Biocomposites from renewable resource and based on cellulose acetate, dioctyl phthalate and shortcurauá fibers were prepared by large scale extrusion and injection molding and their mechanical, mor-phological and thermal properties were studied as a function of plasticizer (dioctyl phthalate) and fibercontents, as well as chemical treatment of the fibers: treatment with NaOH solution or extraction withacetone. The chemical treatments of the fibers play an important role on the mechanical and thermalproperties, increasing the Young’s modulus (up to 50%), the thermal dimensional stability and the ther-mal conductivity (ca. 100%) and decreasing the impact strength (ca. 50%) of the composites in comparisonwith plasticized cellulose acetate. Plasticizer and fibers influence the properties of the biocomposites in
the opposite way. Thus the properties of complete and functional formulations of biocomposites of cellu-lose acetate, plasticizer and curauá fibers with potential of applications and produced by a conventionalpolymer processing such as extrusion and injection molding can be tailored by controlling the amountand the characteristics of the additives. Among semi-empirical models used to describe the mechanicalproperties, the Cox-Krenchel and ROM mathematical model showed to be more suitable to describe the
iocom
Young’s modulus of the b
. Introduction
Composites are multicomponent and multiphase materials thatombine the properties of their constituents in different ways.n general, the properties of a given component is preserved ornhanced in the composites. Thus the properties of composites areetermined by the properties of their components and mainly by
nterfacial adhesion and size, shape, orientation and distributionf the fillers in the matrix (Panigrahi et al., 2007). Polymer matrixomposites (PMC) represent one of the most important classes ofomposites. Among these, the polymer matrix composites rein-orced with fibers are technologically important because of theigh mechanical strength of the fibers compared with the matrixJoshi et al., 2004; Ragoubi et al., 2010). Usually polymer compositesroduced by continuous fiber pultrusion show better mechanicalroperties than those processed with short fibers (short fiber rein-orced composite – SFRC). However, SFRC are more versatile withespect to preparation methods (Joshi et al., 2004; Mano et al.,010).
Traditionally, polymer matrix composites are based on ther-oset resins such as unsaturated polyester and epoxy or on
olyolefins and synthetic fibers, for example glass or carbon fibers.
However, polymer composites (Joshi et al., 2004; Barkoula et al.,2010) and nanocomposites (Abdul Khalil et al., 2012) based on nat-ural fibers have been subject of interest for many researchers in theacademy and in the industry. Most of the recent research with PMCis related to composites of polyolefins and lignocellulosic fibers likebamboo (Phuong et al., 2010; Dagang et al., 2012), curauá (Manoet al., 2010; Castro et al., 2012), banana (Annie et al., 2008), abaca(Rahman et al., 2009) and others. These papers often show improve-ment of the interfacial properties by means of surface treatments ofthe fibers or by using coupling agents. For example, Rahman et al.reported that mercerization decreases the hydrophilicity of abacafibers, improving their adhesion with polypropylene (Rahman et al.,2009). Annie et al. showed the role of interfacial adhesion on theproperties of composites by reporting higher thermal conductivityof the composites of poly(propylene-g-maleic anhydride) and shortbanana fibers in comparison with composites of polypropyleneand the same fiber (Annie et al., 2008). In this case, the anhy-dride groups are responsible for interfacial adhesion resulting fromchemical reactions with the hydroxyls on the surface of the fibers. Inthe last decade, composites based on biodegradable polymers andfillers from renewable sources, known as biocomposites (Mohantyet al., 2000), have received special attention. A series of papers
on biocomposites has been published in recent years, for examplebiocomposites of poly(hydroxybutyrate) (PHB) and bamboo fibers(Krishnaprasad et al., 2009), poly(lactide); PHB and flax fibers (Ariaset al., 2013); poly(lactic acid) (PLA) and continuous fibers of kenaf
Nishino et al., 2003); PLA and jute (Plackett et al., 2003); cellulosecetate and hemp (Mohanty et al., 2004); cellulose acetate and cel-ulose nanowiskers (Zhen-Yu et al., 2013) and polyester and curauábers (Harnnecker et al., 2012).
Cellulose acetate (CA) represents an important class of ther-oplastic derived from cellulose. Cellulose acetate with different
egrees of acetylation are amorphous, biodegradable and non-oxic, presenting many applications as filters, membranes, films,ood packaging and capsules for drug delivery (Chandra and Rustgi,998). This class of polymers has been used as matrices for biocom-osites with lignocellulosic fibers such as hemp (Mohanty et al.,004) and flax (Taha and Ziegmann, 2006). An advantage of theseomposites over the composites based on polyolefins is the naturalnterfacial adhesion resulting from interactions between hydroxylsnd carbonyls of the lignocellulosic fibers and of the matrix (Tahand Ziegmann, 2006).
Curauá fibers are extracted from the leaves of curauá, a plantriginally from the Amazon. These fibers have been used in theutomotive and textile industry due to their high Young’s mod-lus in comparison with other lignocellulosic fibers (Mohantyt al., 2004). This makes curauá fibers promising for replac-ng synthetic fibers in composites (Castro et al., 2012; Zhen-Yut al., 2013). Recently, we reported the mechanical and thermalroperties of biocomposites based on cellulose acetate, curauábers and biodegradable plasticizers prepared using a bench scalextruder (15 cm3 capacity). These composites showed lower ther-al conductivity and higher heat capacity in comparison with theatrix (Gutiérrez et al., 2012a,b). Moreover, they combine typicalechanical properties of dense materials with thermal insulat-
ng characteristics of porous materials such as polyurethane orolystyrene foam and because of this they present a real poten-ial of applications. The present work has the purpose of scaling uphe processing of these biocomposites based on cellulose acetate,lasticizer and short curauá fibers, in order to develop a completend functional formulation with a wide potential of applicationshat could be produced industrially.
. Experimental
.1. Materials
Two commercial grades of cellulose acetates containing 20 wt%nd 30 wt% of dioctyl phthalate (DOP) as plasticizer (Tenite Acetate05, 39.8% acetyl and Mn = 30,000 g mol−1, supplied from Eastmanhemical Company, USA) were used to prepare the biocomposites.
Curauá fibers were obtained from EMBRAPA-PA, milled (aver-ge length (5 ± 1) mm) and dried (F). A portion of the curauá fibers
as subjected to extraction with acetone for 8 h in a Soxhlet appa-
atus, in order to remove any waxes, and air-dried at 100 ◦C for 24 hFA). Another portion of the curauá fibers was treated with 8 wt%queous NaOH solution for 1 h under sonication (2 g of fiber/100 mL
of NaOH solution), washed many times with fresh water until pHnear 7 and air-dried at 100 ◦C for 24 h (FM).
2.2. Preparation of biocomposites
Plasticized cellulose acetate and curauá fibers were fed intoa co-rotating and interpenetrating twin screw extruder (L/D = 44;D = 26 mm) from Coperion Werner Pfleiderer, model ZSK-26. Thefibers were fed to the plasticized cellulose acetate by the side feeder.The extrusion was conducted with degassing under the follow-ing conditions: temperature profile varying from 140 ◦C to 180 ◦Cfrom feed to die and rotation speed of 300 rpm (productivity of10 kg h−1). Table 1 summarizes the composition and nomenclatureused for the biocomposites. The composites were molded into ten-sile test bars (ASTM D638 – specimen type I) and impact resistancebars (ASTM D256) using an Arburg All Rounder M-250 injectionmachine. The temperature of the mold was 50 ◦C, cooling time of25 s, 120 MPa pressure and holding pressure of 70 MPa.
2.3. Characterization
Tensile tests were conducted in a DL2000 EMIC Universal Test-ing Machine with a cell of 5000 N and speed of 50 mm min−1 usingan extensometer (EMIC) to measure the strain. Specimens of bio-composites were conditioned for 72 h at 25 ◦C and 50% moisture,according to ASTM D638. Izod impact resistance tests of notchedinjection molded specimens were made in an EMIC pendulum type(2.7 J) machine, according to ASTM D256.
Thermal dynamic mechanical analyses (DMTA) were conductedon a Rheometrics Scientific DMTA V in the flexural mode (singlebeam cantilever), in the temperature range from −50 to 250 ◦C, fre-quency of 1 Hz, amplitude of 0.01% and heating rate of 2 ◦C min−1.The mean dimensions of the samples between the clamps were:thickness of 6 mm, width of 12 mm and length of 38 mm.
Real and apparent heat capacities were measured by modulateddifferential scanning calorimetry in a MDSC-2910 TA Instrumentsequipment, according to ASTM E1952 06 and following the recom-mendations for calibration from reference (Lopes and Felisberti,2004). The samples were isothermally held at each temperature(20, 40, 60, 80 and 100 ◦C) for 20 min with oscillation amplitudeof ±0.5 ◦C and period of 80 s. Two cylindrical specimens of eachsample with 5 mm of diameter and different heights (0.4 mm and5 mm) were used for each experiment. These cylinders were cutfrom injection molding tensile specimens. Real and apparent heatcapacities were used to calculate the thermal conductivities of thesamples. Thermal expansion coefficients of the composites weredetermined by dilatometry performed on a TA Instruments TMA
2940 thermomechanical analyzer. The analyses were carried from0 to 120 ◦C, at a heating rate of 5 ◦C min−1 using cylindrical speci-mens with 5 mm of diameter and around 7 mm of height, accordingto ASTM E831.
M.C. Gutiérrez et al. / Industrial Crops and Products 52 (2014) 363– 372 365
Fig. 1. Scanning electron microscopy micrographs of the biocomposites: (a) CA30D-10F, (b) CA30D-10FA, (c) CA30D-10FM, (d) CA30D-10F, (e) CA30D-10FA and (f) CA30D-1
fottcMits
3
3
fdimd
0FM.
The morphological analyses of the biocomposites were per-ormed on a JEOL JSM-6390 LV scanning electron microscopeperating at 20 kV. The surfaces of the fractures (perpendicularo the injection flow) resulting from impact resistance tests andhose obtained by cryo-fracturing were sputtered with carbon andoated with gold (Bal-Tec Med 020) prior analysis. A Skyscan-1074icro CT Scanner was used to determine the distribution of fibers
n the matrix of the injection molded specimens for impact resis-ance tests in the following conditions: scan width of 15 mm, expo-ure time of 1.02 s, step size of 1.800◦ and rotation angle of 360◦.
. Results and discussion
.1. Morphology of the biocomposites
The SEM micrographs of the surfaces of the fractures resultingrom impact resistance test of the biocomposites show fibers with
iameters ca. 100 �m (Fig. 1a, b and c) as well as fibrils with approx-
mately 5 �m diameter (Fig. 1d, e and f), distributed throughout theatrix. Therefore, a partial fibrillation occurred and this is possibly
ue to the flow and shear imposed by the extrusion and injection
processes and to the chemical treatment of the fibers This resultis different from that observed for biocomposites prepared using alaboratory scale extruder (Gutiérrez et al., 2012a,b), where almostcomplete fibrillation was observed as a consequence of highershear and longer processing times. Although the morphology ofthe composites seems not to be affected by the chemical treat-ment of the curauá fibers, the surface of the treated fibers presentsa cleaner aspect without particles (Fig. 1). Other researchers alsoobserved changes in the surface roughness of lignocellulosic fiberssubject to extraction with acetone or treatment with NaOH solu-tion (Valadez-Gonzalez et al., 1999; Ouajai et al., 2004). Moreover,the micrographs of the cryofractured surfaces (Fig. 2) for the bio-composite CA20D-20F show that most fibers were cross sectionedand they are aligned in the injection flow direction. Similar resultswere observed for all composites.
Microtomographs for composites CA30D-10F (Fig. 3a and b) andCA30D-10FA (Fig. 3c and d) show the uniform volumetric distribu-
tion of the fibers in the cross section of the injection molded testsamples (3 mm × 13 mm dimension). Comparing both samples onecan observe that the composite with acetone extracted fibers has amore uniform fiber distribution.
366 M.C. Gutiérrez et al. / Industrial Crops and Products 52 (2014) 363– 372
ed sur
3fi
sscirvhfiYoF4c((ifieotDifi
Fi
Fig. 2. SEM micrographs of cryo-fractur
.2. Mechanical reinforcement of cellulose acetate with curauábers
Young’s modulus, tensile strength, strain at break and impacttrength for plasticized cellulose acetate and biocomposites arehown in Table 1. The increase of plasticizer concentration in theellulose acetate from 20 wt% to 30 wt% produces a 100% decreasen Young’s modulus and a 120% increase in impact strength,espectively. These results can be explained based on the freeolume argument: larger plasticizer concentrations result in aigher free volume. The progressive addition of pristine curauáber to the plasticized CA results in a systematic increase in theoung’s modulus and in a decrease in the impact strength, asbserved for other composites of thermoplastics and vegetal fibers.or example, Taha and Ziegmann reported Young’s modulus of.4 GPa and 4.0 GPa for biocomposites of cellulose acetate andontinuous fibers of hemp or flax (20% in volume), respectivelyTaha and Ziegmann, 2006). These values are comparable to that4.2 ± 0.5 GPa) observed in the present work for the biocompos-te containing 20 wt% of DOP and 20 wt% (21 vol%) of short curauábers. The plasticizer and filler effects are counter balanced. Forxample, biocomposites containing 30 wt% of DOP and 10 or 20 wt%f fibers present Young’s modulus similar to cellulose acetate con-
aining 20 wt% of DOP. And biocomposites containing 30 wt% ofOP and 10 wt% of treated fibers present Young’s modulus sim-
lar to composite containing 20 wt% of DOP and 10% of pristineber.
ig. 3. Microtomographs of the injection molded test samples of the biocomposites: (a) Cn two-dimensions and (d) CA30D-10FA in three-dimensions. Dimensions of the cross-se
faces for the biocomposite CA20D-20F.
In general, the mechanical properties of biocomposites are influ-enced by the chemical treatment of the fibers. In the present work,the chemical treatment resulted in a ca. 60% increase of the Young’smodulus and ca. 54% decrease of the impact resistance for bio-composites containing 30 wt% of DOP and 10 wt% of treated fibers(FA and FM) in comparison with the biocomposite with the samecomposition containing pristine fibers. A large number of articlesreporting either, the increase (Castro et al., 2012; Mano et al., 2010;Phuong et al., 2010) or the decrease (Goriparthi et al., 2012; Tomlalet al., 2010) of the impact strength due to the treatment of thefibers can be found in the literature. However, the nature of thechemical treatment (extraction with acetone or NaOH solution)does not influence the tensile and the impact resistance proper-ties of the biocomposites. This improvement of the fiber–matrixinteraction could also be responsible for the more uniform dis-tribution of treated fibers throughout the matrix as observed inthe microtomographs of Fig. 3. The mechanical properties of thecomposites should also reflect the properties of the fibers them-selves. For example, Young’s modulus of the pristine fibers andthose treated with acetone and NaOH solution is (53 ± 8), (48 ± 5)and (81 ± 7) GPa, respectively (Gutiérrez et al., 2012a,b). The highermodulus of the fibers subjected to treatment with NaOH solutionresults from an increase in the crystalline fraction of cellulose in the
fibers as determined by X-ray analysis: 59% for pristine and 70% forNaOH solution treated fibers (Gutiérrez et al., 2012a,b).
Semi-empirical models have been successfully used to describethe mechanical properties of composites. Serial (Eq. (1)), Modified
A30D-10F in two-dimensions, (b) CA30D-10F in three-dimensions, (c) CA30D-10FAction is 13 mm × 3 mm.
M.C. Gutiérrez et al. / Industrial Crops and Products 52 (2014) 363– 372 367
10 15 20 25 30
4
6
8
You
ng M
odul
us (P
a)
% in weig ht of Curaua fib ers
a)
10 15 20 25 30
2
4
6
You
ng M
odul
us (P
a)
% in weig ht of Curaua fib ers
b)
F es wit( h �o =
raY2
E
E
E
E
E
Icvco2aVcfppbdmvfitaamtotf
L ˇL/2
with
ig. 4. Young’s modulus as a function of short curaua fibers content for biocomposit· · ·) Hirsch model with k = 0.12, (·–·) Maxwell model, (··–··) Cox Krenchel model wit
ule-of-mixtures (ROM) (Eq. (2)), Hirsch (Eq. (3)), Maxwell (Eq. (4))nd Cox-Krenchel (Eq. (5)) models are commonly used to predictoung’s modulus (Annie et al., 2008; Choudhury, 2008; Li et al.,000, 2008; Rosenthal, 1992; Thomason and Vlug, 1996).
n these equations EC, EF and EM are the Young’s moduli for theomposites, fibers and matrix, respectively; VM and VF are theolume fractions of matrix and fibers, respectively; k and x areonstants dependent on the orientation, aspect ratio and capabilityf the fibers as reinforcing agent (Choudhury, 2008; Annie et al.,008; Li et al., 2000); �o and �L are the fiber orientation factornd fiber length efficiency factor (Rosenthal, 1992; Thomason andlug, 1996). VF was estimated considering the density of pristineurauá fiber equal to 1.1 g cm3 (Mohanty et al., 2002): 0.12 and 0.21or 10 wt% and 20 wt% of curauá fibers, respectively. Using thesearameters, the Serial model gave unsatisfactory results for com-osites independent of the plasticizer and fiber contents, probablyecause it is more appropriate for composites with short and ran-omly oriented fibers in the matrix (Choudhury, 2008). The Young’sodulus predicted by Maxwell model is close to the experimental
alue for the composite with 20 wt% of DOP and 10 wt% of pristinebers (Fig. 4a). However, this model fails with increasing fiber con-ent (Fig. 4a) due to incomplete wetting of the fibers by the matrix,s pointed out by Sreekumar et al. (2008) and Hu and Lim (2007),nd also with increasing plasticizer content (Fig. 4b). The Hirschodel, adopting x = 0.12, was found to be the best option to predict
he Young’s modulus for composites with 20 wt% of DOP and 10 wt%f pristine fibers and with 30 wt% of DOP. The ROM model predic-ions fit very well with the experimental values of Young’s modulusor the biocomposites with 10 wt% of fibers, assuming k = 0.2; 0.24
h: (a) 20 wt% and (b) 30 wt% of DOP. (—) Serial model, (- - -) ROM model with k = 0.2, 0.2, ( ) Cox Krenchel model with �o = 0.375 and (�) experimental.
and 0.35 for pristine, with acetone extracted and with NaOH solu-tion treated curauá fibers, respectively. According to Choudhuryk = 0.2 is used for composites with aligned fibers (Choudhury, 2008).As reported above, the curauá fibers are aligned with the injectionflow direction (micrographs shown in Fig. 2). The deviation of k val-ues from 0.2 (0.24 and 0.35) for treated fibers can be due to the useof the same values of VF in spite of the changes in the compositionand crystalline degree of the fibers as consequence of the treat-ments. The mechanical properties of the biocomposites (Table 1)and also the thermal properties (Tables 2–4) indicate changes inthe interaction density between matrix and fibers surface. Accord-ing to Choudhury and Sreekumar et al. high k values is associatedto high interfacial adhesion (Choudhury, 2008; Sreekumar et al.,2008), what reinforces the hypothesis that the chemical treatmentsof the fibers improve the interactions in the fiber–matrix interface.
The Cox’s model (Cox, 1952) is commonly used to predict themodulus (and strength) of random-in-plane short fiber composites(Lee et al., 2012; Asgari and Masoomi, 2013). Krenchel modifiedthis model by introducing the fiber orientation factor, �o, into the“rule-of-mixtures”
�o =∑
n
an cos4�n (6)
where an is the fraction of fibers with orientation angle �n withrespect to the reference axis. If the fibers are random in-plane ori-entated �o = 0.375 (Rosenthal, 1992). If not, �o should be calculatedusing Eq. (6) and, for this, others parameters are necessary to esti-mate an values, such as modulus in parallel (EC), 45◦ diagonal (E45)and transverse direction (ET), in relation to the injection flow direc-tion (Rosenthal, 1992).
The fiber length efficiency factor �L is giving by the equation:
� =[
1 − tanh(ˇL/2)]
(7)
ˇ2 = 2D
[2GM
EF ln(√
r/R)
](8)
368 M.C. Gutiérrez et al. / Industrial Crops and Products 52 (2014) 363– 372
Table 2Thermal expansion coefficient (˛) , � relaxation (T�) and glass transition temperature (Tg) of cellulose acetate and biocomposites.
(�m mm−1 K−1) Tg (◦C) Tˇ (◦C)
T < Tg From (E′′) From (tan ı) From (E′′) From (tan ı)
here L, D, r, R , Gm and � are fiber length, fiber diameter, fiberadius, mean center-to-center distance between fibers, shear mod-lus and Poisson ratio (Thomason and Vlug, 1996; Lee et al., 2012),espectively. The r/R factor is related to the fiber volume fractionThomason and Vlug, 1996) by
n(√
r/R) = ln(√
�/Xivf
)(10)
here Xi is a geometrical packing arrangement parameter of thebers.
The �L values for the biocomposites of cellulose acetate calcu-ated using Eq. (7) vary from 0.83 to 0.90 (Table 1). The lower value0.83) was found for the biocomposite containing fibers treatedith NaOH solution (CA30D10FM). The fiber length (L) and diam-
ter (D) values used in Eqs. (7) and (8) are assumed to be the sameor pristine and treated fibers (5 mm and 0.01 mm, respectively).his should lead to errors in the prediction of nL, since lignin andemicellulose are extracted with the treatment of the fibers withaOH solution, changing the composition and the density of thebers.
Young’s moduli of the biocomposites of cellulose acetate were
alculated using Eq. (5). The �o parameter could not be estimatesing Eq. (6) because parallel and 45◦ diagonal modulus of the bio-omposites were not measured. Therefore, Young’s modulus wasalculated using �o = 0.375 (for fiber random orientation) and 0.2
able 4hermal conductivity – � (W m−1 K−1) – at different temperatures for plasticized cellulos
(for comparison with ROM model predictions) with the followingvalues: 0.12 or 0.21 for Vf for composites with 20 wt% and 30 wt%of fibers, respectively; �L, EC and EM are given in Table 1; Xi = 4.0(assuming square packing of fibers (Thomason and Vlug, 1996)) andPoisson ratio of 0.31 (Guo and Barbari, 2010). The Young’s modulusvalues predicted by the Cox-Krenchel model are plotted in Fig. 4.This model overestimates the Young’s modulus for the biocompos-ites of cellulose acetate and curauá fibers for �o = 0.375. However,for �o = 0.2 the predicted values fit very well the experimental dataand better than ROM equation. Comparing Eqs. (2) (ROM) and (5)(Cox-Krenchel) it can be concluded that k = �L �o. Choudhury (2008)adopted k = 0.2 and �o = 0.2 to estimate the Young’s modulus ofhigh-density polyethylene composites reinforced with sisal fibersand in these conditions Eqs. (2) and (5) are equivalent for �L = 1. TheROM predictions for the cellulose acetate biocomposites preparedin this work assumed also �L = 1. The better fit of the Cox-Krenchelprediction in comparison with ROM predictions is possibly due tothe �L, which values vary with the composite composition and fibertreatment (Table 1).
The increase of the tensile strength of the plasticized celluloseacetate by addition of curauá fibers (Table 1) is in agreement withthe reinforcement characteristics of the filler. Moreover, the chem-ical treatment of the fibers has a positive effect on this property, forinstance, the tensile strength of biocomposites containing 30 wt%of DOP and 10 wt% of fibers is around 60% higher for chemically
treated fibers. All these results also indicate an improvement inthe transference of the mechanical energy between polymericmatrix and the fibers. This improvement can possibly be due to theincrease of the interactions density involving hydrophilic groups
n the surface of the fibers (hydroxyls and carbonyls) and in theatrix. This hypothesis could also justify a drop of the strain of
reak of the biocomposites containing treated fibers in comparisonith biocomposites with pristine fibers.
The impact strength of the biocomposites (Table 1) is deter-ined by both variables: plasticizer content and fiber content. The
iocomposite containing 10 wt% of pristine fibers and 30 wt% ofOP (CA30D-10F) presents higher impact strength in comparisonith the biocomposite containing 20 wt% of DOP (CA20D-10F). The
rgument to explain this fact is the same used for plasticized cel-ulose acetate: the increase of free volume and the consequentecrease of the relaxation time. Similar results have been reportedor composites based on plasticized cellulose acetate and clay (Parkt al., 2004; De Lima et al., 2012). The introduction of the fibersecreases the impact strength, independent of the content of plas-icizer. This effect is more pronounced for chemically treated fibers,lthough no significant difference had been observed for the impactesistance of the biocomposites prepared with fibers treated withcetone or NaOH solution.
The impact strength for the CA30D prepared on a labora-ory scale and on a large scale is similar [(327 ± 7) J m−1 and297 ± 5) J m−1]. However, the impact strength for the CA30D-0F prepared on a laboratory scale and on a large scale is quite
different: (73 ± 2) J m−1 and (161 ± 5) J m−1, respectively. Thisdifference is due to the larger extent of fibrillation promoted bythe processing conditions on the laboratory scale, as previouslydiscussed (Gutiérrez et al., 2012a,b).
The storage modulus of the composites in the temperaturerange from 100 to 200 ◦C and from 70 to 200 ◦C for formula-tions containing 20 wt% and 30 wt% of DOP, respectively, is high incomparison with the E′ values for the corresponding plasticized cel-lulose acetate, due to the reinforcing effect of the fibers (Fig. 5a andb). Wong et al. reported similar effects for composites of poly(lacticacid) with flax fibers (Wong et al., 2003). Fig. 5c and d shows lossmodulus (E′′) and Fig. 5e and f shows the tan ı as a function of tem-perature for cellulose acetate and biocomposites plasticized with20 and 30 wt% of DOP. The glass transition temperature (Tg), aswell as the transition temperature (Tˇ), (Table 2) are taken asthe temperatures corresponding to the maximum of the peaks inthe loss modulus curves and tan ı curves. For formulations con-taining 20 wt% of DOP, Tˇ was determined only from tan ı curves,because the peak in the E′′ curves was not well defined. Cellulose
acetate with 20 wt% and with 30 wt% of DOP show Tg at 115 ◦Cand 78 ◦C (data obtained from E′′ curves), respectively. A decreaseof Tˇ with increasing amount of DOP was also observed. Accord-ing to Videki et al. transition is associated with relaxations of
370 M.C. Gutiérrez et al. / Industrial Crops and Products 52 (2014) 363– 372
he glucose monosaccharide unit rings (Videki et al., 2007). The Tg
f cellulose acetate with 20 wt% of DOP shifts 6 ◦C to higher tem-eratures with the addition of 10 wt% or 20 wt% of pristine fibers.omparable increase of Tg of the polymeric matrixes of the compos-
tes, induced by fillers, has been reported in the literature; e.g. foromposites of poly(l-lactide) and carbon fiber (Han et al., 2012) andomposites of poly(3-hydroxybutyrate-co-3-hydroxyvalerate) andanofibrillated cellulose (Srithep et al., 2013). According to theseuthors, the increase of Tg is due to limitation of chain mobilityithin the polymer matrix induced by the filler. On the other hand,
or cellulose acetate biocomposites containing 30 wt% of DOP, theddition of pristine fibers (10 or 20 wt%) practically does not shiftg. However, the replacement of 10 wt% of pristine fibers for the
ame amount of treated fibers causes an increase of 22 ◦C in Tg.ccording to Gerald et al. (1990), Tg and storage modulus of com-osites vary with the “quality” of the interface, being higher forhose with stronger interfacial bonding. Therefore, the increase of
30 wt% DOP
cient (˛) and thermal conductivity (�) of the biocomposites.
Tg of the composite with the fiber treatment is possibly due to theimprovement of the matrix–interface interactions.
3.3. Insulating thermal properties of the biocomposites
Heat capacities (Cp) and thermal conductivities (�) of cellu-lose acetate and its biocomposites at 20, 40, 60, 80 and 100 ◦C arepresented in Tables 3 and 4, respectively. These parameters areimportant to define the processing conditions as well as the poten-tial of applications for these biocomposites (Lopes and Felisberti,2004).
In general, the heat capacity of the composites is higher thanthat of plasticized cellulose acetate. A possible explanation for
this is based on the high heat capacity of the lignocellulosicfibers (1.7 J g−1 K−1) (Wu et al., 2000). Above Tg, the heat capac-ity increases due to the increase of the degrees of freedom. BelowTg, these are restricted to vibrational and rotational modes (Cassu
rops a
aNrpefisdt
bis(Cert
ita1vnawcbcIbt2atiNtcitctt
oerrod0Ttbp0
oeicfitb
M.C. Gutiérrez et al. / Industrial C
nd Felisberti, 2005). Biocomposites containing fibers treated withaOH solution show lower heat capacity in relation to the cor-
esponding plasticized cellulose acetate and biocomposites withristine and with acetone extracted fibers. One hypothesis toxplain this could be the anchoring of the polymer chains on thebers and the lower flexibility of the fibers treated with NaOHolution (higher crystalline degree), restricting chain mobility,ecreasing the heat capacity and increasing the thermal conduc-ivity.
The processing conditions affect the thermal properties of theiocomposites in a complex way. For example, the heat capac-
ty at 20 ◦C of the CA30D-10FA prepared by laboratory and largecales are (1.721 ± 0.001) J g−1 K−1 (Gutiérrez et al., 2012a,b) and1.450 ± 0.002) J g−1 K−1, respectively and for the biocompositesA30D-10FM the values are (1.221 ± 0.004) J g−1 K−1 (Gutiérrezt al., 2012a,b) and (0.974 ± 0.001) J g−1 K−1, respectively. Thisesult shows that not only the processing conditions influences thehermal properties but also the nature of the chemical treatment.
The thermal conductivity of plasticized cellulose acetatencreases with temperature below and near Tg, because the heatransfer occurs through the vibrational modes of the networkccording to the phonons-fractons model (Dashora and Gupta,996). Above Tg a drastic drop is observed, due to the higher freeolume and consequently higher chain flexibilities, which has aegative impact on heat transfer. The argument of chain flexibilitynd free volume can also be used to explain why cellulose acetateith 20 wt% of DOP presents higher thermal conductivity than
ellulose acetate with 30 wt% of plasticizer. Similar behavior haseen reported by Dashora and Gupta (1996) for amorphous non-onductive polymers, such as poly(vinyl chloride) and polystyrene.n multiphase materials, the heat transfer depends on the num-er of phases and consequently on the interfacial area, becausehe interfaces change the rate of transfer (Lopes and Felisberti,004). For the composites, the presence of fibers contributes neg-tively to heat transfer. Comparing composites, those containingreated fibers present higher thermal conductivity than compos-tes with pristine fibers, with the contribution of fibers treated withaOH solution being more significant. Thus, the characteristics of
he fibers play an important role on the thermal properties of theomposites, either due to their own properties or by propertiesnherent to the fiber–matrix interface. Moreover, the increase ofhe concentration of pristine fibers causes a decrease of the thermalonductivity, reinforcing the hypothesis of the negative contribu-ion of the “number” of phases/interphases on the heat transferhrough the material.
As observed for the heat capacity, the thermal conductivityf the biocomposites prepared on a laboratory scale (Gutiérrezt al., 2012a,b) and on a large scale(this work) varies similarly withespect to the chemical treatment of the fibers. Takagi et al. (2007)eported a thermal conductivity of 0.3 W m−1 K−1 for compositesf poly(lactic acid) and flax fibers at 25 ◦C. The biocompositesescribed in this article present thermal conductivity between.08 and 0.15 W m−1 K−1 (Table 1), depending on the composition.he low thermal conductivities of these biocomposites indicatehat they can be used as thermal insulators. Traditionally, non-iodegradable thermal insulation materials like polyurethane orolystyrene foams have thermal conductivities between 0.02 and.06 W m−1 K−1 (Wu et al., 2000).
Thermal expansion coefficients (˛) calculated from the databtained by dilatometry are shown in Table 2. The biocompositesxhibit higher thermal dimensional stability than the correspond-ng plasticized cellulose acetate, since the thermal expansion
oefficient of the composites decreases with increasing amount ofbers. Moreover, the chemical treatment of the fibers contributeso the improvement of the thermal dimensional stability of theiocomposites.
nd Products 52 (2014) 363– 372 371
Finally, Fig. 6 shows the interdependence of the Young’s modu-lus, impact strength, thermal expansion coefficient (˛) and thermalconductivity (�), and summarizes the important characteristics ofthe biocomposites of cellulose acetate, plasticizer and curauá fibers.Although the increase of Young’s modulus and the decrease of theimpact strength and of the thermal expansion coefficient can beeasily explained based on the cellulose acetate chains flexibility andits dependence on the plasticizer and fiber contents, the same argu-ments are not enough to explain the thermal conductivity behaviorof the biocomposites. The thermal conductivity depends also on thepolymer chain flexibility, however, other factors seems to be moreimportant in determining the thermal conductivity of the biocom-posites; for instance the number of phases and the tortuosity ofthe polymeric matrix containing dispersed fibers. This hypothesisis reinforced by the consistent decrease of the thermal conductiv-ity with increasing fiber contents for both biocomposites with 20and 30 wt% DOP. The chemical treatment of the fibers has also animpact on the thermal conductivity and in this case we believe thatthe increase of the thermal conductivity observed for CA20D-10FAand CA20D-10FM in comparison with CA20F-10F should be relatedto the properties of the fibers and also to the improvement of thefiber–matrix adhesion.
4. Conclusions
Biocomposites based on cellulose acetate and short curauá fiberscan be prepared on a large scale by extrusion and their mechani-cal and thermal properties are useful for many applications. Themechanical properties can be tailored by controlling the composi-tion of the biocomposites as well as by chemical treatment of thefibers. The chemical treatments of the fibers result in changes inthe fiber–polymer interface and consequently improve the stiffnessand increase the thermal conductivity and the thermal dimensionalstability of the biocomposites. The impact strength of the com-posites prepared on a large scale is superior in comparison withcomposites prepared in laboratory scale. The thermal properties ofthe composites are influenced by the processing conditions. Thebiocomposites prepared in this work show acceptable dimensionalstability and low thermal conductivity, similar to thermal insulat-ing materials like polyurethane foams or expanded polystyrene, butwith mechanical properties of dense materials. The large extentof fibrillation does not necessarily have a positive effect on thethermal and mechanical properties of the biocomposites.
Acknowledgments
M. Chavez G. acknowledges a fellowship from CAPES – programPEC-PG. The authors acknowledge Prof. C.H. Collins for manuscriptrevision, FAPESP (2010/02098-0 and 2010/17804-7) for financialsupport and EMBRAPA-PA for donation of the fibers.
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