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ke T2 and TB candaries [27]. Thecipitates of the S

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ductile intergranular (with shallow dimples) and brittle intergra-

Contents lists available at ScienceDirect

.el

Materials Science

Materials Science & Engineering A 586 (2013) 418427boundaries. Other precipitates include S and T1 phases within theE-mail address: alexis.deschamps@grenoble-inp.fr (A. Deschamps).nular (featureless except some trace of intergranular precipitates)fracture is observed. In these alloys the common feature is thepresence of a high volume fraction of -Al3Li ordered precipitatesin the matrix, except in a precipitate-free zone close to the grain

0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved.http://dx.doi.org/10.1016/j.msea.2013.06.075

n Corresponding author. Tel.: +33 4 76 82 66 07; fax: +33 4 76 82 66 44.programs, for instance under the commercial name AIRWARE[12]. They attract currently a strong interest in the materials

observed [2835] in various microstructure conditions, from veryunderaged to later stages of ageing. Classically a mixture betweenthe competition between Al alloys and composite materials for air-craft structures, and durably high fuel prices, new alloy develop-ment has led to alloys with optimised compositions where most ofthe issues of former generation alloys have been overcome [47].With the right combination of copper and lithium contents a moreefcient precipitation strengthening can be reached, particularlywith the T1 phase that forms very efcient obstacles to dislocationmotion because of its thin platelet shape of high aspect ratio [811]. These alloys nd a number of applications in new airplane

(nominally Al2CuLi). However, depending on the aand processing conditions other minor phases libe precipitated, particularly at the grain bounpresence of Mg can lead to the formation of pre(Al2CuMg) sequence.

One of the key issues that has been recognisedevelopment of AlLiCu alloys is their propensityfracture. In alloys containing of the order of 1.8AA8090, AA2090, AA2091), intergranular decohestolerance or high costs of processing. However, in the context of phases, the one offering the highest strength being the T1 phaselloy composition1. Introduction

Among precipitation hardeningalloys possess a combination of proattractive for structural applicationsector. Developments conducted in twere based on the main factor that talloy density and increases the mohigh lithium content relative to coppapplications, due to issues related tterm ageing at relatively low tempernium alloys, AlCuLis that has made themcially in the aerospacehties and early ninetiesition of Li decreases the[1,2]. Early alloys withtent have found limitedility losses during long[3], insufcient damage

processing, microstructure development and various properties[1325].

The ternary AlCuLi system experiences a complex precipita-tion sequence, exhibiting aspects of both binary AlCu and AlLisystems (see [26] for a more complete description). The (standard)addition of minor alloying elements such as Mg and Ag providesadditional complexity. The binary sequence of AlCu leads tosingle atomic layer GPI zones as well as GPII and precipitates,whereas the AlLi sequence gives the (Al3Li) phase, the forma-tion of which is known to depend strongly on the Li content. Heattreating the ternary AlCuLi system can generate a number ofOn the role of microstructure in governof an aluminumcopperlithium alloy

B. Decreus a,b, A. Deschamps a,n, P. Donnadieu a, J.C.a SIMAP, INP Grenoble CNRS UJF, BP 75, 38402 St Martin d'Hres Cedex, Franceb Constellium CRV, Voreppe Research Center, BP 27, 38341 Voreppe Cedex, France

a r t i c l e i n f o

Article history:Received 26 April 2013Received in revised form24 May 2013Accepted 18 June 2013Available online 6 July 2013

Keywords:AlCuLiPrecipitationIntergranular fracture

a b s t r a c t

The inuence of precipitatAlCuLi alloy, AA2198. Thby changing the quench rattreatment time. Fracture tevidence clearly the mechasignicantly occurs in all cThis mechanism is mainlydetermine the value of tramainly controlled by the s

journal homepage: wwwg fracture behavior

rstrmb

icrostructure on fracture mechanisms is studied in a recently developedtra-granular and inter-granular microstructures are varied independentlyom the solution heat treatment, the amount of pre-stretching and the heathness is evaluated by short bar chevron tear tests that make possible toms of inter-granular fracture. It is shown that intergranular ductile fractureitions of heat treatment where substantial precipitation has taken place.trolled by the state of inter-granular precipitation and plays a major role toerse fracture toughness, while the strength and ductility of the alloy areof intra-granular precipitation.

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& Engineering A

grains and either phase [28] or Li and Cu containing phases(T2 and TB) [36] at the grain boundaries. Different causes havebeen invoked to explain this extensive occurrence of intergranularfracture and have received extensive attention particularly in the1990s, and have been reviewed in several papers ([29,30,32,34,35]and references therein): presence of intergranular precipitatesassociated to the presence of precipitate free zones (particularlyin terms of -Al3Li precipitation), planar slip (due to the collectiveshearing of precipitates), embrittlement of grain boundaries byliquid metal impurities, or a loss of grain boundary coherency dueto Li segregation. The rst mechanism was deemed more impor-tant in cases where alloys exhibit ductile intergranular fracture, onthe basis of the experimental evidence that toughness andoccurrence of intergranular fracture were well correlated to thegrain boundary precipitate microstructure [28]. Conversely, thelast hypothesis seems nowadays to be favoured in the cases wherebrittle intergranular fracture is observed [34,35].

However, most recently developed alloys such as AA2198 orAA2050 have a Li content limited to lower concentrations (of theorder of 1 wt%), in order to minimise or even suppress the

conditions has proven to be a very efcient way to identify clearlythe intergranular fracture mode.

Following a detailed study of the microstructure developmentinside the grains and related mechanical properties during heattreatment of the AA2198 alloy that has been published recently[26,38], the aim of the present paper is to evaluate the evolutionof fracture mechanisms and related fracture toughness in a largevariety of heat treatments where different parameters are variedacting on the microstructure at the grain boundaries (quench ratefrom the solution treatment temperature, subsequent ageing time)and on the intra-granular microstructure (pre-stretch and ageingtime). The mechanical properties are evaluated by simple tensiletests and by short transverse chevron notch (short bars) samples,giving access to short transverse fracture toughness and relatedfracture surfaces. The microstructure at the grain boundariesis characterised by electron microscopy and the intra-granularmicrostructure has been characterised in the formerly publishedwork [26].

BW2H in Fig. 2. The tests were carried out at a traverse speedof 0.5 mmmin1 with a pre-load of 5 N, according to ASTME1304-97 [39]. The conditional plane strain toughness K was

B. Decreus et al. / Materials Science & Engineering A 586 (2013) 418427 419formation of the phase. These alloys are completely free of and their intragranular microstructure in the T8 aged condition isdominated by the T1 phase and, to a lesser extent, the phase[26]. In such alloys, pronounced intergranuar fracture is stillobserved in the aged T8 temper [23,24]. Similarly, other alloyswith relatively low Li concentration such as AA2020, which areexempt of precipitation have been also shown to be subject tointergranular fracture and delamination [37]. However, the rela-tionship between the microstructure of these low-Li containingalloys and their fracture mode has not been studied in great detail.The limited number of studies concerning the fracture behaviourof these recent alloys [23,24] have reported mechanical testing inthe rolling plane and no detailed microstructural analysis. Slantedfracture is promoted by the intergranular fracture mode, and abrous fracture surface is observed, related to the unrecrystallizedgrain structure, and related very high grain aspect ratio. The shearcomponent of deformation makes it difcult to determine pre-cisely the microstructural features linked with the intergranularfracture mechanism. Although it has been proposed that reducingthe Li content may be an efcient way to inhibit Li segregation tothe grain boundaries [35], it is not clear if these alloys are stillsubjected to this mechanism. In the high Li containing alloys, mostauthors have used short transverse testing in order to evidence themechanisms of intergranular fracture [3134]. Using such testingFig. 1. Optical micrograph of the grain structure.Qv2. Material and experimental methods

Alloy AA2198 has the composition range (all in wt%) [2.93.5]Cu[0.81.1]Li[0.250.8]Mg[0.10.5]Ag[0.040.18]Zr. It wasprovided by Constellium Voreppe Research Centre, France, asrolled sheet of 12 mm thickness with a fully brous grain struc-ture. The sheets showed a strong Brass texture (112 {110}). Fig. 1shows an optical micrograph of the grain microstructure.

The heat treatments consisted rst in a solution treatment anda quench. Unless stated otherwise, the samples were quenchedinto cold water. Some samples were also quenched into hot water(80 1C) or air cooled, in order to vary the grain boundary micro-structure prior to the articial ageing treatment. The samples werethen stretched at least 2% plastic strain and kept several weeks atambient temperature (T351 temper). The articial ageing treat-ment consisted in a heating ramp to 155 1C at 20 K h1, followedby an isothermal hold at 155 1C.

Tensile tests were carried out in the rolling (L) direction of theplates using at tensile tests of section 53 mm2 and gaugelength 60 mm, at a constant strain rate of 1.5104 s1.

Short bar chevron tear tests were carried out with a loadingdirection in the short transverse (ST) direction of the plates withsample dimensions (in mm) of 1912.511 corresponding toFig. 2. Geometry of the short bar chevron notched sample.

calculated as follows:

KQvM YnmFmB

Wp

where Q stands for the conditional nature of the measuredtoughness, and M states for the determination from the maximumload. Ynm is the tabulated minimum stress intensity factor (depend-ing on the sample geometry), and Fm is the maximum loadmeasured during the test. Some conditions must be met tocalculate this toughness value [39]. All the results presented here,except for the observation of fracture surfaces, were obtained invalid conditions; other results are omitted.

Fracture surfaces were observed using a standard scanningelectron microscope LEO Stereoscan 440 at 20 kV. High resolutionimages were obtained on a Zeiss Ultra-55 FEG-SEM at 4 kV within-lens secondary and back scattered detectors.

Conventional transmission electron micrographs were obtainedon electropolished samples with a Jeol 3010 microscope operatingat 300 kV. Complementary observations for the grain boundarymicrostructure were performed in scanning transmission electronmicroscopy (STEM) mode on a FEI Titan 80300 equipped a highangle annular dark eld (HAADF) detector.

3. Tensile tests

Fig. 3 shows the tensile curves obtained for samples heattreated from 1 h to 500 h at 155 1C. The microstructure evolutionduring this ageing treatment has been described in detail in [26].

0

100

200

300

400

500

600

0 2 4 6 8 10 12 14 16

True

stre

ss (M

Pa)

True strain (%)

1h

4h8h

16h100h

500h

Fig. 3. Tensile curves as a function of ageing time at 155 1C.

esol

B. Decreus et al. / Materials Science & Engineering A 586 (2013) 418427420Fig. 4. Fracture surfaces of tensile test samples (SEM micrographs) at low and high r

155 1C.ution for two different states of ageing: (a) and (b) 4 h at 155 1C; (c) and (d) 16 h at

-1/2

), 4*

U.A

. (%

)

B. Decreus et al. / Materials Science & Engineering A 586 (2013) 418427 421300

400

500

d (N

)

1h

8h

16h

100h

4hAfter 1 or 4 h at 155 1C, the microstructure consists mainly of ahomogeneous solid solution, and no signicant fraction of thestrengthening T1 precipitates is present. After 8 h at 155 1C, theformation of the T1 phase is really signicant and the microstruc-ture between 16 and 500 h is mostly stable, consisting on afraction close to equilibrium of very ne T1 platelets resulting inan almost constant yield strength. Fig. 4 shows the fracturesurfaces of the tensile specimens tested in two ageing conditions,namely in the absence of T1 precipitates (4 h (a, b)) and near peak

0

100

200Loa

Displacement (mm)

500h

0 0.5 1 1.5 2

Kqv

M (M

Pa.

m

Fig. 5. (a) Loaddisplacement curves for the short bar tests on samples aged for differenwith corresponding yield stress and uniform elongation of tensile tests in the same age

Fig. 6. Low resolution fracture surfaces (SEM micrographs) of the notch tip of the short bAn increase of the fraction of at intergranular fracture is observed when the ageing ti40

50

60

70

300

400

500

600

Yield stress

Yield stress

Toughnessstrength (16 h (c,d)). In the rst case, fracture appears to becompletely ductile trans-granular. The dimples are oriented dueto the slanted geometry of the fracture surface. In the second case,the fracture surface appears as multiple ne steps of a few mm,with smooth walls between them. These fractographs resemblethat obtained in similar conditions by Chen et al. [23] and Steglichet al. [24] and are characteristic of a fracture controlled byintergranular decohesion with a heavily elongated grain structurein the loading direction. Due to shearing and friction during the

0

10

20

30

0

100

200

1 10 100 1000

(MP

a)

Time at 155C

Elongation

t times at 155 1C. (b) Values of toughness measured from the short bar tests alonging conditions.

ar tests for the following ageing times at 155 1C: (a) 1 h; (b) 8 h; (c) 16 h; (d) 100 h.me proceeds.

B. Decreus et al. / Materials Science & Engineering A 586 (2013) 418427422last stages of the fracture process, it is impossible to describe themicroscopic mechanisms that explain this occurrence of intergra-nular fracture from these observations. Short bar chevron tests

Fig. 7. Detail of ductile fracture area of the short bar sample aged 1 h at 155 1C: (a) in seelectron mode showing the presence of large intermetallics associated with the dimple

Fig. 8. Detail of ductile fracture area of the short bar sample aged 100 h at 155 1C: (aintermetallics on the intergranular surface; (b) medium resolution secondary electrointergranular fracture surface; (c) high resolution secondary electrons image showing tpresented in the next paragraph allow for loading in the directionnormal to the grain boundaries and observing the details of thefracture mechanisms.

condary electron mode showing the presence of dimples and (b) in back-scattereds.

) low resolution image in back-scattered electrons mode showing no large scalens image showing the presence of a high density of small dimples on the athe presence of small particles in these dimples.

B. Decreus et al. / Materials Science & Engineering A 586 (2013) 418427 4234. Chevron notch short bar tests

Fig. 5(a) shows the loaddisplacement curves for the chevronnotch short bar tests after different heat treatment times. For thesamples heat treated 1 h and 4 h at 155 1C where the material isstill very soft, generalised plasticity prevented the measurement offracture toughness from these tests. However, it appears clearlythat the behaviour changes dramatically between 4 and 8 h, thenbetween 8 h and 16 h, and again between 16 h and 100 h, afterwhich it remains stable. Fig. 5(b) shows the evolution with heattreatment time of three parameters of the mechanical tests,namely the yield strength and uniform elongation obtained fromthe tensile tests, and the fracture toughness obtained from theshort bar tests. As evidenced in [26], the evolution of yield stress isessentially representative of the evolution of the volume fractionof T1 precipitates inside the grains. When T1 precipitates rstappear the yield stress starts to increase, and when this fractionsaturates, the yield stress stabilises around 500 MPa. The evolutionof elongation is quite well correlated to this evolution of yieldstress, as discussed in detail in [38]. The conjunction of theincrease in yield strength and the corresponding reduction of thestrain hardening rate capability results in a reduction in uniformelongation. Interestingly, the evolution of fracture toughness isquite different from that of the uniform elongation. The toughnessalso decreases with heat treatment time, but this decrease occurslater. In the rst stages of the increase of yield strength, thestrength increase compensates the reduction in elongation to keepthe toughness roughly constant. It is only after 50 h at 155 1C that

Fig. 9. Transmission electron micrographs showing the microstructure at the grain and susample aged for 16 h at 155 1C; (b) dark-eld micrograph showing a grain boundary inboundaries in the sample aged for 100 h at 155 1C.the toughness decreases signicantly, while the yield strength andelongation (albeit tested with a different loading direction) arenow constant with ageing time.

In order to understand the origin of this toughness drop duringthe peak strength plateau, fractographic observation of the shortbar samples will now be presented.

Fig. 6 presents low resolution images of the fracture surfacesclose to the notch tip, for different heat treatment times (1 h, beforethe increase in strength; 8 h, during the strength increase; 16 hclose to peak strength before the drop in toughness; and 100 h atpeak strength after the drop in toughness). In the rst sample thefracture is entirely ductile transgranular. No at areas are recorded.After 8 h and 16 h at 155 1C, a large majority of the fracture surfaceis still ductile transgranular; however a few at and smooth (at thisscale) surfaces can be observed that correspond to interganularfracture. After 100 h at 155 1C, most of the fracture surface consistsof at and smooth intergranular zones. It is therefore tempting toassociate the decrease in fracture toughness to the occurrence of afracture mode of low energy dissipation [40], namely inter-granulardecohesion along grains that have a large extension normal to theshort transverse loading direction.

The details of the fracture surface are shown in Fig. 7 for thesample aged 1 h at 155 1C, and in Fig. 8 for the sample aged 100 hat 155 1C. In the rst case, the ductile transgranular fracture isclassically characterised by large dimples, which include brokenintermetallic particles. These particles have a bright contrast(Fig. 7b) indicating a larger average atomic number than thematrix that is conrmed by the composition close to Al7Cu2Fe

b-grain boundaries. (a) Dark-eld micrograph showing a sub-grain boundary in thethe sample aged for 16 h at 155 1C; (c) and (d) STEM-HAADF micrographs of grain

ing treatment, the intragranular and intergranular microstructures

rt bf qu

B. Decreus et al. / Materials Science & Engineering A 586 (2013) 418427424given by EDS analysis. In the second case, the situation is entirelydifferent. The at areas of the fracture surface do not present anychemical contrast at low resolution (Fig. 8(a)), which means thatno coarse intermetallic particle is associated with this fracturemechanism. However, at higher resolution a homogeneous dis-tribution of very ne dimples is observed (Fig. 8(b) and (c)). Theirsize is between 100 and 200 nm, and very small particles (lessthan 50 nm in diameter) are observed at the core of the dimples.A quantitative chemical analysis of these particles is clearly out ofthe range of EDS-SEM, but a qualitative comparison between theEDS signal on these particles and on the surrounding matrixclearly shows that they are enriched in Cu and not in Fe. It is notpossible using this technique to estimate a potential enrichment inLi. Although not shown here for space reason, it should bementioned that similar observations were made on all at areasof short bar samples, with a general tendency of decreasing thedimple size with increasing the ageing time [41].

These fractographic observations provide good evidence thatthe intergranular fracture mode responsible for the at areas onthe short bar test fracture surfaces is actually ductile, with adamage initiation related to ne Cu-rich particles lying on the

30

35

40

45

50

55

60

Undeformed

T351

Kqv

M (M

Pa.

m-1

/2)

Microhardness

Def 16h

Undef 16h

Undef 100hDef 100h

Increasing ageing time

100 120 140 160 180 200

Fig. 10. (a) Evolution of the compromise between toughness (as measured by the shocompared to samples aged from an undeformed condition; (b) inuence of the rate otime at 155 1C.grain boundaries. These Cu-rich particles, less than 100 nm indiameter, are likely to be intergranular precipitates formed duringthe articial ageing treatment. In order to conrm this hypothesis,TEM observations are reported in Fig. 9. Fig. 9(a) and (b) showsconventional dark eld micrograph of a sub-grain boundary and agrain boundary after 16 h at 155 1C (at the beginning of the peakstrength plateau). In the matrix, a high density of ne T1 plateletsis observed. At sub-grain boundaries, this density is even higher,but the precipitates have a similar shape and size. Since T1precipitates are known to nucleate on dislocations, sub-grainboundaries can be regarded as regions of particularly high densityof nucleation sites for the precipitates. At the grain boundaries,however (Fig. 9(b)), hardly any precipitation can be observed,which is consistent with the fact that fracture is still mostlytransgranular in the short bar tests. Fig. 9(c) and (d) showsSTEM-HAADF (Z-contrast) images at grain boundaries of thesample aged 100 h at 155 1C. Inside the grains the microstructureis qualitatively identical to that of the sample aged 16 h (it canactually be shown that the two intragranular microstructures arequantitatively almost identical [26]). However the grain bound-aries now include a high density of incoherent particles of sizeapproximately 50 nm. Their bright contrast in the HAADF micro-graphs tells that they must be Cu-rich, in agreement with [36] andevolved concurrently, so that unravelling their respective effectsremains challenging.

In order to provide additional evidence, we have sought tochange independently the intergranular and intragranular micro-structures in two ways. First, we have compared samples aged inthe classical way (namely quenched, stretched and aged) towith similar observations in [13]. These precipitates may be T2 orTB phases; however their crystal structure was not determinedhere. Their size and morphology suit particularly well with that ofthe particles observed at the core of the dimples on the fracturesurfaces (Fig. 8(c)).

5. Effect of the variation of intergranular and intragranularmicrostructures

The preceding experiments show that the evolution of trans-verse toughness is correlated with the state of intergranularprecipitation, which in turn controls the occurrence of ductileintergranular fracture. However, during the studied articial age-

0

10

20

30

40

50

60

70

0 20 40 60 80 100 120

Kqv

M (M

Pa.

m-1

/2)

Time at 155C (h)

Fast quench

Intermediate quench

Slow quench

ar test) and microhardness in the sample aged at 155 1C from the T351 condition asenching on the toughness (as measured by the short bar test) evolution with ageingsamples where the stretching step was omitted. As shown in[42], intragranular precipitation in the absence of pre-deformationis very sluggish and the precipitation kinetics is decreased by afactor of 10. However it is likely that the precipitation kinetics atthe grain boundaries is at rst order unaffected by the absence ofstretching and thus it becomes possible to compare samples withdifferent intra-granular precipitates and similar inter-granularmicrostructure. Secondly, we have compared samples aged bythe same procedure (including the stretch), but having experi-enced different quenching conditions from the solution treatment(intermediate quench, in hot water and slow quench, air cooled).These samples can thus be expected to have similar intragranularprecipitates but different inter-granular microstructures, with ahigher fraction of intergranular precipitates in the slowly cooledmaterials.

Fig. 10(a) shows a graph representing the different sets ofmicrohardness/toughness parameters for samples aged at 155 1Cin the deformed and undeformed materials. In both cases aninverse correlation exists, however very different in the twosamples. As seen in the former sections, for the stretchedsamples, at near peak strength (16 h) the toughness is still veryhigh. At this ageing time, the unstretched material shows acomparable toughness value; however the microhardness is

B. Decreus et al. / Materials Science & Engineering A 586 (2013) 418427 425much lower due to the sluggish intragranular precipitationkinetics. After 100 h at 155 1C, the unstretched material hasreached the hardness value of the stretched material aged 16 h,but has a much lower toughness. Its toughness is now similar tothat of the stretched material aged 100 h at 155 1C. Therefore allthese results are consistent with a hardness controlled by thestate of intragranular precipitation and a toughness mostlycontrolled by the inter-granular precipitation (which is at rstorder controlled by the ageing time).

Fig. 10(b) shows the evolution with heat treatment time oftoughness for samples quenched with different procedures (fast,intermediate and slow). Clearly, decreasing the quench ratereduces very signicantly the toughness value in all ageingconditions. The remaining evolution is more complex, with aconvergence of the intermediate and fast quench samples at longageing time, while the slowly quenched material keeps a lowertoughness at any ageing time. The fracture surfaces of the shortbar tests after 100 h at 155 1C for the intermediate or slowquenched materials are shown in Fig. 11. They present a highfraction of at intergranular areas, with a high density of smalldimples (albeit larger than that of the fast quenched material,presumably due to a larger size of intergranular precipitates), andsome areas of facetted brittle fracture, which could correspond tothe formation at the grain boundaries of a continuous layer of abrittle incoherent intermetallic phase after a sufciently slowquench.

Fig. 11. Fracture surfaces (SEM micrographs) of short bar samples on slowly quenched m6. Discussion

As summarised in the Section 1, the fracture behaviour of AlLiCu alloys containing of the order of 2 wt% Li has been exten-sively investigated in the literature. These alloys have been shownto be prone to intergranular fracture, often of brittle nature. Thisbehaviour was shown to be strongly related to the presence of Liby several mechanisms that depend on the alloy composition (andrelated precipitates formed) and heat treatment, namely Li segre-gation, grain boundary precipitates, zones free of precipitates atthe grain boundaries, and planar slip due to the ordered phases.In the more recently developed alloys where the Li content islimited to about 1 wt% in order to suppress the formation of tothe benet of the T1 phase, the literature on the fracture mechan-isms is still incomplete, although intergranular fracture has alsobeen observed in such alloys to play a prominent role. Themicrostructural observations made in the present paper and inthe two companion papers published earlier on the same alloy[26,38], as well as the fractographic observations reported here,show that these alloys have a very different microstructure ascompared to their higher Li counterparts. This is especiallyimportant with respect to the mechanisms that have been invokedto explain the fracture behaviour. First, they do not show any precipitation, nor any precipitate-free zone (PFZ). After ageing, theT1 platelets have an extension of more than 50 nm, which helpsto explain the absence of PFZ because any precipitate nucleated

aterials: (a) intermediate quench, 100 h ageing; (b)(d) slow quench, 100 h ageing.

grains, so that we believe that it is related to the decohesion along

[12] T. Warner, J.C. Ehrstrm, B. Chenal, F. Eberl, Light Met. Age 67 (2009) 3.[13] N. Brodusch, M. Trudeau, P. Michaud, L. Rodrigue, J. Boselli, R. Gauvin, Microsc.

B. Decreus et al. / Materials Science & Engineering A 586 (2013) 418427426the interface of large facetted precipitates. In all other cases themacroscopically at intergranular fracture surfaces included a highdensity of very small dimples (size of about 100 nm), associatedwith precipitates whose size and composition were consistentwith that observed at the grain boundaries in the transmissionelectron microscope, namely Cu-rich, of average atomic numberhigher than the matrix as proven by the higher brightness in Z-contrast.

Therefore, it can be expected that the mechanisms inducingintergranular fracture may be of a different nature in the high-Liand low-Li containing alloys. Although we cannot rule out com-pletely a loss of the grain boundary cohesion due to the presenceof Li in the alloy, a good indication that such decohesion does notplay a major role in the low-Li containing alloy investigated here isthat we have not observed brittle fracture. In addition the values oftoughness are quite high for transverse toughness tests in pre-cipitation hardening aluminium alloys, which is a further indica-tion that grain boundary cohesion is not affected at rst order.Ductile grain boundary fracture is usually related to three para-meters [28,43]: the presence of grain boundary precipitates, a softprecipitate free zone, and a high stress triaxiality. In the presentcase no signicant precipitate free zone is visible. However thecompetition between intergranular and transgranular fractureappears to be affected by the trixiality ratio, since grain boundaryfracture is systematically observed in the areas of highest triaxi-ality of the chevron notch specimen (close to the initial notch).This effect of external triaxiality may actually help the acceleratedvoid growth in such conditions where no PFZ is present.

The evolution of fracture mechanism as well as the evolution oftoughness with the different microstructural states evaluated hereare consistent with the proposed mechanism controlling inter-granular fracture, namely nucleation and growth of voids at thegrain boundary precipitates. The toughness drops simultaneouslyto the increase in the fraction of intergranular fracture whenageing proceeds, and this happens independently (at rst order) ofthe intragranular precipitation kinetics, which was varied byplaying on the plastic deformation prior to the heat treatment.Similarly, a slower quench rate from the solution treatment, whichfavours a high fraction of grain boundary precipitates, results in alower fracture toughness.

7. Conclusions

In the present work we have attempted to shed light on thefracture mechanisms that prevail in the recently developed AlCuLi alloys with moderate Li content where the microstructure afterarticial ageing is dominated by the precipitation of the T1 phaseand where no phase is formed. In order to study the microscopicmechanisms related to intergranular fracture, which is frequentlymet in these alloys, testing has been performed in the shorttransverse direction. By studying the effect of heat treatment time,pre-stretch before ageing, and quench rate from the solutionat a distance smaller than 25 nm from a grain boundary canextend to the grain boundary provided that some Cu and Lisupersaturation still remains. Secondly, although they are proneto planar slip in the naturally aged condition due to the presenceof solute clusters, when aged, their slip becomes quite homoge-neous [38], because of the particular shearing mechanism of the T1precipitate, which is very different from that of . Thirdly, noevidence of brittle intergranular fracture has been found in thepresent study for all the heat treatments investigated, except forthe slowly quenched material. However in the latter case, thisbrittle fracture was facetted and did not follow the shape of thetreatment, we have been able to evaluate independently the effectMicroanal. 18 (2012) 1393.[14] N. Ward, A. Tran, A. Abad, E.W. Lee, M. Hahn, E. Fordan, O.S. Es-Said, J. Mater.

Eng. Perform. 20 (2011) 989.[15] R.J. McDonald, A.J. Beaudoin, Model. Simul. Mater. Sci. Eng. 18 (2010) 065007.[16] S. Kalyanam, A.J. Beaudoin, R.H. Dodds Jr., F. Barlat, Eng. Fract. Mech. 76 (2009)

2174.[17] I. Eberl, C. Hantrais, J.-C. Ehrtsrom, C. Nardin, Sci. Technol. Weld. Join. 15 (2010)

699.[18] S. Richard, C. Sarrazin-Baudoux, J. Petit, Proceedings of the TMS 2009,

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1763.[20] C. Bitondo, U. Prisco, A. Squilace, P. Buonadonna, G. Dionoro, Int. J. Adv. Manuf.

Technol. 53 (2011) 505.[21] P. Cavaliere, J. Light Met. 2 (2002) 247.[22] P. Cavaliere, M. Cabibbo, F. Panella, A. Squillace, Mater. Des. 30 (2009) 3622.[23] J. Chen, Y. Madi, T.F. Morgeneyer, J. Besson, Comput. Mater. Sci. 50 (2011) 1365.[24] D. Steglich, H. Wafai, J. Besson, Eng. Fract. Mech. 77 (2010) 3501.[25] Y.E. Ma, Z. Zhao, B. Liu, W. Li, Mater. Sci. Eng. A-Struct. Mater. Prop.

Microstruct. Process 569 (2013) 41.[26] B. Decreus, A. Deschamps, F. De Geuser, P. Donnadieu, C. Sigli, M. Weyland,

Acta Mater. 61 (2013) 2207.[27] S.C. Wang, M.J. Starink, Int. Mater. Rev. 50 (2005) 193.[28] A.K. Vasudvan, R.D. Doherty, Acta Metall. 35 (1987) 1193.of intragranular and intergranular microstructure on the strength,the uniform elongation in tension, and transverse fracture tough-ness as well as related fractography.

While the strength and uniform elongation are mainly con-trolled by the state of precipitation within the grains, we havestrong indications that the transverse fracture toughness is largelydetermined by the occurrence of intergranular fracture, which iscontrolled by the presence of incoherent, Cu rich grain boundaryprecipitates. These precipitates serve as damage nucleation sitesand the resulting fracture surfaces present a high density of sub-micrometric dimples. In practice, the choice of a heat treatmentat the very beginning of the peak strength plateau, where theintergranular precipitation microstructure is not yet fully devel-oped, makes possible to avoid a massive appearance of intergra-nular fracture and thus keep a high toughness value. Particularcare must yet be taken to ensure a sufciently fast quench ratefrom the solution treatment so that the grain boundary micro-structure prior to articial ageing is free of precipitates.

Acknowledgements

C. Sigli is thanked for fruitful discussions. L. Charpenay isthanked for helping with the short bar tests. The French researchagency (ANR) is thanked for nancial support under the projectALICANTDE. We are grateful to the Canadian Centre for ElectronMicroscopy, a facility funded by the Canada Foundation forInnovation and the Ontario Government where the HAADF STEMwork was carried out.

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B. Decreus et al. / Materials Science & Engineering A 586 (2013) 418427 427

On the role of microstructure in governing fracture behavior of an aluminumcopperlithium alloyIntroductionMaterial and experimental methodsTensile testsChevron notch short bar testsEffect of the variation of intergranular and intragranular microstructuresDiscussionConclusionsAcknowledgementsReferences