JOURNAL OF MATERIALS SCIENCE 36 (2 0 0 1 ) 451 – 460 Low temperature fracture properties of polymer-modified asphalts relationships with the morphology L. CHAMPION, J.-F. GERARD * Laboratoire des Mat ´ eriaux Macromol ´ eculaires, UMR CNRS 5627, INSA Lyon B ˆ at. 403, 69621 Villeurbanne Cedex, France E-mail: [email protected]J.-P. PLANCHE, D. MARTIN Centre de Recherche Elf de Solaize, Chemin du Canal - 69360 Solaize, France D. ANDERSON The Pennsylvania State University, University Park, PA 16801, USA A methodology for studying the relationships between fracture behavior and morphology of polymer-modified asphalts used as binders was developed by using the linear elastic fracture mechanics (LEFM) method and confocal laser scanning and environmental and cryo-scanning electron microscopies. Different types of polymers were used as modifiers: (i) copolymers from ethylene and methyl acrylate (EMA), butyl acrylate (EBA) or, vinyl acetate (EVA); (ii) diblock or star-shape triblock styrene-butadiene copolymers (SB or SBS * ). The 4 to 6 wt. % blends display an heterogeneous structure with a polymer-rich dispersed phase based on the initial polymer swollen by the aromatic fractions of the asphalt. The fracture toughness of the blends is higher than for the neat asphalt even if KIc of blends remains low compared to usual polymer blends due to the brittleness of the asphalt matrix. The fracture behavior which is strongly dependent on the nature of the polymer is discussed from the toughening mechanisms given for the filled polymers and the polymer blends. The EBA, SB, and SBS-based blends compared to the EMA and EVA-based ones display a higher KIc due to the elastomeric behavior of the polymer phase leading to a more efficient energy dissipation during crack propagation. The sample prepared with 4% crosslinked SB (Styrelf) and the corresponding physical blend (non-crosslinked) display the better fracture properties. C 2001 Kluwer Academic Publishers 1. Introduction Asphalt is a viscoelastic material at room temperature, i.e. it behaves as a viscous fluid at high temperature whereas it is a brittle solid at low temperature. One of the most promising methods for improving the asphalt performances at low temperature is by using additives such as polymers. Several types of polymers have been proposed and used as asphalt modifiers with additive contents as low as 3–6 wt. %, including elastomeric copolymers such as as styrene-butadiene rubbers, SBR [1], poly(styrene-b-butadiene), SB, and poly(styrene- b-butadiene-b-styrene), SBS [1–3]. The final volume fraction of the polymer-dispersed phase is higher that the initial one, about 20% by volume due to the swelling with the aromatic oily species of the original asphalt. The resulting blends, also called as physical blends, are generally unstable mixtures and macrophase separation can occur during long term storage times and/or at high temperatures. To overcome this demixing phenomenon, a dynamic vulcanization process can be done consist- ing in the in-situ crosslinking of the polymer dispersed phase [4]. In the United States, the Government passed legislation to add waste rubber like crumb rubber to their asphalt pavements [5, 6]. Others have modified asphalt with polyolefins such as polyethylene [3, 7, 8], ethylene vinyl acetate, EVA [8, 9], or EMA [10, 11]. The use of crumb rubber or recycled polyolefins is en- vironmentally attractive because it offers alternative for recycling the plastics wastes. The final properties of the asphalt-polymer blends depend on the morphology, i.e. the distribution of par- ticle sizes and composition of the phases, but also on the interface between the asphalt continuous phase and the polymer-rich dispersed phase. From previous stud- ies on styrene-butadiene block copolymer-modified asphalts [12], it was demonstrated that blends dis- plays an emulsion-like morphology and the polymer is swollen by some fractions of the asphalt, mainly by the maltenes. As a consequence, the interfacial tension between the swollen polymer and the asphalt matrix 0022–2461 C 2001 Kluwer Academic Publishers 451
Transcript
JOURNAL OF MATERIALS SCIENCE36 (2001 ) 451– 460
Low temperature fracture properties
of polymer-modified asphalts relationships
with the morphology
L. CHAMPION, J.-F. GERARD∗Laboratoire des Materiaux Macromoleculaires, UMR CNRS 5627, INSA Lyon Bat. 403, 69621Villeurbanne Cedex, FranceE-mail: [email protected]
J.-P. PLANCHE, D. MARTINCentre de Recherche Elf de Solaize, Chemin du Canal - 69360 Solaize, France
D. ANDERSONThe Pennsylvania State University, University Park, PA 16801, USA
1. IntroductionAsphalt is a viscoelastic material at room temperature,i.e. it behaves as a viscous fluid at high temperaturewhereas it is a brittle solid at low temperature. One ofthe most promising methods for improving the asphaltperformances at low temperature is by using additivessuch as polymers. Several types of polymers have beenproposed and used as asphalt modifiers with additivecontents as low as 3–6 wt. %, including elastomericcopolymers such as as styrene-butadiene rubbers, SBR[1], poly(styrene-b-butadiene), SB, and poly(styrene-b-butadiene-b-styrene), SBS [1–3]. The final volumefraction of the polymer-dispersed phase is higher thatthe initial one, about 20% by volume due to the swellingwith the aromatic oily species of the original asphalt.The resulting blends, also called as physical blends, aregenerally unstable mixtures and macrophase separationcan occur during long term storage times and/or at hightemperatures. To overcome this demixing phenomenon,a dynamic vulcanization process can be done consist-
ing in thein-situcrosslinking of the polymer dispersedphase [4]. In the United States, the Government passedlegislation to add waste rubber like crumb rubber totheir asphalt pavements [5, 6]. Others have modifiedasphalt with polyolefins such as polyethylene [3, 7, 8],ethylene vinyl acetate, EVA [8, 9], or EMA [10, 11].The use of crumb rubber or recycled polyolefins is en-vironmentally attractive because it offers alternative forrecycling the plastics wastes.
The final properties of the asphalt-polymer blendsdepend on the morphology, i.e. the distribution of par-ticle sizes and composition of the phases, but also onthe interface between the asphalt continuous phase andthe polymer-rich dispersed phase. From previous stud-ies on styrene-butadiene block copolymer-modifiedasphalts [12], it was demonstrated that blends dis-plays an emulsion-like morphology and the polymeris swollen by some fractions of the asphalt, mainly bythe maltenes. As a consequence, the interfacial tensionbetween the swollen polymer and the asphalt matrix
calculated from the Palierne’ model is very low (about10−5 N ·m−1). In addition, from the swelling of thepolymer by the maltenes, the asphalt continuous phaseis artificially enriched in asphaltenes by a “physical dis-tillation” of the lighter species from the original asphalt,leading to a toughened matrix. Such an understandingof the physical modifications of both the polymer andthe asphalt can be used to select the microstructure, i.e.the chemical nature, the molar mass, etc., of the poly-mer which is used as modifier.
In fact, as asphalt is brittle at low temperature, thethermal cracking of asphalt pavements is a serious prob-lem in cold countries. Nevertheless, the current specifi-cations do not specifically consider the failure mecha-nisms. Hesp [13–15] developed a method based on thelinear fracture mechanics, LEFM, principles in orderto characterize the fracture behavior of neat asphaltsand polymer-modified asphalts at low temperatures.From his work, the fracture toughness of the polymer-modified asphalts at−20◦C are higher than the tough-ness of the neat asphalt. Sabbagh and Lesser [16] alsostudied the mechanical behavior of polyolefin-modifiedasphalts using a three-point-bending beam method. Forsuch materials, the low temperature fracture tests alsoshowed an increase for KIC with increasing the amountof polyethylene (from 0 to 5% by wt.).
The aim of this study was to establish relationshipsbetween the fracture properties, measured from themethod developed by Hesp, and the morphology ofpolymer-modified asphalts. The effects of the chemicalnature and the amount of polymer used for modifyingthe asphalt on the fracture toughness were also stud-ied. For such a purpose, original methods were usedto study the morphology of polymer-modified asphaltssuch as confocal laser scanning microscopy (CLSM).In addition, the environmental scanning electron mi-croscopy (ESEM) and the cryo-scanning electron mi-croscopy (CSEM) were used to examine the fracturesurfaces in order to explain the differences in fractureproperties for different modified binders.
2. Experimental2.1. Materials2.1.1. AsphaltThe neat asphalt used for the modified binders wasa 70/100 penetration grade (penetration at 25◦C: 851/10 mm; ring and ball softening point: 45.6◦C) ob-tained from a Elf-Antar refinery and denoted G0078.The glass transition temperature of the asphalt isabout –27◦C and the crystallized fractions content is5%. The generic composition based on the SARA(Saturates, Aromatics, Resins, Asphaltenes) fractionsof the asphalt is given in Table I. The asphalt canbe considered as a continuum and the SARA frac-tions are defined from the solubility in various sol-
TABLE I Composition of the G0078 asphalt from the SARA fractions
vents [17]. By definition, the asphaltenes precipitatein n-heptane whereas the maltenes are soluble in thissolvent. Coupling thin layer liquid chromatographyand flame ionization detection allows to distinguishmaltene species by using successive solvents such ascyclohexane (saturates), dichloromethane (aromatics),and a dichloromethane/methanol/isopropanol mixture(70 : 25 : 5) (resins). The asphaltenes, having molarmasses between about 800 and 3,500 g·mol−1 areformed of on numerous polycondensed aromatics anddangling aliphatic chains [18]. These species are asso-ciated to form graphitic stacks or micellar structuresin solvents [18] which can be evidenced by small-angle X-Ray scattering (diameter about 2–4 nm). Thesaturates contain few linear alkanes which can becrystallized and have very low molar masses (about600 g·mol−1). The Tg of saturates is about−70◦Cand these ones display a dissolution/crystallization phe-nomenon betweenTg and 100◦C. The aromatics repre-sent the largest fraction of the asphalt. These are basedon less aliphatic chains with slightly condensed aro-matic rings and display aTg at about−20◦C, i.e. closeto that of the whole asphalt [19, 20]. The resins, alsocalled naphteno-aromatics have a composition which isclose to that of the asphatenes with aTg at about 20◦C.
The asphalt can be modeled as a colloidal suspen-sion of asphaltenes peptized by the resins fractions[21, 22]. As a consequence, the structure changes as thetemperature decreases from a newtonian fluid at hightemperature (above 60◦C) to a structure for which thepeptization layers of resins are enough thick to reachpercolation [12].
2.1.2. Polymers used as modifiersTwo different types of copolymers, semi-crystalline andamorphous, were used for modifying the asphalts.
Semi-crystalline copolymers such as poly(ethylene-co-vinyl acetate), EVA, poly(ethylene-co-methyl acry-late), EMA, and poly(ethylene-co-butyl acrylate), EBAwere supplied by ELF Atochem Company. The charac-teristics of the ethylene-based copolymers are reportedin Table II. The weight fraction of the comonomerwas determined by RMN1H. The rate of crystallinitywas calculated from the measured melting enthalpy ofthe polyethylene-enriched phase in the copolymers andfrom the equilibrium melting enthalpy of a pure crystalof linear polyethylene (taken equal to 293 J·g−1 [23]).
The amorphous copolymer SB is a linear styrene-butadiene diblock copolymer whereas SBS∗1, SBS∗2 andSBS∗3 are star-shape triblock styrene-butadiene copoly-mers. The weight fractions of the polystyrene and
TABLE I I Physical properties of the semi-crystalline copolymers
Fraction of the CrystallinityCopolymer comonomer (% wt) Tg (◦C) rate (%)(a)
polybutadiene blocks and the molar masses were deter-mined by RMN1H and size exclusion chromatography,SEC (Table III).
2.2. Processing of the polymer-modifiedasphalts
The asphalt modification was produced by mixing theasphalt and 4 to 6 weight % polymer under moderateshear rate (300 rpm) at 180◦C for few hours. The mix-ture was poured in silicone molds at 180◦C having theshape of the specimens for fracture tests and the sam-ples were cooled down to room temperature at a coolingrate of 2 K·min−1.
As the morphologies of the blends were examined us-ing confocal laser scanning microscopy (CLSM), 0.1%of a fluorescent probe (Rhodamin-B C.I. from Aldrich)was added to the semi-crystalline polymers to stainthe copolymers and increase the contrast. As a con-sequence, the polymer-rich dispersed phase could beobserved in the asphalt-rich phase. The copolymer andthe Rhodamin-B were mixed in toluene at 70–80◦C.Then, the solvent was removed at room temperaturefor one week. Differential scanning calorimetry wasperformed in order to ensure that the toluene was com-pletely removed from the polymer. The stained polymerwas then added to the asphalt in the same way as de-scribed previously. The blends with stained polymerswere stored at−4◦C for one week before observationto prevent the diffusion of the Rhodamin-B from thepolymer-rich phase to the asphalt-rich phase.
2.3. Mode-I fracture testThe fracture test was carried out by using a three pointbending beam method based on the ASTM E399-83procedure. Samples with a V-shape pre-notch (angle90◦) were prepared using the method developed byHesp [13–15]. As mentioned previously, the binder tobe studied was reheated at 180◦C and poured in a sili-cone mold (25× 12.5× 175 mm3) having a 90◦ notchin its centre. The samples were stored at−20◦C for2 hours, removed from the mold, and kept for 18 hoursat the test temperature, i.e.−20◦C. The pre-notch wassharpened with a razor blade immediately before the
test and the new crack length, denoteda, was mea-sured under an optical microscope. The beams wereplaced in the environmental chamber for 10 minutesbefore mechanical testing. The crosshead speed was0.6 mm·min−1 and the sample was loaded until thefracture crack propagated. The critical stress intensityfactor, KIC, was calculated according to the followingequation [24] from tests conducted on 8 samples.
where Pf is the failure load andS the span fixed to100 mm.B andW are the sample depth (12.5 mm) andthe specimen width (25 mm), respectively. The cracklength,a, was measured for each sample.
2.4. Analyses of the morphologiesAn environmental scanning electron microscope(Elec-troscan Explorer 2010)and a cryo-scanning electronmicroscope (Philips XL40 FEG-SEM) were used toexamine the fracture surfaces. ESEM allows the exam-ination of surfaces of practically any specimen, wet ordry, insulating or conducting, by allowing the presenceof a gas in the specimen chamber [25, 26] but the res-olution is limited.
The observation of asphalts with a conventional SEMis difficult because of the high vacuum requirementsand their lowTg. The cryo-preparation and observationequipments for conventional SEM has been availablefor over two decades and applied to biomedical and ma-terials specimens [27]. This innovation allows SEM ob-servation of soft or liquid specimens at low-to-mediummagnification. The fractured samples are coated with4 nm of platinum at−165◦C and transferred to themicroscope.
The fractured specimens were stored for one to threedays at−25◦C before ESEM and CSEM experiments.The morphology was analyzed from the fracture sur-faces of the specimens observed at respectively−5◦Cand−165◦C by ESEM and CSEM.
Another way to analyze the morphology of theblendswas to use the confocal laser scanning microscopy(CLSM). Samples from each type of the polymer-modified asphalts were prepared by squeezing thepolymer-asphalt mixture between glass plates (100µm-thick specimens). A Carl Zeiss laser scan microscopewas used and the images were recorded in transmissionmode using the He-Ne laser (543 nm wavelength) orthe Ar laser (488 and 514 nm wavelengths). CLSM isa relatively new technique which can be used also toobtain 3-D images and has already been applied to theobservation of several types of materials such as poly-mer blends [28], fiber-based polymer composites [29],or porous silicon [30].
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3. Results and discussion3.1. Morphology of the polymer-modified
asphaltsThe confocal laser scanning microscopy allows to ob-serve the morphology of the polymer-asphalt blends bycapturing the contrast between the two phases (Fig. 1).As reported previously, whatever the polymer type,
Figure 1 CLSM photographs. (A): blend based on 4% wt. EBA-35 (λ= 543 nm). (B): blend based on 4% wt. star-shaped SBS∗1 (λ= 488 nm).
(C): blend based on 4% wt. star-shaped SBS∗2 (λ= 514 nm). (D): blend based on 4% wt. star-shaped SBS∗
3 (λ= 543 nm). (E): blend based on 4% wt.SBS1110 (λ= 543 nm). (F): blend based on 6% wt. EMA-28 (λ= 543 nm). (G): blend based on 6% wt. EVA-28 (λ= 543 nm). (H): blend based on6% wt. EBA-35 (λ= 488 and 514 nm).
the volume fraction of the polymer-rich phase in thepolymer-asphalt blends is very high compared to theinitial amount of polymer added to the asphalt [31].The high volume fraction, which cannot be determinedprecisely from the CLSM micrographs, is explainedby the swelling of the polymer by some fractions of
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the asphalt. For example, from the volume fraction ofthe polymer-rich dispersed phase for the 4 wt. % SB-modified asphalt, the swelling rate can be estimated to55%. This phenomenon is confirmed by the fact thatthe two-phase morphology can be evidenced by CLSMeven though no fluorescent probe, i.e. Rhodamine®, isadded to the polymer before mixing with the asphalt. Infact, the polymer is swollen by aromatics species fromthe asphalt, the fluorescence of which allow to revealthe polymer domains [31].
For the ethylene copolymers, i.e. EMA and EVA, andfor the styrene-butadiene diblock, SB, and star-shapedtriblock, SBS, based blends, the CLSM photographsdisplay polymer-rich particles dispersed in a continu-ous asphalt matrix with similar swelling rates of thepolymer as reported in the literature [1, 32]. The in-terface between the two phases is very sharp and therange of the particles size is respectively between 2 and25µm for EMA- and EVA-based blends and between10 and 50µm for SB and SBS-based blends. No sig-nificant differences are observed in between the twoblends based on the styrene-butadiene diblock copoly-mers which although differ from their molar masses.This phenomenon is in agreement with that reported inthe literature [1]. In fact, the molar mass of SB mod-ifiers is not an important parameter for changing themorphology, whereas the particle size is more depen-dent on the amount of polystyrene [1].
On the opposite, for the poly(ethylene-co-butyl acry-late), EBA, based blends, a co-continuous morphologyis evidenced and the interface between the polymer-richdomains and the asphalt-rich matrix is blurred.
3.2. Fracture toughnessThe fracture properties measured at−20◦C for allthe polymer-modified binders are reported in Ta-ble IV. The KIC value of the neat asphalt is equal to48± 9 kPa·m1/2 and is in the same order of magnitudeas the values reported previously by Hesp for other
TABLE IV KI C values for different polymer-modified asphalts mea-sured at−20◦C
Amount of KICMaterial polymer (% wt.) (kPa·m1/2)
Neat asphalt 0 48± 9Asphalt/EVA 18 4 39± 10
6 63± 15Asphalt/EVA 28 4 60± 13
6 74± 20Asphalt/EMA 28 4 47± 13
6 67± 11Asphalt/EBA 35 4 64± 10
6 126± 20Star-shaped SBS∗1 4 85± 19
(28% m. polystyrene)Star-shaped SBS∗2 4 107± 11
(28% m. polystyrene)Star-shaped SBS∗3 4 66± 15
(18% m. polystyrene)Diblock SB 4 111± 16
(12% m. polystyrene)Styrelf 4 113± 21
types of binders [13, 14]. This value is low comparedto those of the polymers which are in the order of mag-nitude of several MPa·m1/2. As a consequence, theasphalt displays the fracture behavior of a very brittlematerial.
Table IV shows that SBS∗ and SB-based blends dis-play higher fracture properties compared to those formixtures produced with polyethylene-based copoly-mers. A blend with 4% of SBS or SB exhibits the sameKIC value as the mixture with 6% of EBA. Neverthe-less, the compatibility between the polymer and theasphalt and the fracture mechanisms can be differentfor these two types of blends. The sample preparedwith 4% crosslinked SB and the corresponding phys-ical blend (non-crosslinked) display similar propertieswhereas their morphologies are probably not the same.
ESEM and CSEM were used to examine the fracturesurfaces in order to explain the differences in fractureproperties observed for the modified-asphalts. The en-vironmental scanning electron micrograph of the neatasphalt shows that the fracture surface is mirror-like (notopographic contrast). As shown earlier from the lowvalue of KIc, the neat binder is a very brittle material.
Three parameters need to be taken in account to ex-plain the fracture mechanisms of the polymer-modifiedasphalts as for common polymer blends: the typeof morphology (dispersed particles vs. co-continuousphases and the distribution of particle size), the vol-ume fraction of the dispersed phase, and the adhesionbetween the two phases. Numerous papers describedin the literature the fracture phenomena occurring asa crack propagates through polymer materials, but ap-parently no work has been done yet on the fracturemechanisms in polymer-modified asphalts. Neverthe-less, the fracture descriptions done for polymer basedmaterials can be of interest for our purpose. The crackfront pinning process, initially proposed by Lange [33]and then modified by Evans [34], supposed that thecrack can be slowed down or hindered by the presenceof particles acting as obstacles. From the creation ofadditional fracture surface, this phenomenon leads toa higher fracture energy but it supposes that the parti-cles are stiffer than the matrix. In the case where theparticles display a ductile behavior and the interfacialadhesion is high, a crack front bridging process is in-volved. The toughening of heterogeneous materials canbe achieved also from the deviation of the crack, fromparticle to particle (crack deflection mechanism). Fi-nally, a mechanism proposed by Kinloch [35], denotedcrack-tip blunting, was proposed to explain the non-stable crack propagation, i.e. the stick-slip mechanism.In the case of polymer-modified asphalts, the stress-strain curve recorded on SEN specimens during bend-ing displays a brittle behavior without a stick-slip prop-agation of the crack. In addition, the fracture behaviorneeds to be considered with respect to the difference be-tween the temperature of testing,−20◦C, and theTg’sof the components. For the asphalt matrix, the temper-ature of testing is located at the beginning of the glasstransition zone and as a consequence, the continuousphase displays a brittle behavior. This brittleness is alsofavored by the physical distillation phenomena which
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leads to the enrichment of the matrix in asphaltenes act-ing as rigid nanofillers. As a consequence, whatever thepolymeric modifier is, the fracture toughness remainsvery low (Table IV).
The difference between the temperature of testingandTg of the polymers depends on the type of polymer(see Tables II and III). In fact, the poly(ethylene-co-vinyl acetate) and poly(ethylene-co-methyl acrylate)are also in a glassy state at the temperature of testing.In fact, their glass transitions are very close to−20◦C,but on the opposite, the poly(ethylene-co-butyl acry-
Figure 2 Fracture surfaces (ESEM). (A): blend based on 4% wt. EVA-18. (B): blend based on 4% wt. EVA-28. (C): blend based on 4% wt. EMA-28.(D): blend based on 4% wt. EBA-35. (E): blend based on 4% wt. star-shaped SBS∗
1. (F): blend based on 4% wt. star-shaped SBS∗2. (G): blend based
on 4% wt. star-shaped SBS∗3. (H): blend based on 4% wt. SB. (I): blend based on 6% wt. EVA-18. (J): blend based on 6% wt. EVA-28. (K): blendbased on 6% wt. EMA-28. (L): blend based on 6% wt. EBA-35.
late), EBA, having aTg of −45◦C is in the rubberystate. The difference inTg for EMA and EBA, canexplain the slighly higher value of KIC for the EBA-modified asphalt (Table IV). In fact, the fracture pro-cess for the EVA or EMA-based binder involves parti-cle pull-out with no deformation in the asphalt matrix(Figs 2 and 3A). For these blends, the crack bypasses thepolymer domains and the fracture occurs at the interfacebetween the two phases. As a consequence, the fracturemechanism is mainly governed by the poor adhesion be-tween the polymer-rich domains and the asphalt matrix.
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Figure 2 (Continued.)
On the opposite, the EBA-based blends display moreplastic deformation on the fracture surface, but the inter-facial adhesion remains poor (Fig. 2D and L). As a con-sequence, the fracture toughness remains slighly higherthan for the EMA and EVA-blends. In addition, fromCLSM, it was demonstrated that the EBA-based blendsseem to display a fine co-continuous structure. Thus, theincrease of the KIC value obtained with EBA-modifiedasphalt compared to EVA- or EMA-modified ones canbe associated with the plastic deformation of the poly-mer phase which is favored by such morphology.
As expected for the polyolefin-based blends, the frac-ture toughness value is higher for the asphalt modifiedwith 6% polymer (Table IV). The improvement of KIC
can be explained by the increase in volume fraction ofthe dispersed phase when the polymer content increasesfrom 4 to 6%. These results cannot be compared eas-ily with other data from the literature as only few pa-pers reported KIc values for polymer-modified asphalts.For a 85–100 grade asphalt modified with 3 wt. % ofchlorinated polyethylene [15], Hesp reported a valueof 154.5 kPa·m1/2. Similar values were obtained bySabbagh and Lesser [13] for LDPE-modified asphalts.The data reported in the literature are in the same orderof magnitude than those reported in this work accordingto the different type of initial asphalts and morpholo-gies. Nevertheless, from the slight differences in thefracture toughnesses of the unstabilized and stabilized
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Figure 3 Fracture surfaces (CSEM). (A): blend based on 6% wt. EVA-28. (B): blend based on 6% wt. EVA-28. (C): blend based on 6% wt. SBS∗2.
(D): blend based on 6% wt. SBS∗2. (E): blend based on 6% wt. SBS∗2. (F): Styrelf.
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emulsion in polyolefin modified asphalts, Sabbagh andLesser concluded that the toughening mechanisms arenot sensitive to the morphology.
For the SBS or SB-based blends, the fracture prop-agation pattern seems to be different as the fracturesurfaces differ from the polyolefin-modified asphalts.For these polymer-asphalt blends, the polymer seems tobe stretched as the crack propagates through the poly-mer domains (Fig. 2C, G, and H). This mechanism,denoted as crack-bridging in polymer blends, requiresa good adhesion between the two phases and that thepolymer displays the behavior of an elastomer at thetemperature of testing. It was previously demonstratedfrom rheological measurements that the interfacial ten-sion between the polymer-rich dispersed droplets andthe continuous asphalt matrix is low [12]. As a con-sequence, one can assume that the hypothesis of ahigh interfacial adhesion is verified. In addition, suchblock copolymers are organized in microdomains anddisplay two glass transitions, corresponding to the PSand PB phases (Table III). In addition, it was demon-strated by transmission electron microscopy that themicrodomains organization of such block copolymersremains in the asphalt [31, 36]. Due to the high amountof polybutadiene in the copolymer, having a low glasstransition temperature, the polymer-rich phase displaysan elastomer-like behavior at the fracture test temper-ature which is in agreement with the crack-bridgingmechanism interpretation. Nevertheless, for the otherpolymer-asphalt blends, for example with star-shapedSBS∗2 modified asphalt, the fracture surfaces demon-strated that particle pull-out occurs (Fig. 2F) but thefracture toughness remains as for the blends based onthe SB diblock copolymers. The higher values of KICfor these blends can be attributed to the ability of thepolymer domains to be plastically deformed before be-ing pulled out (Fig. 3) and to the quite good adhesion atthe interface between the polymer phase and the asphaltbinder. No topographic contrast is observed on the frac-ture surface of the binder modified with crosslinked SBdue to the morphology which is too fine to be exam-ined using ESEM. On the other hand, the observationof the Styrelf by CSEM reveals that the polymer-richdomains are very small compared to those of the cor-responding physical blends (Fig. 3F). The high KICvalue of the chemical blend can be attributed to thisfine morphology.
The fracture surfaces are difficult to observe even us-ing the environmental scanning microscopy. In fact, thesurface can relaxe after that the fracture propagates andduring the time elapsed between the fracture test and theobservation in the microscope. By cryo-scanning elec-tron microscopy, the specimens are observed at lowtemperature. So, the electron beam damage and themodification of the fracture surface are limited. More-over, the resolution is higher in CSEM compared toESEM which allows the observation of the Styrelf andthe polymer-rich phase more accurately.
Tg’s of the asphalt phase and of most of the polymersare in the same range of temperature as the fracturetest one. As a consequence, the fracture mechanismscannot be described as for common polymer blends
where one of phase is in the glassy state. Nevertheless,the crack propagation resistance of polymer-modifiedblends seems related both on the morphology and onthe state of the polymer at the temperature of testing,i.e. in the glassy or rubbery state.
4. ConclusionsA mode-I fracture test on SEN specimens has beenconducted to measure the fracture properties at lowtemperature of neat and polymer-modified blends us-ing the linear elastic fracture mechanics (LEFM). TheKIC values need to be discussed as a function of i)the morphology of the blends, i.e. dispersed phasevs. co-continuous structure, volume fraction, distribu-tion of particle size, state of dispersion, and composi-tion of polymer-rich domains and asphalt matrix, andii) the interfacial adhesion. In fact, as reported previ-ously, the volume fraction of the polymer-rich phasein the blend is higher than the initial amount due tothe swelling of the polymer by the aromatic fractionsof the asphalt. On the other hand, from the physicaldistillation of the asphalt when mixed with a poly-mer, the asphalt matrix is enriched with asphalteneswhich can be considered as stiff nanofillers. This studyconducted on various asphalt/polymer blends basedon different types of polymers, polyolefins or styrene-butadiene block copolymers, demonstrates that the ad-dition of polymer to asphalt increases the fracturetoughness at low temperature. However, the improve-ment is higher with styrene-butadiene copolymers thanwith polyethylene-based copolymers due to the differ-ent toughening mechanisms involved during the crackpropagation. Using environmental and cryo-scanningelectron microscopy and taking into account the tough-ening mechanisms described for filled polymers andpolymer blends, it was shown that for mixtures withEVA and EMA, the crack propagates at the interfacebetween the polymer-rich phase and the asphalt-richmatrix which displays a brittle behavior whatever thepolymer is. This fracture behavior can be explained bythe fact that all the components of the blend are at thebeginning of their glass transition region at−20◦C, i.e.in the glassy state, and by the poor adhesion betweenphases. On the opposite, the EBA-based blends displaya higher plastic deformation induced in the vicinity ofthe polymer phase in the asphalt matrix, which can ex-plain the KIC increase compared to other polyethylene-based copolymers. This effect is related to the lowerTg of poly(ethylene-co-butyl acrylate) which is in arubbery state at the temperature of testing and by theco-continuous structure observed by CLSM. Such amorphology can contribute also to a higher fracturetoughness of the resulting blends. Concerning blendswith SB or SBS, the improvement can be explainedby a better adhesion between the phases due to a bet-ter compatibility between the polymer and the binderwhich governs the volume fraction and the propertiesof the polymer phase. In addition, as the microdomains-based structure of the SB or SBS copolymers remainin the blends, the increase of KIC can be explainedby the crack-bridging mechanism reported for polymer
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blends. In fact, due to the lowTg of the polybutadi-ene domains, the polymer-rich domains display an elas-tomeric behavior at−20◦C. Such a behavior is assumedin the crack-bridging mechanism.
This methodology involving the transposition offracture mechanics from polymer-based materials toasphalt-based ones and original microscopies, confocallaser scanning and environmental and cryo-scanningelectron microscopies, is helpful for designing the mi-crostructure of the polymer for an efficient reinforce-ment of the polymer-modified blends at low temper-ature. In fact, the compatibility with the asphalt, i.e.the morphology of the polymer-rich phase, and the in-terfacial adhesion can be defined in order to involveenhancement of the fracture toughness from efficientenergy consuming mechanisms.
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Received 20 September 1999and accepted 16 May 2000
