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Materials Science and Engineering A 437 (2006) 406–412

Determination of mechanical properties of aluminiummatrix composites constituents

J. Rodrıguez, M.A. Garrido-Maneiro, P. Poza ∗, M.T. Gomez-del RıoDepartamento de Ciencia e Ingenierıa de los Materiales, Universidad Rey Juan Carlos, Escuela Superior de Ciencias

Experimentales y Tecnologıa, C/Tulipan s.n. 28933 Mostoles, Madrid, Spain

Received 2 February 2006; received in revised form 26 July 2006; accepted 28 July 2006

bstract

This work presents nanoindentation experiments with Berkovich tips, carried out on an Al–Li 8090 alloy reinforced with 15 vol.% SiC particleso determine hardness and Young’s modulus of the metal matrix and the ceramic reinforcements. The influence of pile-up effect has been estimatederforming a finite element model of the nanoindentation tests to provide values representative of the materials’ behaviour. Results indicate a clear

ncrease in matrix hardness in comparison with the unreinforced alloy and a slow decrease with the distance to the reinforcement particles. Pilingp of the material tested overestimates the results obtained through nanoindentation. This effect could be corrected by means of the simulated testnalysis.

2006 Elsevier B.V. All rights reserved.

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(wtsoo[fiotoddma

c

eywords: Metal matrix composites; Micromechanical characterization; Nano

. Introduction

The reinforcement of metallic alloys (mainly aluminium andagnesium) with hard ceramic particles (SiC or Al2O3) has

een extensively studied. The addition of ceramic reinforce-ents improves the stiffness as well as the wear and creep

esistance, and to a minor extent the strength [1–3]. Furthermprovements in the elastic properties could be attained by usingl/Li alloys as matrices in these composites. The elastic mod-lus of aluminium alloys is increased by approximately 6% forvery 1 wt.% of lithium and the density is reduced by 3%, whichmproves significantly the specific modulus of these materials4]. For instance Al/Li–SiC composites exhibit elastic modu-us over 100 GPa with relative densities around 2.6 leading topecific stiffness 50% above that of standard aluminium anditanium alloys [5,6]. In addition, the presence of lithium in theluminium matrix helps to strengthen these materials throughhe precipitation of the ordered �′ (Al3Li) phase coherent with

he matrix [7,8], and the Al–Li/SiC system presents an excellenterformance/price ratio for applications where the stiffness isritical, as it is in the aerospace industry.

∗ Corresponding author. Tel.: +34 914 887 179; fax: +34 914 888 150.E-mail address: pedro.poza@urjc.es (P. Poza).

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921-5093/$ – see front matter © 2006 Elsevier B.V. All rights reserved.oi:10.1016/j.msea.2006.07.118

ation; Finite elements method

The mechanical properties of metal matrix compositesMMC) have been analysed through micromechanical modelshich can be broadly divided into three groups. Models based on

he modified shear-lag approach [9,10] provide simple expres-ions for the stresses acting on the reinforcements and for theverall composite behaviour, but they cannot cover the evolutionf stresses and strains during deformation. Mean field models11–14] use volume-averaged values for the stresses and strainselds in the two phases, matrix and reinforcement. Models basedn the finite element analysis of a unit cell, representative ofhe composite, have been developed from the initial predictionsf the stress–strain curve [15,16], which did not include anyamage mechanism, to more refined methods which includeamage mechanisms like reinforcement fracture [17–19] andodels based on multiparticle unit cells which could take into

ccount the particle distribution [20,21].These methods to predict the mechanical behaviour of real

omposites are based on the properties of the matrix and theeinforcements and the damage micromechanisms experimen-ally observed. The mechanical response of the matrix is usuallyssumed as the behaviour of the unreinforced alloy and the par-

icle properties as the bulk ceramic. However, the microstructuref the matrices is modified due to the reinforcement and soheir mechanical response. The grain size of the matrices iseduced respect to the unreinforced alloys, it is observed a higher

e and

dhdm[fa

pttmaottlswfotplvtaotnstt

mpr[bmntisdcpsmoSmofiitl

tn

2

abodabt2A4rusds8dmedo2ttwa8(pc〈1 0 0〉 directions in the matrix. The main difference betweenthe composite and the unreinforced alloy was the developmentof a �′ precipitate free zone (PFZ), located along the matrix-reinforcement interface and around 200 nm thick. In addition,

J. Rodrıguez et al. / Materials Scienc

islocation density in the matrices and the nucleation of inco-erent precipitates is favoured by the presence of reinforcementsue to the higher defect’s density [1,6,22]. For these reasons theatrices are expected to be harder than the unreinforced alloys

1,22]. In addition, the mechanical properties of the ceramic rein-orcements could be modified from those of the bulk ceramicss their processing and treatment could be rather different [23].

The evaluation of in situ mechanical properties of the com-osites’ constituents is necessary to improve the accuracy ofhe predictions obtained by the models and so, to increasehe knowledge of the micromechanisms responsible for the

echanical response in metal matrix composites. However, thisnalysis has experimental limitations due to the small scalef the components, which makes difficult the use of conven-ional microindentation techniques to estimate the hardness ofhe individual constituents. The measurement areas are usuallyarger than the reinforcements. The matrix measurement pointshould be selected at the positions where the reinforcementsere widely separated, which is often quite difficult, especially

or composites with high volume fraction of ceramics. Furthern, the reinforcements should be located at the subsurface ofhe indentation area. To overcome these limitations it is pro-osed the use of nanoindentation, a method developed over theast decade for probing the mechanical properties of materials atery small scale and used in the mechanical characterization ofhin coatings. Nanoindentation uses high resolution sensors andctuators to continuously monitor the loads and displacementsn an indenter as it is driven into the material studied. Analysis ofhe load–displacement curve provides information about hard-ess and elastic modulus of the material tested at the sub-microncale [24,25]. Imaging of the nanoindentation is not required andhe area involved in the measurement is much smaller than inhe microindentation tests.

Nanoindentation has been used recently to evaluate in situechanical properties of metal matrix composites in a few

apers. Matrix hardness has been estimated [26] and its valueelated to the accelerated ageing behaviour of these materials27,28]. The effect of particle clustering on residual stresses haseen also studied [29]. The matrix-reinforcement interface inaterials processed using coated particles has been probed by

anoindentation to analyse the reactions between the matrix andhe ceramic reinforcements [30]. In addition, in situ mechan-cal properties of fibre reinforced composites have been alsotudied [31,32]. However, there is not a systematic use of nanoin-entation to evaluate the in situ mechanical properties of theonstituents in discontinuously reinforced metal matrix com-osites and this is the aim of this investigation. Nanoindentationhould provide the hardness and elastic modulus of the metalatrix and the ceramic reinforcements. This work was carried

n a commercial Al/Li alloy 8090 reinforced with 15 vol.% ofiC particles and the unreinforced alloy was used as a controlaterial. The investigation was completed with the simulation

f a nanoindentation test on the unreinforced alloy using the

nite elements method. The mechanical properties of the unre-

nforced alloy are well known from tensile tests performed onhis material [6,33] and these properties could be used to simu-ate the nanoindentation test. The simulated test provides addi-

Fe

Engineering A 437 (2006) 406–412 407

ional information to understand the materials’ behaviour duringanoindentation.

. Materials and experimental techniques

The 8090 Al alloy reinforced with 15 vol.% of SiC particlesnd the unreinforced alloy were supplied by Cospray (Ban-ury, United Kingdom) in the form of extruded rectangular barsf 25.4 mm × 62.5 mm cross section. Both materials were pro-uced by spray codeposition of the matrix and the particles ontosubstrate [5]. They were extruded at 420 ◦C into a rectangularar, and the extrusion ratio was 25:1. The bars were solution heatreated at 530 ◦C for 2 h, water quenched and cold stretched up to% to relieve the residual stresses introduced during quenching.fterward, they were artificially aged at 170 ◦C during 32 and8 h for the composite and the unreinforced alloy, respectively, toeach the peak-aged condition (T651). The grain structure of thenreinforced alloy was highly anisotropic and grains larger thaneveral hundreds of micrometers were found in the longitudinalirection. The average grain sizes in the long and short transver-al directions, perpendiculars to the extrusion axis, were 21 and�m, respectively. The ceramic particles homogenized the grainimensions in the composite and the average grain size of theatrix was 12 �m in the longitudinal direction, parallel to the

xtrusion direction, and 6 �m in the long and short transversalirections, perpendiculars to the extrusion direction. The sizef the reinforcements was 7.5 ± 2.4 �m and the aspect ratio.4 ± 1.2. The particles were oriented with the longer axis inhe extrusion direction (Fig. 1). The main precipitate found washe metastable and ordered �′(Al3Li) phase, which is coherentith the aluminum matrix and nucleates homogeneously, with

n average diameter around 25 nm. The presence of Zr in the090 alloy to refine the grain size lead to the development of �′Al3Zr) dispersoid. S′ (Al2CuMg) phase was also observed. Thisrecipitate, semicoherent with the matrix, has an orthorhombicrystal structure and is observed as needles oriented along the

ig. 1. Longitudinal section of the composite showing grain structure and thelongated SiC particles, which were oriented in the extrusion direction.

408 J. Rodrıguez et al. / Materials Science and

Table 1Macroscopic mechanical properties of the unreinforced alloy and the compositeobtained from tensile tests [6]

Material E (GPa) σy (MPa) σu (MPa) εu (%)

88

thfmaet[yg

apppTarorpauil

H

wc

E

wb

S

β

r

wrctb

Bf

h

w

un1Fmto(tations. These observations were also used to analyse the influ-ence of distance to reinforcement particles on aluminium matrixproperties.

090 T651 81 552 581 2.3090 + SiC T651 102 523 568 1.5

he dislocation density observed in the composite was muchigher than in the unreinforced alloy, as it could be expectedrom the thermal expansion coefficient mismatch between theetal matrix and the ceramic reinforcements. Further details

bout the microstructure of these materials have been describedlsewhere [6]. The macroscopic mechanical properties at roomemperature of these materials have been evaluated previously6,33] and are summarized in Table 1, were elastic modulus, E,ield strength, σy, tensile strength, σu, and failure strain, εu, areiven.

Metallographic samples were prepared in the longitudinalnd transversal directions for nanoindentation tests. The sam-les were cut from the bars with a diamond saw and initiallyolished on SiC paper to 500-grit finish. This was followed byolishing on a diamond slurry (up to 1 �m) and, finally, on silica.he nanoindentation tests were performed on both sections usingMTS Nanoindenter XP with a diamond Berkovich tip 100 nm

adius. The tests were carried out on the unreinforced alloy andn the two constituents of the composite: the matrix and theeinforcement. During the nanoindentation tests the load on sam-le and the penetration depth were recorded. Young’s modulusnd hardness were estimated from the load–displacement curvesing the Oliver and Pharr [24,25,34] method for a Berkovichndenter. These properties are calculated according to the fol-owing expressions:

= Pmax

A, (1)

here Pmax is the maximum load and A is the projected area ofontact at maximum load. On the other hand,

∗ = S1

2

√π√A

1

β, (2)

here S is the slope of the load-depth of penetration curve at theeginning of the elastic unloading

= dP

[dh] (h=hmax)(3)

is a constant dependent on the indenter geometry and E* is theeduced or effective modulus determined by:

1

E∗ = 1 − ν2

E+ 1 − ν′2

E′ , (4)

here E, ν, E′ and ν′ represent Young’s modulus and Poisson

atio of the material and the indenter, respectively. In the pre-eding expressions, the projected contact area, A, is evaluatedhrough the indenter shape function, A = A(hc), a relationshipetween A and the contact depth, hc (A = 24, 5h2

c for an ideal

Ftm

Engineering A 437 (2006) 406–412

erkovich tip). Finally, this last magnitude can be calculatedrom the expression:

c = hmax − εPmax

S, (5)

here ε is again a constant dependent on the indenter geometry.Initially, maximum loads ranging from 1 to 100 mN were

sed. The influence of load on the properties measured wasegligible at lower loads, although the properties obtained at00 mN differed slightly from those got at 1.5 and 10 mN.inally, maximum load of 5 mN was chosen to characterize theseaterials. The indentations were distributed homogeneously

hroughout the materials studied and the samples tested werebserved in an environmental scanning electron microscopeESEM) Philips XL 30 to analyse the morphology of the inden-

ig. 2. Experimental load–depth of penetration curves after nanoindentationests. (a) Ceramic reinforcements and (b) unreinforced alloy and composite’s

atrix.

J. Rodrıguez et al. / Materials Science and Engineering A 437 (2006) 406–412 409

Table 2Elastic modulus and hardness of the unreinforced alloy and the composite constituents measured by nanoindentation. Reference values for bulk SiC are also included

Bulk SiC [35] SiC particles Al 8090 alloy Composite’s matrix

EH

3

tnlefntuimpdevtofip

mTisFatum

Fc

htafdahdue to the high experimental scatter.

The unreinforced alloy macroscopic tensile properties(Table 1) could be compared with the micromechanical

(GPa) 414 300 ± 50(GPa) 33 31 ± 6

. Micromechanical characterization

Nanoindentation tests performed on the longitudinal andransversal directions of the original extruded bars showedo appreciable differences on the properties measured. Theoad–displacement curves obtained for the reinforcementsxhibited a typical elastic behaviour, as it could be expectedor ceramic particles (Fig. 2a). The elastic modulus and hard-ess, estimated from this curve (Table 2), are slightly smallerhan those found in the literature [35] for bulk SiC, this could benderstood as the processing of ceramic reinforcements couldntroduce more defects than the processing routes used for bulk

aterials. The unreinforced alloy and the composite’s matrixresented a plastic behaviour (Fig. 2b) which results slightlyifferent when both curves are compared. These small differ-nces observed are enough to provide significantly differentalues of hardness (Table 2) and the matrix was harder thanhe unreinforced alloy. The elastic modulus measurements wereverlapped and no differences should be expected 200 nm awayrom the reinforcements, where the distribution of �′ precipitatess similar in both materials. However, the presence of sub-surfacearticles could increase the elastic modulus and also the scatter.

The influence of distance from particle reinforcements on theatrix properties was analysed using the ESEM observations.he distribution of tests through the matrix could be observed

n Fig. 3 where an array of indentations in the composite ishown. A summary of the experimental results is presented inig. 4: both Young’s modulus (Fig. 4a) and hardness (Fig. 4b)re represented versus the distance to the nearest reinforcement

ogether with the experimental bands corresponding to thenreinforced alloy. The mechanical behaviour of the aluminiumatrix differs from that observed in the unreinforced alloy. The

ig. 3. Secondary electron image showing an array of indentations on the matrixomposite.

F(a

90 ± 2 98 ± 151.86 ± 0.05 2.2 ± 0.2

ardness of the matrix slightly decreased with the distance tohe nearest reinforcement particle, but even its lower value wasround 15% higher than the hardness measured on the unrein-orced alloy. This result is in agreement with the microstructureifferences usually found between metal matrix compositesnd the corresponding unreinforced alloys [1,22]. On the otherand, the elastic modulus variation is not clearly appreciated

ig. 4. Influence of distance from the nearest particle on matrix Young’s modulusa) and hardness (b). Experimental bands, with the properties of the unreinforcedlloy, have been also included.

410 J. Rodrıguez et al. / Materials Science and Engineering A 437 (2006) 406–412

Fp

ptmetiue[iodaempea

4

siaumrtTithttm(csd

Fs

twttTmctTo

imt

swnctf

aaa

vftmbalT

ig. 5. Secondary electron image showing a nanoindentation on the matrix com-osite. The material tested is piled up around the nanoindentation.

roperties measured in situ. Nanoindentation tests overestimatehe elastic modulus from 81 to 90 GPa. The Oliver and Pharr

ethod, used in this work to obtain hardness and modulus,stimates the contact area, A, through the indenter shape func-ion which only considers the contact depth, hc, for a particularndenter, as it has been explained above. If the contact area isnderestimated using this approach the micromechanical prop-rties obtained through Eqs. (1) and (2) are clearly overestimated24,25,34]. The nanoindentations’ morphology was analysedn the ESEM and images similar to that shown in Fig. 5 werebserved. The material tested was piled up around the nanoin-entation. As a consequence, the area estimated using the Olivernd Pharr method is smaller than the real contact area. Thisffect, known as pile-up, has been observed previously in ductileaterials [25] and leads to overestimate the micromechanical

roperties measured by nanoindentation. To analyse how thisffect takes place and its influence on the properties measured,nanoindentation test was simulated on the unreinforced alloy.

. Finite element modelling

An elastic–plastic indentation with a Berkovich indenter wasimulated in a two-dimensional model assuming that the actualndenter provides similar results that a conical tip with a 70.3◦ngle. With this simplification, computation time was reducedsing the axisymmetric capabilities of the ANSYS® finite ele-ent code [37]. The indenter and specimen were treated as

evolution bodies. The indenter was modelled as a rigid conicalip with a semivertical angle θ = 70.3◦ edged radius of 100 nm.hese are the geometrical characteristics of the actual tip. The

ndenter and the specimen are shown in Fig. 6, together withhe boundary conditions: the nodes along the axis symmetryave their movements limited along the y-axis; the nodes athe bottom of the mesh are fixed along the y-axis. To modelhe specimen, a high order two-dimensional eight-node ele-

ent was used having two degrees of freedom at each node

translation in the nodal x and y directions) with compatibleontact surface elements. A surface-to-surface algorithm waselected with the augmented Lagrangean method and rigid-eformable contact pair behaviour. Two adjustable parameters,

mcdp

ig. 6. Finite elements mesh used for the simulated test including indenter andpecimen.

he contact stiffness (FKN) and the penetration tolerance (FTOL)ere selected in such a way that the simulated an experimen-

al curves were matched. Acceptable results were obtained whenhe values applied were FKN = 0.3 and FTOL = 0.1, respectively.hese parameters are independent of the indented material’sechanical properties, as they are only related to the equilibrium

onditions for the rigid-plastic contact. The interface betweenhe indenter and the specimen was considered to be frictionless.his assumption is valid for a Berkovich indenter as the effectf friction is negligible for this geometry [38].

The material behaviour is elastic plastic with multi-linearsotropic hardening and assumed to be initially stress free. The

aterial properties used in the model are those correspondingo the Al 8090 (Table 1).

A downward displacement was imposed to the indenter toimulate a typical indentation process. The corresponding loadas achieved by summing the reaction forces on the bottomodes for a given penetration length. The load–displacementurve during the simulated process can be plotted and comparedo the experimental one. The Von Mises yield criterion is appliedor determining the occurrence of the plastic deformation.

Numerical model was checked comparing the experimentalnd numerical load–displacement curves in both, the loadingnd the unloading branches. Fig. 7 shows the close agreementchieved between numerical and experimental measurements.

The analysis of the numerical load–displacement curves pro-ides the Young’s modulus and hardness of the material indentedollowing the same procedure as in the experimental case. Addi-ionally, the finite element simulation makes possible the esti-

ation of the pile-up influence. The actual area of contact cane determined identifying in the model the last point of contactt maximum load (Fig. 8). This value can be used to recalcu-ate hardness and modulus. The results obtained are included inable 3.

This procedure can be applied to the unreinforced alloy. Their

echanical properties, known from macroscopic tension tests,

an be supplied to the finite element code to simulate a nanoin-entation experiment. In the case of the composite’s matrix, itsroperties are unknown (this is, actually, the objective of this

J. Rodrıguez et al. / Materials Science and

Fig. 7. Simulated and experimental load–displacement curves. Close agreementis observed between both.

Fc

wdcfacfam

TEm

EH

5

mtnidttriigasdsmticaso

bcTsmd[dudtthtadp

mf

ig. 8. Contact profile obtained from the simulated test. The actual area ofontact is determined identifying in the last point of contact at maximum load.

ork) which impedes a direct application of the same proce-ure. To overcome this difficulty the following hypothesis wasonsidered: the pile-up amount is similar for both, the unrein-orced alloy and the composite’s matrix. The ratio between thectual area and the nominal one is considered the same in bothases. With this assumption, the finite element model developedor the unreinforced alloy provides information about the pile-up

mount and this same value is applied to correct the composite’satrix data. The results obtained are presented in Table 3.

able 3lastic modulus and hardness of the unreinforced alloy and the composite’satrix considering pile-up correction

Al 8090 alloy Composite’s matrix

(GPa) 81 83(GPa) 1.38 1.6

mfpImv

bmcm

Engineering A 437 (2006) 406–412 411

. Discussion

The hardness measured using nanoindentation in the alu-inium matrix was more than 15% higher than that measured in

he unreinforced alloy. In addition, the aluminium matrix hard-ess was reduced as the distance to the nearest reinforcement wasncreased. Although this absolute values could be overestimated,ue to the pile-up effect observed around the nanoindentations,he relative variations clearly indicates a different behaviour ofhe matrix compared to the unreinforced alloy. Further on, theseesults are in agreement with the microstructural differences typ-cally observed in MMC. In particular, for the materials studiedn this investigation, the aluminium matrix exhibited a smallerrain size and a higher dislocation density than the unreinforcedlloy [6]. A reduction in grain size represents an increment intrength through the Hall-Petch effect. In addition, the higherislocation density reduces the dislocation’s mean free pathtrengthening the aluminium matrix [36]. This strengtheningechanism is more important close to the ceramic particles, as

he dislocations are emitted from the reinforcements during cool-ng due to the different thermal expansion coefficients of bothonstituents. The relationship between the indentation resultsnd elastoplastic properties like yield stress or ultimate tensiletrength is complicated, but it seems necessary to consider thesebservations in the MMC micromechanical models.

The differences observed in the Young’s modulus determinedy macroscopic mechanical tests and nanoindentation, is an indi-ation of how the pile-up phenomenon affects the data reliability.he results obtained in this work are based on the assumption ofimilar influence in the unreinforced alloy and the composite’satrix. Several works have addressed the pile-up development

uring indentations of elastic–plastic solids during the last years39–41]. Most of them report that the pile-up effect is depen-ent on the mechanical properties, but no abrupt variations weresually observed. Although the aim of this work is to captureifferences between the properties of the unreinforced alloy andhe composite’s matrix, radical changes are not expected. Thus,he underlying hypothesis is rather plausible. In addition, thisypothesis leads to a matrix hardness about 20% higher thanhat estimated for the unreinforced alloy (Table 3). These valuesre reasonable considering the strengthening of the compositeue to the microstructural changes discussed in the previousaragraph.

A reasonable alternative would be to consider similar Young’sodulus instead of comparable pile-up amount. This is also a

easible assumption, but it is more difficult to apply. The Young’sodulus provided by the nanoindentation may be affected by

actors difficult to assess, such as the presence of reinforcementarticles under the zones where the measurement is performed.n fact, the modulus measurements carried out in the composite’satrix are affected by a clearly upper scatter and with average

alues quite higher than those of the unreinforced alloy.As it has been explained, the simulations carried out are

ased on a continuum mechanics approach and the indentedaterial plastic deformation is controlled by the von Mises

riterion. This approach reproduces quite accurately the experi-ental load–penetration curve. However, the size of the material

4 e and

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R

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12 J. Rodrıguez et al. / Materials Scienc

ndented is comparable to the grain size and further develop-ents could include crystal plasticity to simulate the material

ested. Although the benefits of this refinement should be bal-nced with the complexity involved.

. Conclusions

This investigation has analysed the in situ mechanical prop-rties of an Al–Li alloy 8090 reinforced with 15 vol.% SiCarticles using nanoindentation. This technique has been ableo probe the mechanical properties of the metal matrix compos-te constituents despite the small scale characteristic of these

aterials. The mechanical properties measured on the ceramiceinforcements were slightly smaller than those found in the lit-rature for bulk materials, probably due to the different manufac-uring process used. The mechanical behaviour of the aluminium

atrix is rather different from that observed in the unreinforcedlloy. The hardness of the matrix was higher than that of the unre-nforced alloy and the values measured decreased as the distanceo the nearest reinforcement was increased. The unreinforcedlloy and the composite matrix piled up around the indenta-ion and this effect overestimated the properties measured byanoindentation. The effect of pile-up could be evaluated inhe unreinforced alloy by a simulated test carried out using thenite elements method. Assuming similar pile-up in the matrixomposite and in the unreinforced alloy the micromechanicalroperties of both materials could be corrected. The compos-te’s matrix hardness is about 20% higher than that estimatedor the unreinforced alloy. These results are in agreement withhe microstructural characteristics typically observed in MMCnd should be considered in the micromechanical models ofhese materials.

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