Mechanical Behavior and Fracture Toughness of Poly(L-lactic acid)Natural Fiber Composites Modified...

Mechanical Behavior and Fracture Toughness of

Poly(L-lactic acid)-Natural Fiber Composites Modified

with Hyperbranched Polymers

Susan Wong, Robert A. Shanks,* Alma Hodzica

CRC for Polymers, Applied Science, RMIT University, GPO Box 2476V, Melbourne, Victoria, Australia, 3001Fax: þ61 3 9925 2255; E-mail: [email protected]

Received: December 3, 2003; Revised: February 19, 2004; Accepted: February 20, 2004; DOI: 10.1002/mame.200300366

Keywords: biofibers; biomaterials; composites; mechanical properties; toughness

Introduction

Green composites that consist of biodegradable compo-

nents have received much attention in the past decade. The

use of natural fibers (such as flax, hemp, and sisal) in combi-

nation with biodegradable polymers (synthetic and natural)

and/or petrochemical polymers has been shown to yield

properties suitable for low stress applications such as auto-

motive door panels and rear-parcel shelf panels. The use

of natural fibers in place of synthetics such as glass or

carbon fibers offers the advantage of high specific stiff-

ness and strength, desirable fiber aspect ratio, low density,

biodegradability, lower cost per unit volume, and sound

absorption.[1,2]

Of the many biodegradable thermoplastics available,

poly(L-lactic acid) (PLLA) has been themost popular due to

Summary: The use of hyperbranched polymers (HBP) withhydroxy functionality as modifiers for poly(L-lactic acid)(PLLA)-flax fiber composites is presented. HBP concen-trations were varied from 0 to 50% v/v and the static anddynamic tensile properties were investigated along withinterlaminar fracture toughness. Upon addition of HBP, thetensile modulus and dynamic storage modulus (E0) both di-minished, although a greater decline was noticed in the staticmodulus. The elongation of the composites with HBP show-ed a pronounced increase as large as 314% at 50% v/v HBP.The loss factor (tan d) indicated a lowering of the glass tran-sition temperature (Tg) due to a change in crystal morphologyfrom large, mixed perfection spherulites to finer, smallerspherulites. The change in Tg could have also resulted fromsome of the HBP being miscible in the amorphous phase,which caused a plasticizing effect of the PLLA. The inter-laminar fracture toughness measured as the critical strainenergy release rate (GIC) was significantly influenced byHBP. At 10% v/v HBP, GIC was at least double that of theunmodified composite and a rise as great as 250% wasachieved with 50% v/v. The main factor contributing to highfracture toughness in this study was better wetting of thefibers by the matrix when the HBP was present. With im-proved ductility of the matrix, it caused ductile tearing alongthe fiber-matrix interface during crack propagation.

ESEM photograph of propagation region of the interlaminarfracture toughness specimens with 30% v/v of HBP.

Macromol. Mater. Eng. 2004, 289, 447–456 DOI: 10.1002/mame.200300366 � 2004 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

Full Paper 447

a Current address: Mechanical Engineering, School of Engineer-ing, James Cook University Townsville, Queensland, Australia,4811.

its high mechanical strength, which can be potentially used

for structural materials.[3] PLLA is an aliphatic polyester

derived from corn and sugar beets[2] and degrades to non-

toxic compounds in landfill. It can be synthesized from

direct condensation of lactic acid and by ring-opening poly-

merization of the cyclic lactide dimer.[2] Of the two isomers

available (D- and L-), the polymer of the L-enantiomorph

is a hard, transparent, and highly crystalline polymer.

Relatively large spherulites are achievable as a result of slow

crystallization and as a consequence, brittleness can occur,

whereby toughness is catastrophically reduced. Toughness

in composites is important to ensure that during use del-

amination and crack propagation within the composite is

minimized.

To overcome the brittleness of PLLA, a large range of

plasticizers has been usedwith some success. These include

citrate esters,[4] 1,2-propylene glycol, glycerol,[5] poly-

(ethylene glycol),[6] glucose monoesters, and fatty acids.[7]

These small molecules can cause significant changes to the

thermal and mechanical properties of PLLA. But due to the

boiling temperatures of these plasticizers being similar to

the melting temperature of PLLA, their concentrations

could vary due to some evaporation during processing.With

long-term use of the composites, the plasticizers could also

leach, which would cause embrittlement. An alternative

to plasticizers is the use of larger polymeric modifiers to

perform the same functions. It is common practice to blend

different polymers to alter their properties. Hyperbranched

polymers (HBP) have attracted attention recently as tough-

ening modifiers for thermosets (epoxy resins) in place of

commercial tougheners such as rubbers, thermoplastics,

or glass particles.[8]

HBP are highly branched macromolecules described as

having a globular structure with a multiplicity of reactive

end groups. They are synthesized by a one-step polycon-

densation reaction, producingmolecules with a high degree

of irregular branching, a broad molecular weight distribu-

tion, and significantly lower viscosity compared with linear

polymers of the samemolecular weight.[9] They are flexible

polymers with low glass transition temperatures (Tg often

below room temperature) and lowmelting temperatures (Tm)

usually below 60 8C. Due to the high degree of branching,

limited crystallization or interchain entangling is possible.

Consequently, this gives rise to poor mechanical properties

but good solubility and reduced melt viscosity.[10]

WhenHBPare utilized in epoxy resins, their outer shell is

functionalized with epoxy groups to ensure that the necess-

ary mechanical interactions are present for good load

transfer from the matrix to the modifier. With only 5% of

HBP, the critical strain energy release rate (GIC) of the resin

is increased by a factor of 6.[11] The toughening effect was

postulated to be induced by a finely-dispersed particulate

structure. Despite the toughening effect, Young’s modulus,

the glass transition temperature (Tg), and the thermo-

mechanical properties were unaffected.

In this study, hyperbranched polyesters were blended

with thermoplastic PLLA in flax fiber composites. At

present to our knowledge, there is limited literature about

hyperbranched polymers behavior in thermoplastics and

in composites. We hypothesize that through blending

with HBP the brittleness of PLLA will be reduced by

altering the morphology of the matrix, which should result

in improved toughness in the prepared composites. The

toughness was evaluated by mode I critical strain energy

release rate (GIC). The tensile static and dynamic mechan-

ical properties were also studied in order to observe if

greater favorable effects on the mechanical properties

would accompany toughness.

Experimental Part

Materials

Hyperbranched polyester of third generation (32 hydroxylfunctionality) (H30) was supplied by Perstop, Sweden. Thehyperbranched polyesters are prepared from condensation ofethoxylated pentaerythritol as a tetrafunctional core and 2,2-dimethylol propionic acid (bis-MPA) as the repeating unit.[9,12]

A commercial sample of poly(L-lactic acid) (PLLA) (ResomerL206) obtained from Boehringer Ingelheim (Ingelheim,Germany) was used for this study. Its physical properties aregiven in Table 1. Flax fibers (Durafiber Grade One of 95%purity) were obtained from Durafiber Inc. (Cargill) and weredried in a vacuum oven at 100 8C for 2 h before preparation ofcomposites.

Preparation of Composites

Blends of hyperbranched polyesters and PLLAwere solutionmixed prior to incorporating with flax fibers to form compo-sites. The hyperbranched polyesters were initially dissolved in1,4-dioxane separately to obtain a concentration of 1% w/v.The same concentration of PLLA solution was preparedby dissolving in chloroform under reflux. The appropriatevolumes were combined to give blends of 10:90, 30:70, and50:50 v/vHBP to PLLA.Dried flax fiberswere incorporated bymixing while the blends were in solution and were cast ontoglass plates and the solvents were evaporated. The fiber:poly-mer blend ratio was kept constant at 1:1 v/v. Residual solventwas eliminated by drying in a vacuum oven at 50 8C for 3 h.

Table 1. Physical properties of materials.

Material Density Mw Tg Tm Hf

g/cm3 8C 8C J/g

PLLA 1.206a) 110 000a) 56a) 176a) 58.0a)

HBP 1.295c) 3 500b) 10c) 57c) 9.1c)

a) Data obtained from: Y. Choi, S.Y. Kim,M.-H.Moon, S. H. Kim,K.-S. Lee, Y. Byun. Biomaterials 2001, 22, 995.

b) Data obtained from manufacturer (Perstorp).c) Data obtained from ref.[9]

448 S. Wong, R. A. Shanks, A. Hodzic

Macromol. Mater. Eng. 2004, 289, 447–456 www.mme-journal.de � 2004 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim

After elimination of residual solvents, the polymer-fibersheets were cut to smaller pieces and heat-pressed betweenTeflon sheets at 175 8C for 3 min under 1 t pressure to obtaincomposite sheets for static and dynamic tensile mechanicaltesting. The selected molding temperature and low residencetime used was to minimize degradation of PLLA by chainscission. The composite sheets were 0.7mm thick. All sampleswere then thermally treated to obtain the same thermal historyby annealing at 50 8C for 24 h.

For the specimens used in the interlaminar fracture tough-ness tests, dried fibers that were randomly oriented and conti-nuous were placed in an aluminium mold. The solution blendsat the same concentrations as described in the previous twoparagraphs were poured into the mold and left to dry. Thepolymer-fiber sheets were removed from the mold and heatpressed at 175 8C for 3 min under 1 t pressure. A thin piece ofaluminium foil of thickness about 20 mm was inserted in-between two polymer-fiber sheets to introduce a pre-crack of50 mm and heat-pressed as previously to form a plaque ofthickness of 2.8 mm. Annealing was also carried out at 50 8Cfor 24 h.

Mechanical Characterization

The stress-strain behavior of the compositeswasmeasured by adynamic mechanical analyzer (Perkin-Elmer DMA 7e) undertension in static mode. Samples of 5 � 12 � 0.7 mm3 werestretched from50 to 1 000 kPa at a rate of 50 kPa/min at a gaugelength of 7 mm. The measurements were performed at roomtemperature and at least eight replicates were averaged. Thetensile modulus and elongation were determined from thestress-strain curves.

Samples with the above dimensions were subjected to dyna-mic testing in tensile mode. The instrument was set to performwith a constant dynamic strain of 0.1% with a static:dynamicforce ratio of 110%. These parameters were selected afterdetermining the linear viscoelastic range and were within thelinear viscoelastic limits. The sampleswere subjected to a fixedfrequency of 10Hz and the temperature range investigatedwasfrom �50 8C to 100 8C at a heating rate of 2 8C/min undernitrogen atmosphere. The results presented are averaged fromat least three replicates. Calibrations for the height, force,furnace temperature, and eigendeformation were performedbefore data acquisition using the three point bending measur-ing system. The height calibration standardized the displace-ment transducer to measure the position and the amplitudeaccurately. The force calibration sets the static and dynamicforces of the displacement transducer. The furnace temperaturecalibration involved matching the thermocouple readings toreferencematerials of indium and zinc then to the programmedfurnace temperature. The eigendeformation calibration adjustthevery smallmovements of the analyzerwhen large forces areapplied so that increasing force with no sample yielded zerodisplacement.

The interlaminar fracture toughness of the prepared compo-sites was evaluated in accordance with ASTM D 5528-01 bythe double cantilever beam (DCB) tests. Samples were of con-tinuous, random orientated fibers with dimensions of 140 � 20 �2.8 mm3 and a pre-crack (a0) of 50 mm. Due to the samplesbeing inflexible, irreversible deformation of the specimen arms

occurred before the desired total crack length was achieved(100 mm). To eliminate this problem, two thin sheets of‘‘Perspex’’ were bonded to the specimens with ‘‘Supaglue’’ toaid with specimen flexing during the test. Piano hinges wereattached to the specimens and the applied load was introducedby an Instron (model 4065) universal test instrument at thestrain rate of 3 mm/min. The critical strain energy rate (GIC)was calculated by the modified beam theory from the follow-ing equation:

GIC ¼ 3Pd2bðaþ DÞ ð1Þ

where P is the applied load, d is the load point displacement, bis the specimenwidth, a is the delamination length, andD is thepoint of x intersection when y¼ 0 from the cube root ofcompliance (C1/3) and delamination length (a) plot. The com-pliance, C, is the ratio of load point displacement to load. GIC

values presented are those obtained at the non-linear point ofthe load (P) vs. displacement (d) plots.

Visual Examination of Composites

The crystal morphology of the composites was examined bypolarized optical microscopy (OM) and the fracture surfacesfrom the DCB tests were observed by environmental scanningelectron microscopy (ESEM). OM was performed using aNikon Labophot 2 optical microscope and the photographswere captured on a Nikon digital camera. Special samples forOM consisting of a single fiber in the blend of interest wereprepared so that the interaction between thematrix and the fibercould be seen more clearly. Films of the blends were cut andplaced on a glass slide with a single fiber and melted on aMettler FP82 HT hot stage at 175 8C for 3 min. The glass slidewas then removed from the hot stage and left to crystallize.ESEM micrographs of the fractured surfaces of the samplesfrom the DCB tests were conducted using conventionalsecondary electrons with an accelerating voltage of 20 kV at0.5 Torr.

Results and Discussion

Static Tensile Behavior

The static tensile behavior of the composites with HBP is

shown in Figure 1 and Table 2. It is apparent that as the

proportion of HBP increases, the tensile modulus decreases

as expected, which shows that the HBP has a softening

effect on the composites. The tensile modulus was reduced

by 4%with 10% v/v of HBP compared with the unmodified

PLLA composite and a reduction as great as 52% was

achieved with 50% v/v of HBP. Consequently, the elonga-

tion at maximum stress was increased with increasing pro-

portion of HBP. With 10% v/v of HBP, the elongation was

increased by 21% and a pronounced increase of 214% is

resulted with 50% v/v of HBP. The dramatic effects on the

tensile properties were a result of the dissimilarity in crys-

tallinity (as indicated by the heat of fusion, DHf (Table 1))

Mechanical Behavior and Fracture Toughness of Poly(L-lactic acid)-Natural Fiber Composites Modified . . . 449


between HBP and PLLA. The low crystallinity of HBP was

brought about by the linear segments in the globular struc-

ture that fold back on themselves and crystallize. The

globular structure also does not allow for chain extension

and orientation, which prevents strain hardening from

occurring.[9]

Dynamic Mechanical Behavior

The dynamic mechanical behavior in terms of storage

modulus (E0) and loss tangent (tan d) is presented in Figure 2and 3. The storage modulus or elastic modulus (E0) indi-cates the elastic portion of the complex modulus (E*),

which is the vector sum of the elastic and loss moduli. The

E0 is a measure of the stiffness of the viscoelastic material.

The loss modulus (E00) (not presented here) is related to thework dissipatedwithin thematerial during one load cycle. It

is a measurement of the viscous component or oscillation

energy dissipated that is unrecoverable during a load cycle.

The loss tangent or damping factor (tan d) is the ratio of theE00 and E0, which is a measure of the energy loss in relation

to the recoverable energy.

Upon addition of HBP, the E0 of the composites all

showed diminished values over the temperature range in-

vestigated, revealing that the stiffness of the composites

decreased accordingly. This coincided with the changes

in the static tensile moduli measured. Before reaching the

glass transition region (temperatures below �40 8C), thedecrease in the E0 was only about 0.3 to 0.4 GPa relative tothe unmodified PLLA composite. In dynamic mode, the

apparent reduction in E0 was not as significant compared

with the static tensile modulus.

Figure 1. Stress-strain curves of composites containing 0 (&),10 (*), 30 (~), and 50 (!)% v/v HBP.

Table 2. Static tensile properties of composites.

[H30] Tensile modulus Elongation

% v/v MPa %

0 11.92 (0.01)a) 0.070 (0.002)10 11.45 (0.04) 0.085 (0.009)30 9.21 (0.02) 0.106 (0.009)50 5.73 (0.01) 0.22 (0.03)

a) Standard deviation given in brackets.

Figure 2. Storage modulus with temperature of compositescontaining 0 (&), 10 (*), 30 (~), and 50 (})% v/v HBP.

Figure 3. Loss tangent of composites containing 0 (&), 10 (*),30 (~), and 50 (})% v/v HBP.



One distinct feature of the E0 curves of the composites

with HBP as opposed to the unmodified PLLA composite is

an inflection at temperatures above 60 8C. The inflection

was attributed to recrystallization upon heating. This was

also observed in the melting scans of the composites from

differential scanning calorimetry (Figure 4) where crystal-

lization proceeded over the Tg in the presence of HBP.

Further discussion of the thermal properties is presented at a

later section. The inflection is most pronounced at 10% v/v

of HBP and the crystallization temperature (Tc) shifted to

lower temperatures as the proportion of HBP increased.

The occurrence of the inflection was a consequence of

the reduction of the glass transition temperature (Tg) upon

blending with HBP, which enabled PLLA molecules to

obtain the mobility to start crystallizing. The Tg’s of the

HBP composites was reduced by at least 5 8C compared

with the unmodified PLLA composite as shown in Figure 3

as a peak in the tan d curves. As only one peak is observed ineach curve, it implied that some dissolution of HBP into the

PLLA matrix has occurred. The decrease in the tan d peak

temperatures at increasing proportions of HBP and the

small reduction in the crystallization temperatures (Tc),

presented in Table 4 (see below), also suggested this

behavior. However, the decrease in Tg is not as high as one

would expect for a totally miscible blend therefore the

blends can be inferred as partially miscible. The misci-

bility will be further clarified by observation of the crystal

morphology with optical microscopy.

The peak intensity of the tan d curves of materials can

generally give some indication of the extent of the bonding

at the fiber-matrix interface. It is well recognized that com-

posites with diminishing peak intensity at a given tempe-

rature indicate a strong bonded fiber-matrix interface as

a greater number of polymer chains is restricted in their

movement.[13] However, a strong interface usually results

in low toughness as it does not allow the stress to be relieved

via interfacial debonding.[14] It is evident that a careful

balance of the interfacial properties should be considered in

order for a composite to ideally show relatively good interfa-

cial bonding so that high strength is achieved, but at the same

time, the toughness is not detrimentally compromised.

From the loss tangent curves (Figure 3), a remarkable

increase in the peak intensities was achieved with increas-

ing proportions of the HBP. The peak intensity of the com-

positeswith 10, 30, and 50%v/v ofHBP showed an increase

of about 4, 5, and 6 times greater than the unmodified com-

posite, respectively. As previously discussed, the increase

in the magnitude of the tan d peaks implies an increase in

toughness, therefore by blending with HBP the results in-

dicate that a favorable effect on the toughnesswas achieved.

This will be further realized by the interlaminar fracture

toughness tests.

Interlaminar Fracture Toughness

The interlaminar fracture toughness expressed in terms of

the critical strain energy release rate (GIC) is an important

property that not only gives a measure of toughness (resis-

tance to crack propagation) but is also very sensitive to

the matrix crystal morphology as well as the fiber-matrix

interface.[15] In mode I crack propagation three general

types of behavior have been observed as shown in Figure 5.

Crack propagation can occur continuously (Figure 5(a))

where the force gradually declines as the displacement in-

creases. In the second instance (Figure 5(b)), ‘‘stick-slip’’

propagation can occur whereby the force becomes unstable

as displacement increases due to crack initiation followed

by crack arrest. The third behavior shown (Figure 5(c)) is a

hybrid of the first two, where continuous propagation occurs

along with stages of crack initiation and crack arrest.[16]

In the prepared composites, crack propagation generally

occurred in a stable, continuous manner as illustrated in

Figure 6 by the force-displacement (P-d) plots. It is worthnoting that at the end stages of the unmodified PLLA com-

posites, some instability was observed, which could be in-

duced by voids present in one particular replicate rather

than its true behavior.

Comparedwith the unmodifiedPLLAcomposite, all com-

posites with HBP displayed a greater opening displacement

before reaching the maximum load, which also increased

correspondingly with HBP content. With 10% v/v HBP, a

small incremental increase in the maximum load of about

9% was observed compared with the unmodified compo-

site. Greater increases of load were realized for composites

with 30 and 50% v/v of HBP with magnitudes of 30 and

68%, respectively. From these observations, it can be de-

duced that the composites becamemore ductilewith greater

proportion of HBP, which was consistent with the static

Figure 4. DSCmelting scans of PLLAcomposites and its blendswith hyperbranched polyesters.



tensile moduli previously described. Generally, it requires

greater force to cause delamination between ductile mater-

ials as opposed to a brittle one due to the ductile material

being capable of drawing before complete fracture.

The resistance curves (R-curves) of the composites illus-

trated in Figure 7 show the energy strain release rate (GIC)

with specimen crack growth. Initially, the GIC values were

relatively low and then rapidly rose to a maximum value

Figure 5. Schematics of continuous (a), stick-slip (b), and hybrid (c) crack propagationbehavior.

Figure 6. Load-displacement curves (P-d) of composites con-taining 0 (&), 10 (~), 30 (*), and 50 (}) % v/v HBP.

Figure 7. Resistance curves (R curves) of composites containing0 (&), 10 (~), 30 (*), and 50 (}) % v/v HBP.



where crack growth through the specimen occurred. After

the maximum, theGIC values reached a stable plateau where

the crack growth propagated in a constant manner.

From the R-curves, the GIC values of the composites

increased with increasing proportions of HBP. The GIC

values of the composite with 50% v/v was the most pro-

nounced of all composites which was not surprising as the

matrix was the most ductile. By blending the PLLA with

HBP, the composites showed an improvement in toughness

by at least double in contrast to the unmodified composite.

From the GIC values measured at the non-linear point

(Table 3), the increase in toughness at 10% v/v HBP was

about 1.9 times greater than the unmodified composite.

With 30% v/v HBP, the toughness was favorably increased

by about 2.1 times. At 50% v/v of HBP, the increase was

3 times greater. At this stage, HBP proved to be effective

toughness modifiers as a dramatic influence on the tough-

ness was achieved, but the mechanical properties of PLLA

were not detrimentally compromised when used in small

proportions (up to 30% v/v), as demonstrated in the static

and dynamic mechanical properties.

Visual Examination of Composites

Observation of Crystal Morphology

The morphology of the various systems using optical

microscopy is presented in Figure 8(a–d). It is evident that

without HBP, the typical crystals of PLLA were large and

with mixed perfections. Due to the large crystals formed,

large cracks were apparent throughout the matrix. This is

typical of a brittle matrix with a high Xc. Transcrystallinity

is also observed along one side of the fiber but voids resulted

from thermal contraction of the matrix from the fiber were

also evident. Brittleness of the matrix was inferred by the

formation of coarse, large spherulites with poor intercon-

nected crystal structures, which led to a high tensilemodulus

and E0, but low fracture toughness for the unmodified com-

posite. As Xc was high in the unmodified composite, a high

Tg (tan d peak) was detected in the dynamic mechanical

analysis and the change in E0 during the glass transition wasnot as dramatic as those when HBP was present.

For the systems with 10% v/v HBP, some transcrystals

were formed along the fiber, but not in the systems with

higher proportions of HBP. This coincided with a reduction

in Xc (Table 4). Good crystal structure was also achieved

for all systems with HBP in contrast to the unmodified sys-

tem due to the HBP yielding smaller and finer PLLA spher-

ulites. Cracks and thermal contraction of the matrix were

also eliminated or reduced in the presence of HBP, which

may account for the favorable effects on the fracture tough-

ness. The presence of the HBP in the systems can be seen as

Table 3. GIC values of composites at non-linear point.

[H30] GIC

% v/v J/m2

0 38.9 (8.4)a)

10 73.2 (7.9)30 83.3 (5.5)50 115.4 (6.9)

a) Standard deviation given in brackets.

Figure 8. OM photographs of crystal morphology of (a) 0, (b) 10, (c) 30, and (d) 50% v/vHBP-PLLA flax systems.



‘‘droplets’’ in the photographs. At 10 and 30% v/v HBP,

good dispersion of HBP within the PLLA matrix was

achieved at these concentrations with minimal agglomera-

tion observed. The morphology of the systems with 10 and

30% v/v HBP are very similar but the appearance of many

dark spots are evident in 30% v/v, which may be attributed

to some HBP inhibiting crystallization of PLLA. This was

confirmed by a decrease in Tc (Table 4). However, at 50% v/

v significant phase separation occurred. This appeared as

dark regions in the photograph near the fiber, as HBP do

not crystallize. The large increase in GIC values of the

composite with 50% v/v relative to other concentrations

may result from the properties of the HBP overshadowing

those of PLLA when present in large quantities. These

observations supported the suspicion that the blends are

partially miscible, as the Tg’s observed by DMA do not

decrease in magnitude coinciding with the relative propor-

tion of HBP.

In a study by Perrin et al.[15] it was concluded that com-

posites exhibit higherGIC values when two general require-

ments are met; the composite has good fiber distribution

and the crystal structure is well-interconnected. Poor fiber

distribution results in fiber-rich and softer matrix-rich re-

gions, which can consequently cause a crack to deviate

along its path, yielding lower fracture toughness. In general

large, coarse sized spherulite distribution with weak inter-

spherulitic regions will lead to reduced transfer of stress

between the matrices, which contributes to low fracture

toughness. It was found that fine spheruliticmicrostructures

resulted in a transpherulitic fracture path that displayed

higher fracture toughness than an interspherulitic fracture

path that resulted from a coarse spherulitic microstructure.

As fine and small spherulites were exhibited when blended

with HBP in the composites, this could provide the favor-

able fracture toughness observed.

Fracture Surface Analysis

The surface morphology of the interlaminar fracture tough-

ness specimens in the initiation region are presented in

Figure 9(a–d) as observed using ESEM. The initiation reg-

ion was in the area where the thin foil film was introduced

Table 4. Thermal properties of PLLA composites andhyperbranched polyesters.

[H30] Tga) Tc Tm DHf Xc

% v/v 8C 8C 8C J/g

0 76 110 176 32.6 0.7010 63 96 177 30.0 0.5830 57 109 175 20.0 0.3050 58 109 179 16.2 0.17

a) Tg taken from the maximum peak temperature of tan d curves.

Figure 9. ESEM photographs of initiation region of the interlaminar fracture toughnessspecimens with (a) 0% v/v, (b) 10% v/v, (c) 30% v/v, and (d) 50% v/v of HBP.



to promote a pre-crack, and this region is located on the

right side of each photograph.

The most obvious feature in these photographs is the

difference in the degree ofwetting of the fibers by thematrix

as seen in the left side of each photograph. The wetting of

the fibers was greatest for those composites with higher

proportions of HBP as with 30 and 50% v/v of HBP as

opposed to the unmodified composite. When HBP was

present in the PLLA matrix, the viscosity of the matrix

was noticed to have decreased during processing, which

would be consistent with the enhanced wetting. The en-

hanced wetting of the fibers by the matrix could provide a

contribution to the interlaminar fracture toughness, as the

calculatedGIC is very sensitive to the fiber-matrix interface.

Another distinguishing feature of the photographs was

evidence of matrix deformation revealed at higher magni-

fication in the propagation region, which was taken from

about themid-length of the specimen (Figure 10(a–d)). The

occurrence of plastic deformation on the specimen frac-

ture surface was greatest at higher proportions of HBP as

indicated by white fibrils. For the unmodified composite

(Figure 10(a)), some exposed fiber surface is observed,

which inferred that crack propagation may have proceeded

via the fiber-matrix interface as well as through the matrix

phase. In this instance, the bonding between the fiber and

the matrix maybe less than the bulk matrix, therefore

delamination also occurred via the interface. Some matrix

rupture is also evident (upper side of photograph) although

the matrix does not appear to have undergone plastic defor-

mation due to a less ductile matrix.

With 10% v/v of HBP, crack propagation may have pro-

ceeded through the matrix as no fiber surface was exposed.

This indicated that the strength between the fiber and the

matrixwas greater thanwithin the bulkmatrix and therefore

thematrix failed before the interface. The rough topography

of the matrix resembled a material with low ductility, but

not to the same extent as the unmodified composite as the

edges are not as sharp.

The propagation regions of the composites with 30 and

50% v/v of HBP showed similar behavior to that of 10% v/v

HBP but the ductility of the matrix became apparent with

increased proportions of HBP. At 30% v/v HBP, matrix

rupture still appeared to be less ductile but some plastic

deformation of thematrixwas observed as fibrils along each

side of the fiber. At 50% v/v HBP thematrix became ductile

and some drawing was visible, as seen in the upper left of

Figure 10(d). An increase in ductility may result in higher

strength of the matrix, contributing to higher GIC values.

Crack propagation also occurred along the fiber-matrix

interface for the composites with 30 and 50% v/v HBP, as

shown in Figure 8(c) and (d), respectively, with much resis-

tance and high energy release rate caused by ductile tearing

Figure 10. ESEM photographs of propagation region of the interlaminar fracturetoughness specimens with (a) 0% v/v, (b) 10% v/v, (c) 30% v/v, and (d) 50% v/v of HBP.



of thematrix at the interface. This particular failuremode in

conjunction with increased strength of the matrix contri-

buted to the significant increase in fracture toughness for

the composites with 30 and 50% v/v HBP.

Thermal Properties of Composites

The thermal properties of the PLLA composite with and

without HBP are presented in Table 4. The Tg’s presented

were taken from the maximum peak temperature from the

loss tangent curves.

As mentioned before, the Tc’s of the composite had

marginally decreased with HBP, which indicated that some

miscibility was achieved to cause inhibition of the crystal-

lization process of PLLA. Consequently, the Tg had

decreased from 76 8C for the unmodified composite to

58 8C at 50% v/v HBP. In an according manner, Xc of the

unmodified composite also was reduced markedly from

0.70 to 0.17 at 50% v/v HBP. Surprisingly, Tmwas approxi-

mately the same regardless of Xc or Tg. From these thermal

results, it can be deduced that the matrix had changed from

a brittle nature with high Tg and Xc to a ductile nature with

low Tg and Xc. The ductility may contribute to the overall

toughness observed, resulting in a rise in the loss tangent

peak heights and the higher GIC measured.

Conclusion

Hyperbranched polyesters of generation 3 were used as

toughness modifiers in PLLA-flax fiber composites, varying

between 10% v/v and 50% v/v. As the proportion of HBP

increased, the tensile modulus showed an opposing trend

whilst the elongation at maximum stress increased corres-

pondingly.A reduced tensilemoduluswithHBPand greater

elongations correlated well with a decreased Tg and Xc. A

dramatic influence of these properties was observed when

theHBP content was high (at 30%v/v and higher) due to the

HBP suppressing PLLA crystallization, as indicated by a

reduction in Tc.

Before the glass-rubber transition, the storage modulus

(E0) of the composites with HBPwere reduced but not to the

extent as those observed in the static modulus. At the glass-

rubber transition, the decrease in E0 of the composites with

HBP was more pronounced than with the unmodified com-

posites, as the molecules exhibited greater mobility. Conse-

quently, cold crystallization occurred as demonstrated by

the inflection after the glass-rubber transition and the occur-

rence of a Tc during the melting scan in DSC.

The loss factor (tan d) peak showed that there was a

diminishing effect on the Tg with greater proportions of

HBP. The peak intensity was also amplifiedmarkedly as the

HBP content increased, which implied that the toughness

of the composites was enhanced. This was further clarified

by the GIC calculated from the interlaminar fracture tough-

ness tests.

The interlaminar fracture toughness tests proved that

crack propagation generally proceeded in a stable manner

and the load versus displacement plots demonstrated that

larger forces were required to cause the same opening dis-

placement when the HBP concentration was higher.

Accordingly, the resistance curves showed higher GIC

values with an increase in HBP. The fracture toughness

was beneficially influenced, withGIC values at least double

even at the lowest HBP concentration tested. The main

contributing factor to high fracture toughness in this study

was due to better wetting of the fibers by the matrix when

HBP was present and with improved ductility, it caused

ductile tearing along the fiber-matrix interface during the

crack propagation. With the use of HBP as a toughness

modifier, the static and dynamic mechanical properties

were not detrimentally compromised to gain a large effect

on the fracture toughness.

Acknowledgement: Scholarships from the AustralianPostgraduate Award and the Co-operative Research Centre forPolymers (CRC-P) for Susan Wong are acknowledged.

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Mechanical Behavior and Fracture Toughness of Poly(L-lactic acid)Natural Fiber Composites Modified...

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Transcript of Mechanical Behavior and Fracture Toughness of Poly(L-lactic acid)Natural Fiber Composites Modified...