EUROPEAN

European Polymer Journal 42 (2006) 2140–2152

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POLYMERJOURNAL

Morphology and fracture behaviour of polyethylene/Mg–Allayered double hydroxide (LDH) nanocomposites

F.R. Costa a, B.K. Satapathy b, U. Wagenknecht a,R. Weidisch a,b,*, G. Heinrich a,*

a Leibniz-Institute of Polymer Research, Hohe Strasse 6, 01069 Dresden, Germanyb Institute of Materials Science and Technology, Friedrich-Schiller-University Jena, Löbdergraben 32, 07743 Jena, Germany

Received 8 December 2005; received in revised form 3 April 2006; accepted 3 April 2006Available online 23 June 2006

Abstract

Fracture behaviour of polyethylene (PE)/Mg–Al layered double hydroxide (LDH) based nanocomposites has beenstudied by essential work of fracture (EWF) approach. Transmission electron microscopy (TEM and X-ray diffraction(XRD) analysis have been used to investigate the morphological features of these nanocomposites. A maximum in thenon-essential work of fracture was observed at 5 wt.% LDH demonstrating enhanced resistance to crack propagation com-pared to pure PE. Morphological analyses of the nanocomposites show that the dispersed LDH platelets are partially exfo-liated and also forms clusters with polymer chains remaining entrapped within. Rheological analyses show that the typicallow-frequency Newtonian flow behaviour, as observed in unfilled polymer, shifts to shear-thinning behaviour with increas-ing LDH concentration. At 5 wt.% LDH a ductile-to-brittle transition has been observed. Fracture surface investigation bySEM reveals the arresting of the plastic crack growth by the LDH particle clusters, which is more significant at 5 wt.%LDH content. At higher LDH concentrations, the number of such particle clusters increases causing decrease in the aver-age distance between them. As a result large-scale plastic deformation of the matrix at higher LDH concentration is effec-tively arrested favouring small strain failure and this in turn reaffirms the possible existence of a ductile-to-brittletransition. The study in general reveals that the resistance against crack initiation (essential work of fracture: EWF)and crack propagation (non-essential work of fracture: bwp) in these nanocomposites are structurally correlated withthe matrix behaviour and the morphology (state of LDH particle dispersion) respectively.� 2006 Elsevier Ltd. All rights reserved.

Keywords: Morphology; LDH; Essential work of fracture; Crack toughness; Nanocomposites

0014-3057/$ - see front matter � 2006 Elsevier Ltd. All rights reserveddoi:10.1016/j.eurpolymj.2006.04.005

* Corresponding authors. Tel.: +49 351 4658 361; fax: +49 3514658 362.

E-mail addresses: Roland.Weidisch@uni-jena.de (R. Wei-disch), gheinrich@ipfdd.de (G. Heinrich).

1. Introduction

Layered double hydroxides (LDH) are wellknown for their applications in the fields like heter-ogeneous catalysis [1,2], biomedical applications[1,3], heat stabilizers and halogen scavengers forpolyvinyl chloride [4], etc. The use of LDH materials

.

mailto:Roland.Weidisch@uni-jena.demailto:gheinrich@ipfdd.de

F.R. Costa et al. / European Polymer Journal 42 (2006) 2140–2152 2141

in preparing polymer nanocomposites is rather anew field of application. The potential outcomescould be the development of flame retardant com-posites, improvement of conductivity, mechanicalreinforcement along with improved thermal stabil-ity, etc. Our interest in LDH is due to its high boundwater content and endothermic decomposition simi-lar to metal hydroxides, like Mg(OH)2 and ATH,which makes LDH suitable as flame retardant addi-tive in polymer matrices. The use of Mg–Al LDH asflame retardant in polyolefin composites is longknown, but its use as nano clay in such compositesafter suitable modification is very recent. This con-cept could be very useful in improving the dispersionof metal hydroxide type flame-retardants and hencereduce their loading requirement to obtain satisfac-tory flame retardancy. The study of mechanicaland fracture behaviours of such flame retardantcomposites is also important as these materialsoften find applications where load bearing capacitycould provide extra advantages. Specially, duringburning process the structural integrity of thefinished parts is necessary to stop rapid flame prop-agation through dripping. The investigations ofthe fractured surface morphology also provide usinsights of filler–polymer interaction and how thefiller–polymer interface behaves during loadingprocess. In the present article, we have reported indetail fracture mechanical behaviours and the frac-tured surface morphology of polyethylene/Mg–AlLDH based nanocomposites prepared by melt mix-ing technique.

The LDH as a class belongs to anionic clay andare represented by a general chemical formula½MII1�xMIIIx ðOHÞ2�

xþðAn�Þx=nmH2O, where MII is

divalent metal ion (like, Mg+2, Zn+2, etc.), MIII istrivalent metal ion (like, Al+3, Cr+3, etc.) and A isan anion with valency n (like CO�23 , Cl

�, etc.). Thehydroxide layers of LDH have close similarity tothe mineral brucite or Mg(OH)2, where some M

II

ions of brucite layers are substituted by a trivalentcation yielding positively charged ½MII1�xMIIIx -ðOHÞ2�

xþ layers. These positively charged layersremain stacked in z-direction with anions and thewater molecules in the interlayer space [1,5]. Thecharacteristic that makes LDH materials suitablefor polymer nanocomposite synthesis is theexchangeable nature of the inter layer anions bysuitable oligomeric or polymeric anion. The pristineMg–Al LDH clay was modified with an anionic sur-factant, sodium dodecylbenzenesulfonate (SDBS),using the regeneration method, which is based on

the well known ‘memory effect’ shown by LDHclays [1,6]. This modification process have alreadybeen discussed in detail elsewhere [7,8].

In recent years, characterization of fracturemechanical behaviour of polymer nanocompositesusing EWF concept has been reported for differentsystems [9–12]. Usually, in polymer nanocompositesaddition of nanofiller beyond certain concentrationmakes the composite preferably brittle and thus lim-its their applications where load-bearing ability isimportant. Hence it is worthwhile to investigatethe effects of filler loading on the fracture toughnessof the composite materials. The present studyfocuses on the deformation and fracture behaviourof PE–(Mg–Al) layered double hydroxide nanocom-posite systems using the essential work of fracture(EWF) approach, based on post-yield fracturemechanics (PYFM) concept.

2. Experimental and methodology

2.1. Materials

LDPE (density 0.9225 g/cm3, melt flow index,MFI 3.52 g/10 min) was purchased from ExxonMobil, the compatibilizer, maleic anhydride graftedpolyethylene (PE-g-MAH, polybond 3109, density0.926 g/cm3, MFI 32.2 g/10 min, MAH index 2.7),was purchased from Crompton Europe and sodiumdodecylbenzene sulfonate (SDBS) was purchasedfrom Aldrich chemical company. Mg–Al basedLDH was obtained from DUSLO a.s. Sala, Slova-kia and contained MgO/Al2O3 molar ratio 4.4,which corresponds to Mg/Al ratio of about 2.2.

2.2. Preparation of polyethylene/LDHnanocomposites

The preparation of the nanocomposites was car-ried out in a twin-screw extruder fitted with tightlyintermeshing co-rotating screws (Leistritz Micro27) and having screw diameter of 27 mm and L/Dratio of 36. The compounding was carried in twosteps: first, a master batch of the SDBS modifiedLDH in PE-g-MAH was prepared using 1:2 weightratio of the two components and second, the masterbatch was diluted with unmodified polyethylene bydifferent extent to obtain various concentrations ofLDH in the final composites. The nanocompositescontaining 1, 2.5, 5, 10 and 15 wt.% LDH were pre-pared in this way. Both steps i.e., master batch prep-aration and the dilution, were carried out in the

2142 F.R. Costa et al. / European Polymer Journal 42 (2006) 2140–2152

same twin-screw extruder using a temperature pro-file of 180–200 �C and feed rate of 6 kg/h. The nano-composite samples have been designated by‘PEXLDH’, where the letter ‘X’ is a number indicat-ing the weight percent of LDH presents in therespective sample.

2.3. Characterization of polyethylene/LDH

nanocomposites

2.3.1. Morphological analysis

The morphological analyses of the nanocompos-ites were carried out using both direct (X-raydiffraction, XRD and transmission electron micros-copy, TEM) and indirect methods (rheological anal-ysis). The XRD analysis over 2h = 1.8–40�, in stepsof 0.02� was carried out using X-ray diffractometerP4 (Siemens AG Karlsruhe, now BRUKER axsKarlsruhe) with Cu-Ka radiation (k = 0.154 nm,monochromatization by primary graphite crystal)generated at 30 mA and 40 kV. The XRD patternswere interpreted with respect to the position of thebasal peak (003), which depends on the distancebetween two adjacent metal hydroxide sheets inthe LDH crystal lattice. The higher order peaks ofthe same hkl series (00 6, 009 and so on) were alsoreported as they indicate the presence of repeatingcrystal planes and symmetry in a specific crystallo-graphic direction. The TEM was carried out atroom temperature using transmission electronmicroscope TEM LEO 912 with acceleration volt-age of 120 kV and bright field illumination. Theultra thin sections of samples were prepared by ult-ramicrotomy at �120 �C with a thickness of the sec-tion 80 nm.

Researchers have long established the rheologicalanalysis of polymer nanocomposite melt at lowdeformation rate as a very efficient technique tounderstand the state of clay particle dispersion inmatrix [13–16]. The melt rheological properties ofthe material were determined using an ARES-rhe-ometer (Rheometrics Scientific, USA). The mea-surements were performed in the dynamic modeusing 25 mm parallel plates geometry with gapsettings about 2 mm under nitrogen atmosphere.The strain amplitude was kept below 5% in thewhole frequency range to ensure viscoelastic linear-ity. The frequency sweep was carried out in therange 0.056–100 rad/s at temperature 240 �C.

The morphological analysis of the fractured sur-face was carried out by scanning electron micros-copy (SEM) using LEO 435 VP, Zeiss.

2.3.2. Fracture behaviour

Tensile properties of the nanocomposites weremeasured using universal tensile testing instrumentZwick 1456. The tests were carried out with dumb-bell shape specimen using standard ISO 527 –2/5A/5. The tensile properties were reported interms of tensile strength, elongation at break andE-modulus.

The EWF was determined from rectangular com-pression moulded sheets of 80 mm · 20 mm ·0.5 mm dimension. The samples were pre-notchedwith different ligament lengths varying from 2 mmto 10 mm. The fractures of these pre-notched speci-mens were carried using universal tensile testinginstrument using constant extension speed to obtainindividual load displacement curves. The essentialwork of fracture (EWF) method has been usedallowing distinguishing between two terms repre-senting the resistance to crack initiation and crackpropagation corresponding to inner fracture processzone and outer plastic deformation zone. The appli-cability of EWF to nanostructured polymers andnanocomposites has been demonstrated in recentpapers [10]. The precondition for the validity ofEWF approach has been demonstrated by the self-similarity nature of the load–displacement diagramsof these blends. In this study, the fracture mechani-cal tests (test speed: 1 mm/min, room temperature)were performed on double-edge-notched-tension(DENT) specimens by a universal testing machinewith mechanical grips (Z010 Zwick) and interfacedwith a software-assisted video monitoring system.Double edge-notched tension (DENT) specimenswith total length of 80 mm, thickness of 0.5 mmand width of 20 mm, respectively, were cut fromthe compression moulded plates. The clamp dis-tance was 40 mm. For notching, a special devicewith fresh razor blades with notch tip radius of0.20 lm was used to realize that both notches aresimilarly sized. For each material, at least 10 speci-mens were tested with different ligament lengthbetween ca. 2 mm and ca. 10 mm.

For plane-stress conditions, the total work offracture W dissipated in a notched specimen canbe divided into a component We characterizing theinner or fracture process zone and another one Wpcorresponding to an outer zone

W ¼ W e þ W p ¼ EWF � B � lþ bwp � B � l2 ð1Þ

where B, l and b are specimen thickness, ligamentlength and shape factor of the plastic zone, respec-

F.R. Costa et al. / European Polymer Journal 42 (2006) 2140–2152 2143

tively. After dividing W by the ligament (notched)area B Æ l the specific work of fracture w is obtained

w ¼ EWFþ bwp � l ð2Þ

Based on the fact that the intrinsic fracture pro-cess takes place in the inner zone, the term EWF,the essential work of fracture, is experimentallydetermined by extrapolation of w as a function ofl to zero ligament length. For this, several similar-sized specimens with different ligament lengths aremonotonically loaded. bwp (non-essential work offracture) is the slope of the linear fit extrapolation.

3. Results and discussion

3.1. XRD and TEM analysis

The XRD patterns for the modified LDH clayand the nanocomposites are shown in Fig. 1. Thebasal peak (003) position in the SDBS intercalatedLDH corresponds to an interlayer layer distance of29.6 Å, which is nearly four-fold in comparison tothe pristine clay material, where it is 7.7 Å. In thenanocomposites, the presence of basal peak indi-cates that the LDH layers are not exfoliated. How-ever, a small shift (by about 3.2 Å) to lower 2hvalue is observed in case of all the nanocompositesamples. Another interesting observation is the dis-appearance or weakening of the higher order peaksin the nanocomposites, which could be due lossof crystalline symmetry in the stacking direction ofthe hydroxide layers and lowering of number of

2 4 6 8 10 12 14 16

(009)(006)

(003)3.27 nm 2.95 nm

Inte

nsity

(a.

u.)

2 theta (deg)

LDH-SDBS PE1LDH PE2.5LDH PE5LDH PE10LDH PE15LDH

Fig. 1. XRD pattern for SDBS modified Mg–Al LDH (LDH-SDBS) and filled nanocomposites.

hydroxide layers [17]. The XRD patterns of thesekinds, therefore, point towards the intercalated orpartially exfoliated nature of the polyethylene/LDH nanocomposites. The TEM analysis supportsthis observation. The TEM micrographs for thenanocomposites containing 5 and 10 wt.% LDH(Fig. 2) show the presence of both exfoliated layersand intercalated tactoids of LDH crystallites. Thetactoids are mostly in the form of thin platelets hav-ing thickness in the range 30–50 nm and lateral

Fig. 2. TEM micrographs of polyethylene/LDH nanocomposites(a: 5 wt.% and b: 10 wt.% LDH loading) showing the partiallyexfoliated nature. The presence of intercalated tactoids is alsoapparent.

2144 F.R. Costa et al. / European Polymer Journal 42 (2006) 2140–2152

dimension ranging from few hundred nm to about2 lm. It has also been observed that the exfoliatedlayers are quite uniformly dispersed within thematrix, whereas the tactoids tend to form domainsof physically network structure especially at higherLDH concentration, resulting in an inhomoge-neous distribution of LDH particles throughoutthe matrix.

3.2. Rheological analysis

The linear viscoelastic response of the nanocom-posite melt is strongly influenced by the factors likestate of particle dispersion, inter-particle and parti-cle–polymer interactions. Therefore, rheologicalanalysis in the linear viscoelastic regime is widelyused to characterize polymer nanocomposites. Inaddition, the processing behaviour of these materi-als could also be obtained from their rheologicalcharacteristics. The complex viscosity versus fre-quency plots for the polyethylene/LDH nanocom-posites are shown in Fig. 3. The unfilled matrix(PE and PEPB) show classical viscoelastic behav-iour characterized by transition from low frequencyNewtonian flow behaviour to high frequency shearthinning nature. In case of the nanocomposites,the presence of LDH significantly changes thisbehaviour, especially in the low frequency regionand above certain level of LDH concentration. Inthe low-frequency region, it is observed that theNewtonian flow behaviour is gradually changed toshear-thinning behaviour when the melt shows atransition from liquid-like state to pseudo-solid-likestate. The degree of this transition with increasingLDH concentration could be studied by an index

0.1 1 10 100102

103

104

105

Materials -n in |η*|~ ωn

PE 0.13PEPB 0.04PE1LDH 0.19PE2.5LDH 0.23PE5LDH 0.41PE10LDH 0.71PE15LDH 0.82

Shear thinning behaviour

Newtonian behaviour

PE PEPB PE1LDH PE2.5LDH PE5LDH PE10LDH PE15LDH

Com

plex

vis

cosi

ty, |

η*|

(Pa

.s)

Frequency, ω (rad/s)

-

Fig. 3. Complex viscosity versus frequency plots for polyethyl-ene/LDH nanocomposites. (PEPB is the blend of polyethyleneand PE-g-MAH corresponding to the nanocomposite containing15 wt.% LDH). The variation of shear thinning exponent ‘n’ withLDH concentration has been shown in the table right.

called shear-thinning exponent, ‘n’, which isobtained by fitting the low-frequency data inFig. 3 to power law equation of the type jg*j � xn.Newtonian liquids have the value of n equal to zero.The unfilled polyethylene matrix behave more likelyso with small negative value of n, whereas the nano-composites show significant deviation from theNewtonian behaviour. With increasing LDH con-centration the negative value of n increases steadily(Fig. 3) showing the developing shear-thinning nat-ure. This type of low-frequency rheological responseis typical for polymer nanocomposites, where relax-ation of the system is strongly lowered due to teth-ering of polymer chains on the particle surface andparticle–particle interactions [13]. While analysingthe rheological behaviour of three component poly-mer nanocomposite system, like the present case,the effect of compatibilizer phase should be kept inmind, especially when its molecular weight is muchsmaller than the main polymer. The MFI value ofPE-g-MAH used in the present studies is aboutten times higher than that of the polyethylene. Asa result the average molecular weight of the matrixdecreases with increasing weight fraction of PE-g-MAH and hence lowers the overall viscosity of thematrix phase. This is evident in Fig. 3 and to someextent counter balance the effect of LDH loading onoverall increase in the viscosity of the system. How-ever, the comparison in terms of shear-thinningexponent provides a better understanding of theeffects of LDH loading on low frequency rheologi-cal behaviour. Wagner and Reisinger [18] haverelated the increasing value of ‘n’ with the increasingdegree of exfoliation of the dispersed clay particles.The XRD patterns of the nanocomposites do notshow enhancement in the extent of exfoliation ofLDH layers with increasing LDH concentration.Therefore, the change in the shear-thinning expo-nent with LDH concentration is not due to higherextent of exfoliation. Rather, the increasing LDHconcentration causes two changes in the system.Firstly, an increasing number of polymer chains/segments becomes tethered on the LDH particlesurface or in the inter gallery space and secondly,the average distance between the dispersed particleis decreased [19].

3.3. Mechanical properties

The tensile properties of the polyethylene/LDHnanocomposites are shown in Fig. 4. It is apparentthat the LDH clays does not cause any mechanical

PE PE1LDH PE2.5LDH PE5LDH PE10LDH PE15LDH PEPB10

12

14

16

18

200

250

300

350

400

Elongation at break (%

)

Ten

sile

str

engt

h, E

-mod

ulus

(M

Pa)

Tensile strength E-Modulus

20

40

60

80

100

120 Elongation at Break

Fig. 4. The tensile properties of the polyethylene/LDH nano-composites and their comparison with unfilled matrices.

0 1 2 3 4 5 6 70

10

20

30

40

50

60

70

(a)

(b)

PE5LDH

E

D

CB

A

Loa

d (N

)

Displacement (mm)

A ~ 2mmB ~ 4mmC ~ 6mmD ~ 8mmE ~ 10mm

2 4 6 8 10

0

10

20

30

40

σ n (

N/m

m² )

Ligament length (mm)

PE PE1LDH PE2.5LDH PE5LDH PE10LDH PE15LDH

Fig. 5. (a) The load–displacement diagram for polyethylene/LDH nanocomposites showing self similarity; (b) Hill’s analysisshowing the non-dependence of net section stress on ligamentlength indicating plane-stress criterion.

F.R. Costa et al. / European Polymer Journal 42 (2006) 2140–2152 2145

reinforcement of the composites compared to theunfilled PE matrix. Apparently the tensile strengthshows a decreasing trend with an increase in LDHconcentration. However, increasing LDH concen-tration brings about steady increase in modulusand also a sharp decrease in the elongation at break.In the nanocomposites, the percentage of lowmolecular weight compatibilizer (that is the maleicanhydride grafted polyethylene) is also increasedalong with the increasing LDH concentration. It isbelieved that this has a deteriorating effect on thetensile strength of the matrix (the comparison ofpure polyethylene and polyethylene/compatibilizerblend in Fig. 4 demonstrates this fact). In fact thechange in the tensile properties of the compositeswith increasing LDH concentration is a combinedeffect of dispersed LDH particles and the compatibi-lizer. The two influence the tensile properties inopposite directions. The reinforcing nature of theLDH clay could actually be realized in the composi-tion containing 15 wt.% LDH, when compared withsimilar unfilled matrix composition (PEPB inFig. 4). The comparison further reveals that E-mod-ulus is increased by about 40% and tensile strengthis increased from about 12.17 MPa in PEPB toabout 15.17 MPa in case of the nanocomposite with15 wt.% of LDH content. However, the elongationat break is reduced drastically. This also indicatesthat at 15 wt.% LDH concentration, the energyabsorbed by the nanocomposites drops significantlyand the material tends to become even more brittlein comparison to the composites with lower LDHcontent and the unfilled matrix.

3.4. Fracture behaviour of the nanocomposites

3.4.1. Load–displacement diagramsIn general, polyethylene/LDH nanocomposites

have shown thermoplastic behaviour and the valid-

ity of the EWF methodology through the self-simi-lar nature of the force–displacement curves and areshown in Fig. 5a. On the other hand the plane-stresscriterion was confirmed by Hills analysis [20], whichrevealed that the net-section stress remained inde-pendent of the ligament length (Fig. 5b). From thein situ video equipment observation during the testit is confirmed that the failure of the DENT speci-mens occurred after ligament yielding, which is alsoapparent from their corresponding strain fieldimages. Full yielding of the ligament in the DENTspecimens occurred at maximum load (Fmax) andbefore crack propagation started. Clearly thisbehaviour observed for these nanocomposite sys-tems is in agreement with conditions of the EWFconcept.

0 2 4 6 8 10 12 14 160.01

0.02

0.03

0.04

0.05

0.06

0.07

0.08Ligament length ~ 6mm

(b)

s max

/Fm

ax (

mm

/N)

LDH content (wt%)

0 1 2 3 4 5 6 70

5

10

15

20

25

30

35

40

(a)

F

E

D

C BA

Loa

d (N

)

Displacement (mm)

A - PE0LDHB - PE1LDHC - PE2.5LDHD - PE5LDHE - PE10LDHF - PE15LDH

Fig. 6. The influence of LDH concentration on: (a) load–displacement behaviour and (b) smax/Fmax ratio.

2146 F.R. Costa et al. / European Polymer Journal 42 (2006) 2140–2152

It is evident from the Fig. 6a that the incorpora-tion of LDH into PE causes a gradual transition inthe nature of the load–displacement curves (for lig-ament length of ca. 6 mm) characterized by a notice-able increase in the load at low LDH loading and areduction in the maximum displacement. This givesa qualitative indication of the existence of a ductile-to-brittle transition with increasing LDH content.As a result of the increasing load and decreasing dis-placement the ratio of maximum displacement(smax) and maximum load (Fmax) (ductility ratio)decreases as well (Fig. 6b). When the LDH contentwas increased from 2.5 to 5 wt.% the maximum dis-placement (smax) decreased from 6 to 3 mm and themaximum force (Fmax) decreased from 36 to 32 N.On further increasing the LDH content from 5to 10 wt.% the maximum displacement (smax)decreased from 3 to 1.5 mm, while the maximum

force (Fmax) decreased to some extent indicating adecrease in the crack tip plasticity leading to brittlefailure of the materials (Fig. 6b). In case of nano-composites with a higher LDH concentration thanthe critical level, the entrapment of the matrix PEchains between the LDH platelets and the clusteringof such platelets become decisive in controlling theoverall deformation behaviour of the material.These clusters restrict the plastic deformation atthe cluster–matrix interface causing a sharpdecrease in ductility of the composite matrix athigher LDH concentration.

3.4.2. Crack toughness behaviour

The crack toughness behaviour was studied onthe basis of essential work of fracture concept,where the non-essential work of fracture (bwp) is ameasure of the resistance against crack propagation.The slope of the linear-fit plot of specific work offracture against the ligament length represents bwpand has equivalence in the J-integral concept [21].The equivalence between EWF concept and R-curveconcept using the J integral has been shown by Maiand Cotterell [21,22]. They found, that EWF = JIc,the critical J value as a measure of the stable-crackinitiation toughness, and bwp = 1/4 Æ dJ/da (forDENT specimens) where dJ/da is the slope of theJ–R curve (a – actual crack length). Some authorsconfirmed these findings experimentally correspond-ing to EWF and bwp [23] or numerically (EWF) [24].

The plot of specific work of fracture against theligament length is shown in Fig. 7a. The variationin EWF with LDH concentration is shown inFig. 7b. It has been clearly observed that the magni-tude of EWF remained comparatively unaffected upto 2.5 wt.% LDH content suggesting the crack initi-ation to be primarily matrix-controlled. However, asharp drop in EWF could be observed on furtherincreasing the LDH concentration to 5 wt.% indi-cating a decreased resistance to crack initiation. Itshould be noted that when the tensile strength wasnot showing any significant variation (Fig. 5),EWF approach could distinctly reflect the transitionin the nature of material response with increasingLDH content. The slope of the linear regression fitof the data points in the plot represents the non-essential work of fracture (bwp) decreases by�30% on incorporation of 2.5 wt.% of LDH ascompared to pure PE (Fig. 7c). Interestingly, on fur-ther increasing the amount of LDH to 5 wt.% bwpincreased nominally when compared to the nano-composite with 2.5 wt.% LDH, though it still

2 4 6 8 100

10

20

30

40

50

60

70

(a)

PE PE1LDH PE2.5LDH PE5LDH PE10LDH PE15LDH

Spec

ific

wor

k of

fra

ctur

e (k

J/m

2 )

Ligament length (mm)

0 2 4 6 8 10 12 14 16

0

5

10

15

20

25

30

(b)

Ess

entia

l wor

k of

fra

ctur

e (N

/mm

)

LDH content (wt.%)

0 3 6 9 12 15

0

1

2

3

4

5

6

(c)

βwp

(N/m

m2 )

LDH content (wt.%)

Fig. 7. Plots showing (a) specific work of fracture with ligamentlength and LDH concentration and (b) EWF with LDHconcentration and (c) non-essential work of fracture (bxp) withLDH loading in the nanocomposites.

0 2 4 6 8 10 12 14 16

0.00

0.05

0.10

0.15

0.20

0.25

Cra

ck o

peni

ng d

ispl

acem

ent,

CO

D (m

m)

LDH concentration (wt%)

Fig. 8. The influence of LDH concentration on the crack openingdisplacement (COD) in the polyethylene/LDH nanocomposites.

F.R. Costa et al. / European Polymer Journal 42 (2006) 2140–2152 2147

remained lower than that of pure PE. However, asharp drop in the bwp values by �60% was observedwhen the LDH content was increased to 10 wt.%.

The existence of such a non-linear dependence ofbwp on the content of LDH indicates a tough-to-brittle transition in these polymer nanocomposites.Qualitatively, at 5 wt.% of LDH such tough-to-brit-tle transition might be due to inherent morphologi-cal changes. The dependence of bwp on the amountof LDH might be attributed to a reduced mobilityof the PE chains at higher LDH concentrationsdue to an overall increase in the interfacial area.

The crack opening displacement (COD) has beencalculated based on the empirical relationshipwe = Mry COD, where ry is the yield stress and Mis the plastic constraint factor which is taken as1.15 for double edge notched tension (DENT) spec-imens [22,25]. The COD as a function of the LDHcontent is plotted in Fig. 8 and exhibits a similarityin the trend to that of bwp. The COD decreased (likebwp) with the increase in the LDH content to2.5 wt.%, while on increasing the content to5 wt.% the COD increased and subsequentlyshowed a decreasing trend as the concentration ofLDH was increased further up. Conceptually,COD is a measure of the inherent resistance of thematerial to stable crack propagation by crack blunt-ing (alternatively, COD is related to deformationcapacity of the material) and is related to the sizeof outer plastic deformation zone in the post-yieldfracture process zone. The observed local maximumin bwp at 5 wt.% LDH which has a correspondenceto the COD primarily indicates that the crackpropagation behaviour in these nanocomposites iscontrolled by the blunting of the advancing crackin the frontal process zone. The increase in theCOD values (crack-blunting efficiency) in the

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concentration range of 2.5–5 wt.% LDH, which isseen from Fig. 8 indicates a critical change in thedispersion state of the LDH particles. The fracturesurface morphology, which will be described indetails later, shows that within this concentrationrange the primary LDH particles (platelets) startforming secondary structures or clusters that entrappolymer chains. The interface also shows better par-ticle–polymer adhesion allowing more efficientstress transfer from the matrix to the particleclusters. However, on further increase in LDH con-centration, the average distance between these sec-ondary structures reduces drastically resulting fastercrack propagation and brittle failure of the material.This is also evident in the load–displacement dia-gram in Fig. 6a, where the load after the maximumvalue drops sharply to zero value at concentrationsabove 5 wt.% indicating less energy required forcrack propagation. Similar dependence of the cracktoughness behaviour has also been reported in caseof SBR/carbon black, PMMA/silica and in PC/MWNT based nanocomposites [12,26–29].

The effect of particle–particle cluster formation isalso evident from the sharp increase in the shear-thinning exponent above 5 wt.% LDH concentra-tion as discussed in the previous section. The partialexfoliation of the LDH clay together with anincreasing number of particle clusters causes inho-mogeneous distribution of filler beyond 5 wt% con-centration. The strong particle–particle interactionand their thermodynamic incompatibility withnon-polar polyethylene matrix act as the drivingforce behind this kind of particle clustering at higherfiller concentrations. Such inhomogeneous distribu-tions might also cause high strain localization andare also attributed to the extreme modulus mis-match between the LDH and the PE matrix [30].The decrease in the J-integral fracture toughnessabove a critical clay content has also been reportedin case of PP/CaCO3 nanocomposites [31]. Thus thenanocomposites with a composition range of 2.5–5 wt.% LDH indicate the possible composition win-dow in which the polymer–filler interaction buildsup while above 5 wt.% LDH two competitive eventsin the form of polymer–LDH and filler–filler inter-action occur. At a concentration above 10 wt.%LDH the filler–filler type network formation iskinetically favoured promoting the nanocompositesto undergo brittle failure with unstable (high speed)crack propagation. Thus these nanocompositesundergo a tough-to-brittle transition at 5 wt.%LDH.

This could be attributed to the fact that up to thecritical concentration range (2.5–5 wt.%) the nano-composites undergo preferably plastic deformationas the dispersed primary particles or exfoliated frag-ments do not show strong interfacial adhesion withthe matrix. At higher LDH concentration, immobi-lization of the matrix takes place around the second-ary particle–particle structure due to entrapment ofthe polymer chains/segments within such structures.Such immobilization restricts the plastic deforma-tion of the matrix and crack blunting at the interfa-cial region. Similar observations have been reportedfor other polymer-based composites reinforced withactive fillers, which form a surrounding layer ofimmobilized polymer (such as rubber/carbon blackand PMMA/silica). The crack initiation toughness(as resistance against stable or unstable crack initia-tion) of these materials is increased as a function offiller volume fraction primarily due to increasingenergy dissipation capacity of the composite, whichis evident from the increasing loss modulus (G00) inDMA [27–29]. However, in case of nanocompositeswith LDH contents more than 5 wt.% the formationof large number of secondary filler particle struc-tures (particle clusters) and their close proximityto each other in the matrix leads to large magnitudecrack propagation instability. This inevitably reaf-firms the existence of a ductile-to-brittle transitionat 5 wt.% LDH as already discussed in the contextof non-essential work of fracture (bwp) andductility.

Thus the different crack toughness behaviour ofthe composites with varied LDH contents may beattributed to the change in matrix morphologyand the filler–matrix interaction. At higher LDHcontents crazing becomes predominant where theinitiation and growth of such crazes are affectedby both the matrix morphology and the filler–matrix interaction. Crack growth in these nanocom-posites is primarily determined by the crack tip plas-tic deformation which in turn is depending on twocompetitive events occurring at the crack tip i.e.,viscous and shear deformation across the length ofthe developing cracks giving rise to characteristicfractured surface morphology patterns which is dis-cussed in the subsequent section.

3.5. Fracture surface morphology

The SEM analysis of the post-yield fractured sur-faces of the nanocomposites is shown in Fig. 9. Theunfilled polyethylene and the nanocomposites with

Fig. 9. Fracture surface morphology of polyethylene/LDH nanocomposites: a – PE0LDH; b – PE1LDH; c – PE2.5LDH; d – PE5LDH;e – PE10LDH; f – PE15LDH (magnification bar in A is 10 lm and others 20 lm). The highlighted circular regions in d, show the possibleplatelet clusters and their good adhesion with matrix.

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1–2.5 wt.% LDH content have primarily undergoneductile failure accompanied with plastic and sheardeformation. The volume of polymer that under-goes yielding determines the total energy absorptionand the ultimate mode of fracture in case of filledthermoplastics. With concentration of LDH below5 wt.% the fractured surface reveals the presenceof primary LDH particles in the form of platelets,which is due to the lower tendency to form agglom-erates at these low concentrations. It has alreadybeen confirmed from TEM that the LDH plateletsare quite well intercalated and also partially exfoli-ated giving nano scale dispersion of the particle

fragments. The composites containing 5 wt.% orhigher amount of LDH show the presence of sec-ondary structures formed by clustering of LDHplatelets. The large scale plastic flow characteristicsas observed in unfilled polyethylene and at lowLDH concentration are greatly reduced in thesecases as revealed from the failure surface topogra-phy indicating small scale yielding of the polymeraround the particle clusters in Fig. 9d–f. Thematrix–cluster interface in the fractured surfacecontains large number of pullout strands of matrixpolymer indicating bridging effect. Fig. 10 shows amagnified view of such a cluster, where this bridging

Fig. 10. SEM image showing the nature of matrix failure at theinterface between a LDH particle cluster and the bulk matrix.

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effect and also the presence of large number ofmicro voids at the interface are distinctly visible.Though the LDH clay is premodified with surfac-tant SDBS, the hydrocarbon chain of SDBS is tooshort (only slightly longer than C12 length) to giveentanglement effect with the matrix. Also singleLDH platelet does not show such interfacial adhe-sion with matrix (Fig. 9c). Therefore, only possibil-ity exists is the entrapment of considerable amountof matrix polymer within the particle clusters. Thisalso facilitates the micro-void formation aroundthe particle clusters favouring inter particle matrixyielding and thereby increasing toughness in caseof filled polymers [32,33]. Conceptually, the absenceof agglomerates/clusters favours multiple yieldingwhich eventually supports the dissipation of thedeformation energy and hence causes the reductionin the crack growth driving force. This observationof increased toughness in case of the nanocompositewith 5 wt.% LDH content is in agreement with theobserved ductile-to-brittle transition. However, thetendency to form clusters has been observed to beenhanced as the content of LDH is increasedbeyond 5 wt.% facilitating extensive voiding (cavita-tions) which are mainly initiated at the polymer–LDH cluster interface eventually leading to theformation of highly developed localized strains asevident from the fibrilar polymer network as shownin Fig. 9e and f. In such cases the crack growthoccurs easily through filler agglomerates that lieacross the crack propagation path causing pre-mature low-energy failure [34,35]. On furtherincreasing the concentration (say 15 wt.%), thenanocomposites undergo brittle fracture becauseof larger aggregates/clusters of LDH making the

crack propagation path much more vulnerable tolow-strain failure (Fig. 9f).

4. Conclusion

A tough-to-brittle transition was observed in thepolyethylene/Mg–Al LDH based polymer nanocom-posites between 5 and 10 wt.% LDH concentration.At 5 wt.% LDH a local maxima in the non-essentialwork of fracture (bwp) was observed indicatingincreased resistance against crack propagation. Thenon-essential work of fracture (bwp) has been foundto be a sensitive parameter for characterizing cracktoughness behaviour of polymer–LDH clay nano-composites. The study further reveals that highercontent of LDH could be detrimental to the fracturetoughness in this type of nanocomposites. This couldbe of immense significance in developing flame-retardant composites having good mechanical andfracture properties. The high magnification TEMand low magnification SEM images show that theLDH particles exist in the matrix in different forms,starting from nanoscale exfoliated fragments tomicron sized particle clusters. These different statesof particle dispersion influence the failure behaviourof the matrix in different manner: the primary parti-cle usually show poor interfacial adhesion, whileparticle clusters show evidence of stress transferfrom the matrix phase to the filler (particle) phasethrough the interface. The nanocomposites showpreferably brittle failure above 5 wt.% LDH concen-tration. Rheological analysis also indicates thatabove this concentration the melt deviates largelyfrom the low-frequency Newtonian behaviour indi-cating the possible formation of some unstable fillernetwork structure that breakdowns on shearing.However, formation of 3D filler network structuresseem unlikely due to thermodynamic reason ratherformation of localized domains of such networkstructures or clusters by the primary LDH particlesbrings about specific changes in the material behav-iour. Fracture surface investigation by SEM showsarresting of the plastic crack growth by these clustersat 5 wt.% LDH content. The correspondence ofEWF and bwp to deformation capacity of pure PEand the state of dispersion of the LDH in the PEmatrix respectively. This explains the correlation ofresistance against crack initiation with matrix behav-iour and the resistance against crack propagationwith morphology. The increased toughness due toaddition of LDH at 5 wt.% has been correlated withthe effective resistance to damage initiation caused

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by the stress transfer at the matrix–cluster interfaceand hence arresting the crack. The present studyshows the applicability of the EWF concept toinvestigate the fracture toughness behaviour of poly-mer nanocomposites and its correlation with thedispersion state of the clay particles in the matrix.Our results also suggest that higher content ofLDH can potentially limit the application of LDHbased nanocomposites as flame-retardant compos-ites where improved fracture toughness and failureproperties are desired. However, more detailedinvestigations are necessary regarding the nature ofsurfactant used for modification of pristine LDHand the influence of processing conditions as bothfactors could largely influence the final morphologyof the nanocomposites and hence the fracturemechanical properties.

Acknowledgements

Authors gratefully acknowledge DUSLO a.s.Sala, Slovakia for providing Mg–Al based LDHmaterial and Dr. Dieter Jehnichen of IPF Dresdene.V. for XRD measurements.

References

[1] Cavani F, Trifiro F, Vaccari A. Hydrotalcite-type anionicclays: preparation, properties and applications. Catal Today1991;11:173–301.

[2] Tichit D, Coq B. Catalysis by hydrotalcites and relatedmaterials. Cattech 2003;7(6):206–17.

[3] Choy JH, Kwak SY, Park JS, Jeong YJ. Cellular uptakebehavior of [c32-P] labeled ATP-LDH nanohybrids. J MaterChem 2001;11:1667–71.

[4] Lin YJ, Li DQ, Evans DG, Duan X. Modulating effect ofMg–Al–CO3 layered double hydroxide on thermal stabilityof PVC resin. Polym Degrad Stab 2005;88(2):286–93.

[5] Reichle WT. Synthesis of anionic clay minerals (mixedmetal hydroxides, hydrotalcite). Solid State Ionics 1986;22:135–41.

[6] Miyata S. Physico-chemical properties of synthetic hydro-talcites in relation to composition. Clays Clay Miner1980;24:50–6.

[7] Costa FR, Leuteritz A, Wagenknecht U, Jehnichen D,Heinrich G. Intercalation of Mg–Al based layered doublehydroxide by anionic surfactants containing different func-tionalities using regeneration method. Appl Clay Sci,submitted for publication.

[8] Costa FR, Goad MA, Wagenknecht U, Heinrich G.Nanocomposites based on polyethylene and Mg–Al layereddouble hydroxide: I. Synthesis and characterization. Polymer2005;46:4447–53.

[9] Qiao Y, Avlar S, Chakravarthula SS. Essential fracture workof nylon 6-silicate nanocomposites. J Appl Polym Sci2005;95:815–9.

[10] Bureau MN, Perrin-Sarazin F, Ton-That MT. Polyolefinnanocomposites: essential work of fracture analysis. PolymEng Sci 2004;44:1142–51.

[11] Lach R, Schneider K, Weidisch R, Janke A, Knoll A.Application of essential work of fracture concept to nano-structured polymer materials. Eur Polym J 2005;41:383–92.

[12] Satapathy BK, Weidisch R, Pötschke P, Janke A. Cracktoughness behaviour of multiwalled carbon nanotube(MWNT)/polycarbonate nanocomposites. Macromol RapidCommun 2005;26:1246–52.

[13] Krishnamoorti R, Gianeelis EP. Rheology of end-tetheredpolymer layered silicate nanocomposites. Macromolecules1997;30:4097–102.

[14] Ren J, Silva AS, Krishnamoorti R. Linear viscoelasticity ofdisordered polystyrene – polyisoprene block copolymerbased layered – silicate nanocomposites. Macromolecules2000;33:3739–46.

[15] Lee KM, Han CD. Rheology of organoclay nanocomposites:effects of polymer matrix/organoclay compatibility and thegallery distance of organoclay. Macromolecules 2003;36:7165–78.

[16] Solomon MJ, Almusallam AS, Seefeldt KF, Somwangtha-naroj A, Vardan P. Rheology of polypropylene/clay hybridmaterials. Macromolecules 2001;34:1864–72.

[17] Vaia RA, Liu W. X-ray powder diffraction of polymer/layered silicate nanocomposites: model and practice. J PolymSci Part B: Polym Phys 2002;40:1590–600.

[18] Wagner R, Reisinger TJG. A rheological method to comparethe degree of exfoliation of nanocomposites. Polymer2003;44:7513–8.

[19] Costa FR, Goad MA, Wagenknecht U, Heinrich G.Nanocomposites based on polyethylene and Mg–Al layereddouble hydroxide: II. Rheological characterization. Polymer2007;47:1649–60.

[20] Hill RH. On discontinuous plastic states, with specialreference to localized necking in thin sheets. J Mech PhysSolids 1952;1:19–30.

[21] Mai YW, Cotterell B. On the essential work of ductilefracture in polymers. Int J Fract 1986;32:105–25.

[22] Hashemi S. Work of fracture of high impact polystyrene(HIPS) film under plane stress conditions. J Mater Sci2003;38:3055–62.

[23] Chen YH, Mai YW, Tong P, Zhang LC. Numericalsimulation of the essential work of fracture method. In:Williams JW, Pavan A, editors. Fracture of polymers,composites and adhesion. ESIS Publication, 27. Amster-dam: Elsevier; 2000. p. 175.

[24] Arkhireyeva A, Hashemi S, O‘Brien M. Factors affectingwork of fracture of uPVC film. J Mater Sci 1999;34:5961–74.

[25] Hashemi S, Williams JG. Plast Rubb Compos 2000;29:294.

[26] Grellmann W, Cäsar T, Heinrich G. Mechanische Eigens-chaften, Rissinitiierung und molekulare Struktur gefüllterVulkanisate. Kauts Gummi Kunst 1999;52:37–43.

[27] Grellmann W, Heinrich G, Cäsar T. Crack initiation, wearand molecular structure of filled vulcanized materials. In:Grellmann W, Seidler S, editors. Deformation and fractureof polymers. Berlin, Heidelberg: Springer; 2001. p. 479.

[28] Reincke K, Lach R, Grellmann W, Heinrich G. Investigationsof crack propagation behaviour of unfilled and filled vulca-nizates. In: Grellmann W, Seidler S, editors. Deformation

2152 F.R. Costa et al. / European Polymer Journal 42 (2006) 2140–2152

and fracture of polymers, Springer, Berlin, Heidelberg, 2001.p. 493.

[29] Lach R, Kim GM, Michler GH, Grellmann W, Albrecht K.Indentation fracture mechanics toughness assessment ofPMMA/SiO2 nano-composites, Macromol Mater Eng, inpreparation.

[30] Ajayan PM, Schadler LS, Braun PV. Nanocomposite scienceand technology. Wiley-VCH Verlag; 2003.

[31] Chan CM, Wu J, Li JX, Cheung YK. Polypropylene/calciumcarbonate nanocomposites. Polymer 2002;43:2981–92.

[32] Suetsuga Y. In: Conf proc MOFFIS, 37, Namur, Belgium,1993.

[33] Seelig T. On micromechanical modeling of tougheningmechanisms and failure in amorphous thermoplastic poly-mer blends. Vom Fachbereich Mechanik der TechnischenUniversität Darmstadt genehmigte Habilitationsschrift,2004.

[34] Haworth B, Raymond CL, Sutherland I. Polyethylenecompounds containing mineral fillers modified by acidcoatings. 2: factors influencing mechanical properties. PolymEng Sci 2001;41:1345–64.

[35] Kinloch AJ, Young RJ, Fracture behaviour of polymers.Applied science. London; 1983.

Morphology and fracture behaviour of polyethylene/Mg-Al layered double hydroxide (LDH) nanocompositesIntroductionExperimental and methodologyMaterialsPreparation of polyethylene/LDH nanocompositesCharacterization of polyethylene/LDH nanocompositesMorphological analysisFracture behaviour

Results and discussionXRD and TEM analysisRheological analysisMechanical propertiesFracture behaviour of the nanocompositesLoad-displacement diagramsCrack toughness behaviour

Fracture surface morphology

ConclusionAcknowledgementsReferences