IntroductionAgricultural fibres, such as cellulose and protein fibres,provide an environmentally friendly reinforcementoption for commodity thermoplastic materials such aspolyethylene (PE)1,2 and polypropylene (PP)3–12 and area pathway to obtain materials from sustainableresources. Agricultural fibres typically are of lowerdensity than inorganic reinforcements such as talc, glassor calcium carbonate. The absolute physical propertiesof agricultural fibres might not be as high as inorganicreinforcements, but specific or density normalisedproperties are comparable with inorganic fillers.Agricultural fibres are polymeric and have densitieslower than inorganic fibres so synthetic polymers can bereinforced with little or no cost to unit weight.

The key to successful enhancement of polymers withfibres of higher modulus or strength is to achieve goodpolymer/fibre interaction.13 Good fibre/polymer interac-tion can be expected if the two components arechemically compatible. Plant based cellulosic fibres suchas wood, flax, jute and sisal, and inorganic fibres andfillers are hydrophilic. However, commodity thermo-plastics such as PP are hydrophobic. To get chemicalcompatibility between the two, additives such ascoupling agents or compatibilisers are used.1,2,7,13

Maleic anhydride has been used to increase the proper-ties of cellulose fibre reinforced PE1,2 and PP.4–9 In most

cases, ultimate strength of the composites increased atmaleic anhydride concentrations of 0.1–2%, dependingon the maleic anhydride substitution in the olefin andthe concentration of maleic anhydride modified olefin inthe composite. Scanning electron microscopy (SEM)revealed increased fibre/polymer interactions.1,4,5

Protein fibres are mixed hydrophobic/hydrophilicdepending on the amino acid sequence. Keratin fibresobtained from poultry feathers have ,60% hydrophobicamino acids in the amino acid sequence, with thebalance being hydrophilic amino acids.14 So it can beexpected that there will be some chemical compatibilitybetween PP and feather keratin fibres. Recently, Baroneand Schmidt showed that it was possible to increase thetensile properties of low density polyethylene (LDPE)using keratin feather fibre without the use of couplingagents or compatibilisers.15 SEM showed good fibre/polymer interaction.

Maleic anhydride modified polypropylene (MaPP)has been used in keratin feather fibre–PP composites.Bullions et al. used MaPP to increase interactionsbetween PP and a mix of keratin feather fibre and kraftpulp fibre.10 Schuster made MaPP–keratin feather fibrecomposites and observed an increase in properties overneat PP.11 The proposed interaction mechanism betweenMaPP and the fibres is hydrogen bonding betweenhydroxyl groups on the MaPP and fibre and covalentbonding between hydroxyl groups on the fibre and thecarbonyl on the maleic anhydride. The PP portion of theMaPP is free to entangle with the PP composite matrix.

Investigation of the fibre surface in PP and cellulosefibre composites shows a ‘transcrystalline’ layer devel-ops.7,8,12 The fibre surface serves as a nucleation site for

USDA/ARS/ANRI/EMBUL, BARC West, Bldg. 012, Rm. 1–3, 10300Baltimore Ave., Beltsville, MD 20705, USA

*Corresponding author, email [email protected]

� 2006 Institute of Materials, Minerals and MiningPublished by Maney on behalf of the InstituteReceived 19 June 2006; accepted 10 August 2006DOI 10.1179/174328906X146478 Plastics, Rubber and Composites 2006 VOL 35 NO 6/7 287

polymer crystals and the crystals that form at the fibresurface are different from the crystals formed in the bulkpolymer. This means that the polymer volume immedi-ately surrounding the fibres has properties different fromthe bulk polymer. The transcrystallinity can be affectedby shear deformation at the fibre/polymer interface, thechemical nature of the interface and the cooling rateafter processing.7,8,12,16 Joseph et al. performed a verycomplete study on the properties of the transcrystallinelayer in PP/sisal composites.7 The study showed that theaddition of a urethane compatibiliser initiated crystal-lisation at the fibre/polymer interface in a short amountof time, similar to using shear deformation at theinterface. Crystallisation at the interface occurred overmuch longer time periods in absence of shear deforma-tion or compatibiliser. The melting point and amount ofcrystallinity as studied by differential scanning calori-metry (DSC) increased with fibre and compatibilisercontent. Nunez et al. showed that a b crystalline fractionappeared at the PP/wood flour interface when cooledquickly from the melt.8 These researchers also showedthat the melting point and crystalline fraction increasedwith filler and MaPP content.

In the present paper, keratin feather fibre of 0.2 cmlength is incorporated into blends of PP and MaPP withthe MaPP content varying from 0 to 10 wt-%. Thephysical properties of the composites are characterisedusing tensile testing, DSC and SEM. Fibre–polymerinteractions are assessed in terms of the transcrystallinitythat develops at the interface.

Experimental

Keratin feather fibreKeratin feather fibre was obtained from Featherfiber(Nixa, MO) using a process patented by the UnitedStates Department of Agriculture (USDA).17 Thefeather fibre was semicrystalline and had a diameter of,5 mm and a density of 0.89 g cm23 (Ref. 15). A moredetailed description of feather fibre can be foundelsewhere.10,11,15,18 Fraser and MacRae report modulusvalues of 3.4 and 5.2 GPa and yield stress values of 100and 200 MPa at relative humidity values of 100 and 65%respectively for feather rachis.19 Modulus and strengthvalues of polyolefins are less than the feather values, soreinforcement of polyolefins with keratin feather fibre ispossible.

Fibres of 0.2 cm length were made by grindingfeather fibre using a Retsch ZM 1000 centrifugalgrinder. The rotational velocity of the instrument was15 000 rev min21 and contained a torque feedback so asnot to feed in too much material and overload themotor. The fibre was fed in slowly to avoid motoroverload and to minimise frictional heating of theinstrument and the fibre. Sieving analysis and SEMimaging of the fibres showed that the majority of thefibres were in the 0.2 cm length range with a fewpercentages of smaller ‘fine’ particles.

Composite preparationThe PP was ExxonMobil PP9505E1, a random copoly-mer PP with melt flow index (MFI)530 g/10 min at230uC and 2.16 kg. The MaPP is ExxonMobil ExxelorPO1020 and had a reported maleic anhydride substitu-tion of 0.5–1.0 wt-%. The fibre fraction of the compo-sites was held constant at 20 wt-%. The MaPP content

represented the amount of MaPP relative to the totalsample weight. The blends were prepared by mixing PP/MaPP in ratios of 100:0, 98:2, 96:4, 94:6, 92:8 and 90:10.Composites were prepared by dry blending the PP andMaPP at PP/MaPP ratios of 100:0, 97.5:2.5, 95:5,92.5:7.5, 90:10 and 87.5:12.5. The polymer was firstadded into a Brabender mixing head set at 160uC androtating at 75 rev min21. Immediately after adding thepolymer, 20 wt-% keratin feather fibre was added intothe mixing head. The total sample weight of eachcomposite was 40 g, which represented a degree of fill of70% of the total volume of the mixing head. The melttemperature was monitored independently. The polymercontrol samples all achieved melt temperatures of,174uC and the composites all achieved melt tempera-tures of ,183uC. The total mixing time for each samplewas 15 min.

Following mixing, each sample was sandwichedbetween Teflon coated aluminium foil and pressed intothree thin sheets in a Carver Press Autofour/30 Model4394 at 180uC and 1 MPa for 18 s. The film was thenremoved and cooled under an aluminium block until itreached room temperature. After pressing, each thin filmwas inspected to note feather fibre dispersion. Previouswork showed that pressing did not affect fibre disper-sion.15,20 Good dispersion was observed in all cases.

To prepare samples for testing, the three thin sheetswere cut into quarters, stacked on top of each other,sandwiched between Teflon coated aluminium foil, andpressed in the Carver Press at 180uC and 0.7 MPa for1.5 min. After pressing, the films were air cooled untilthey reached room temperature. The resulting plateswere ,0.3 cm in thickness. Type 4 dogbone samples fortesting according to ASTM D638 were machined fromthe plates.

Composite testingComposite samples were allowed to sit at ambientconditions for one week before testing. Uniaxial tensiletesting was performed using a Com-Ten Industries 95 RCTest System. The applied test speed was 2.5 cm min21.Four samples of each composite were tested.

Thermal analysisThe thermal properties of the polymers and compositeswere assessed using DSC. A TA Instruments DSC 910swas used according to the procedures outlined in ASTMD3417 and D3418. Sample sizes of ,5 mg were used in aN2 atmosphere. The first heating cycle proceeded from30 to 225uC at 10uC min21. The first heating cycle wasfollowed by a cooling cycle from 225uC back down to20uC. The cooling rate could not be controlled. Theheater simply shut off and the DSC cell cooled to 30uC.Although there may have been some variation in thecooling step from sample to sample, the room tempera-ture was always 20uC; so it was assumed that the coolingsteps were similar. A second heating cycle thenproceeded from 30 to 225uC. Peak assignments andareas were determined according to the ASTM proce-dures. The first heating cycle yielded information aboutthe state of the polymer as a function of processingconditions. The heat of fusion of the first heating cycleDQm,1 and the melting temperature during the firstheating cycle Tm,1 were determined because these valuesrelate to the crystalline structure of the as testedcomposite. The crystalline fraction of each polymer

Barone and Gregoire Characterisation of fibre–polymer interactions and transcrystallinity

288 Plastics, Rubber and Composites 2006 VOL 35 NO 6/7

and composite after the first heating cycle X1 wasdetermined from the DSC results using

X1~DQm,1

DQf (1{mf )(1)

where DQf is the theoretical heat of fusion of 100%crystalline PP, DQf5(5798 J mol21)/(42 g mol21)5138 J g21 (Ref. 21), and mf is the mass fraction of fibre.Joseph et al. determined the crystallinity of PPcomposites in a similar manner.7

Scanning electron microscopy (SEM)The fracture surfaces were excised from the failed tensilebars using a scalpel blade and transferred into amodified specimen carrier. The specimen carrier wasknown as an ‘indium vise’ because the dissected pieceswere clamped between sheets of indium metal andplunge cooled in liquid nitrogen to 2196uC. The cooledholder was then transferred to an Oxford CT1500 HFcryopreparation system attached to a Hitachi S-4100SEM. The sample temperature was raised to 290uC for10 min to remove surface water from the sample surface.The sample was then cooled to below 2120uC andcoated with ,5 nm of platinum metal using a magne-tron sputter coater. Coated samples were transferred tothe cold stage in the SEM at 2170uC and observed withan electron beam accelerating voltage of 2 kV.

Results and discussion

Effect of MaPP on tensile propertiesFigure 1 shows the elastic modulus of the PP/MaPPblends and PP/MaPP/20 wt-% keratin fibre composites.The percentage of MaPP relative to the total amount ofPP or PP/fibre was kept constant to note how thereplacement of PP with fibre affected tensile properties.However, on a phr basis, the samples had overlappingranges, 0–14 phr MaPP for the composites and 0–11 phr

MaPP for the blends. There was perhaps a minimalenhancement of the elastic modulus as a function ofincreasing MaPP concentration at a constant fibreloading. The peak stress and strain at peak stressshowed noticeable differences as MaPP concentrationincreased as shown in Figs. 2 and 3 respectively. Theaddition of MaPP to PP did not affect the peak stress.At low MaPP concentration, the composite peak stresswas lower than the PP/MaPP blends. However, the peakstress of the composites was increasing with increasing

1 Tensile modulus versus MaPP concentration for PP–

20 wt-% 0.1 cm keratin feather fibre composites

2 Peak stress versus MaPP concentration for PP–20 wt-%

0.1 cm keratin feather fibre composites

3 Strain at peak stress versus MaPP concentration for

PP–20 wt-% 0.1 cm keratin feather fibre composites


Plastics, Rubber and Composites 2006 VOL 35 NO 6/7 289

MaPP concentration. At MaPP concentrations .4 wt-%,the peak stress increased discontinuously to peak stressvalues above the PP/MaPP blends. There did notappear to be a strong dependence of stress on MaPPconcentration above 4 wt-%, indicating that this mayhave been the limiting value necessary to get optimalfibre–polymer interactions. It should be noted that thesamples with higher MaPP contents had some surface‘freckles’ after moulding and it was possible that thismay have contributed to a premature failure. Therefore,theoretical peak stress values could be higher than theones reported if samples could be prepared to avoidsurface anomalies.

In Fig. 3, the strain at peak stress did not appear to begreatly dependent on MaPP concentration for the PP/MaPP blends. The blends were more ductile than thefibre composites. For the composites, the strain imme-diately decreased upon introduction of MaPP but slowlyincreased with increasing MaPP back to the 0 wt-%MaPP value. This indicated that better fibre–polymerinteractions required more strain to break. Therefore,addition of MaPP increased not only the strength butalso the toughness of the composites.

Effect of MaPP on thermal propertiesThe effect of MaPP and fibres on PP Tm,1 and X1 isshown in Figs. 4 and 5 respectively. The PP/MaPPblends had increased Tm,1 up to 6 wt-% MaPP; then theTm,1 values decreased at ,6 wt-% MaPP. The fractionof crystallinity showed a similar trend with maximumvalues at ,2 wt-% MaPP. For pure MaPP, Tm,15

163.48uC and X150.57. Therefore, the PP–MaPP blendsreached a maximum in Tm,1 and X1 at 2–6 wt-% MaPP,with the maximum values being higher than purePP9505E1 or MaPP PO1020. At greater MaPP con-centrations, the PP matrix has inhibited crystallinityprobably because of a branched chain architecture thatthe blends assume.

For the composites, the melting temperature increasedquickly from 0 to 6 wt-% MaPP; then the meltingtemperature increase was much slower at MaPPconcentrations of .6 wt-%, i.e. the thermal behaviourappeared to reach a limiting value for Tm,1 withincreasing MaPP concentration. The PP–MaPP blendsand composites showed a similar trend in Tm,1 up to 4–6 wt-% MaPP. At greater MaPP concentration, thecomposites had greater X1 and Tm,1 values. Larger X1

and Tm,1 values meant that there were more largecrystals in the composites. Figure 5 shows that thecrystalline fraction of the composites increased up to 4–6 wt-% MaPP, and then there was a discontinuousincrease in the crystallinity of the composite after whichthe crystalline fraction remained relatively unchanged.The same dependence on MaPP was observed in thepeak stress characteristics of the composites. The fibreloading was kept constant for each composite, so thefibre surface area for crystal nucleation was the same.Liu et al. increased the wood flour content in compositesand saw increased crystallinity because of increasednucleating surface.22 Again, MaPP loading of ,4 wt-%was a critical point in the PPzMaPP/fibre composites.However, unlike the blends without fibres, an increasedcrystallinity was observed because of the presence offibres.

Observation of fibre/polymer interactionsFigure 6 shows the fracture surface of a 0 wt-% MaPPcomposite magnified 62000. At 0 wt-% MaPP, therewas little fibre–polymer interaction. There were voids inthe polymer matrix left by fibres that ‘pulled-out’ andvoids around identifiable fibres. The fibres had ‘clean’surfaces, indicating that the PP did not adhere to thefibre. Figure 7a shows a macroscopic view of thefracture surface of a 4 wt-% MaPP composite at6300. There were voids where ‘pulled-out’ fibres onceresided. Closer inspection of the surface at 61000 as

4 Melting temperature from DSC first heating cycle Tm,1

versus MaPP concentration5 Fraction crystallinity from DSC first heating cycle X1

versus MaPP concentration



shown in Fig. 7b shows that some fibres were stronglyadhered to the PP matrix. At 6 wt-% MaPP, the fibreswere also strongly adhered to the polymer matrix asshown in Fig. 8 for a fracture surface magnified 63000.There were not many voids from a macroscopic view.The micrographs showed that, concurrent to thethermal analysis and physical property data, a definitivetransition in the properties of the composites occurredafter 4 wt-% MaPP. This concentration was highenough for most of the fibres to bind to the matrixwhereas anything less did not provide enough materialfor all of the fibres to bind and there was still some fibrepull-out.

Effect of crystallinity on propertiesThe increased crystallinity observed in the PPzMaPP/fibre composites manifested as an increased stress tobreak. The data in Figs. 2, 4 and 5 suggested that anincreased population of large crystals resided near thefibres. SEM analysis showed increased fibre–polymerinteractions. The purpose of maleic anhydride is toincrease chemical interactions between the hydrophobicPP and hydrophilic fibres (in the case of cellulosic orglass fibres) or partially hydrophilic fibres (in the caseof keratin). However, the results obtained in thepresent study suggest that there may be a physicalmechanism of interaction manifesting through an

increased transcrystallinity at the fibre/polymer interfacethat results in an increased stress at break.23

To test this, annealing experiments were performed onPP and PP/20 wt-% 0.2 cm keratin fiber composites notcontaining MaPP. In this manner, pure PP was used butthe transcrystallinity affected through thermal treatmentrather than chemical treatment with maleic anhydride.PP9505E1 and a composite of PP9505E1 and 20 wt-%0.2 cm keratin feather fibre were prepared and annealedat 0.93Tm5135uC in a convection oven for 8 h. At theend of 8 h, the oven was turned off and allowed toslowly cool to room temperature. The results of theannealed PP and annealed PP/20 wt-% 0.2 cm fibrecomposites are plotted in Figs. 1–5 as open and closedsquares respectively. The DSC first heating cycle curvesfor the annealed and unannealed PP and PP compositeswithout MaPP are shown in Fig. 9. Analysis of the DSCcurves in Fig. 9 shows that annealing did not measur-ably affect the average melt temperature Tm,1, depictedas the middle of the peak and plotted in Fig. 4.However, annealing removed the low Tm,1 or smallcrystal size fraction thus narrowing the crystal sizedistribution. Addition of fibres further narrowed thecrystal size distribution after annealing. Annealing thepure PP increased the percentage of crystallinity asshown in Fig. 5. Annealing the PP–fibre compositeincreased the percentage of crystallinity even further to

7 Image (SEM) of tensile bar fracture surface for 4 wt-% MaPP composite at a 6300 and b 61000

6 Image (SEM) of tensile bar fracture surface for 0 wt-%

MaPP composite at 62000

8 Image (SEM) of tensile bar fracture surface for 6 wt-%

MaPP composite at 63000



the values obtained with MaPP and no annealing, soaddition of fibres again acts as a preferential nucleatingsurface. Although the crystallinity of an annealed com-posite was similar to that of a composite with .4 wt-%MaPP, there was no increase in peak stress. In fact, therewas a decrease in the peak stress of the annealedcomposite. Therefore, for PP–keratin fibre composites,an increased transcrystallinity and larger crystals near thefibre/polymer interface are not sufficient criteria for stresstransfer across the interface. While this feature of thetranscrystalline region has been shown to increasephysical fibre–polymer interactions and thereforeincrease stress to break in the case of PP, increasedchemical interactions, i.e. use of MaPP, is required to getincreased stress at break.23 It was possible that thecrystallinity values obtained in the annealed compositewere enough to overcome any chemical interactionsbetween the fibre and polymer and pulled the matrix fromthe fibre surface, as has been observed previously.23

The key to polymer composites is still efficient stresstransfer across the fibre/polymer interface throughincreased adhesion, so the higher yield stress fibres canbear more load. The role of the transcrystalline layer inincreasing the composite properties is still arguable.Wang and Hwang argue that after fibre debonding,larger transcrystalline layers provide more frictionbetween fibre and polymer.16 Research has shown thatthick transcrystalline layers have increased fibre–polymer interactions, which manifest as lower fragmentlengths in single fibre fragment tests.12 The results shownin the present study indicated that the increased peakstress of the composites with .4 wt-% MaPP may havebeen partially because of a thick transcrystalline region,i.e. more large crystals near the fibres. However, theincreased fibre/polymer adhesion from maleic anhy-dride, as shown first hand by the SEM images, createdmuch more efficient stress transfer across the interface sothe fibres could bear load.

ConclusionsAlthough an increased population of large crystals neara polymer/fibre interface can physically increase stress

transfer across that interface, for the PP compositesstudied here, increased chemical interactions are alsoneeded for this efficient stress transfer. This wasperformed using maleic anhydride modified PP. SEManalysis showed this increased interaction and itmanifested as increased stress at break. Thermal analysisshowed that increasing MaPP concentration gave anincreased crystallinity with larger crystals in the tran-scrystalline region. The presence of fibres furtherincreased crystallinity and resulted in larger crystalsnear the fibres. Best properties were observed at MaPPconcentrations of .4 wt-% but increasing MaPP furtherdid not have a marked effect on physical properties.

Acknowledgements

Mention of trade names or commercial products in thisarticle is solely for the purpose of providing specificinformation and does not imply recommendation orendorsement by the US Department of Agriculture. Thepresent paper is based on a presentation at theAutoPolymers workshop held in Charleston, SC, USAin October 2005.

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