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Available online at www.sciencedirect.com

ScienceDirect

Additive Manufacturing 6 (2015) 1–5

Production of high strength Al85Nd8Ni5Co2 alloy by selective laser melting

K.G. Prashanth a,∗, H. Shakur Shahabi a, H. Attar a,b, V.C. Srivastava c, N. Ellendt d,V. Uhlenwinkel d, J. Eckert a,e, S. Scudino a

a IFW Dresden, Institut für Komplexe Materialien, Postfach 27 01 16, D-01171 Dresden, Germanyb School of Engineering, Edith Cowan University, 270 Joondalup Drive, Joondalup, Perth, WA 6027, Australia

c Metal Extraction & Forming Division, National Metallurgical Laboratory, Jamshedpur 831007, Indiad Institut für Werkstofftechnik, Universität Bremen, D-28359 Bremen, Germanye TU Dresden, Institut für Werkstoffwissenschaft, D-01062 Dresden, Germany

Accepted 5 January 2015Available online 13 January 2015

bstract

Bulk high strength and thermally stable Al85Nd8Ni5Co2 samples have been prepared by selective laser melting (SLM). The alloy shows aomposite-like microstructure consisting of submicron-sized stable intermetallic phases dispersed in an Al matrix, which leads to high compressive

trength (1–0.5 GPa) at elevated temperatures (303–573 K). These results indicate that SLM is an effective alternative to conventional routes forroducing dense, thermally stable and near net shaped components from high strength Al-based alloys.

2015 Elsevier B.V. All rights reserved.

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eywords: Additive manufacturing; Aluminum alloys; Intermetallic compound

. Introduction

Aluminum and its alloys are among the most widely usedaterials for structural and functional applications, due to their

igh specific strength, high corrosion resistance and good pro-essability [1]. The strength of Al-alloys can be increased byhe formation of amorphous/glassy (MG), nanocrystalline (NC)r ultrafine grained (UFG) structures [2–4]. The MGs can bebtained by non-equilibrium processing techniques like meltpinning, mechanical alloying, gas atomization or copper moldasting [5–7]. However, Al-based MGs have poor glass forma-ility and hence their size is limited, typically <1 mm [8,9]. TheC/UFG alloys exhibit thermal instability due to the excess

nthalpy associated with the high density of grain boundaries;hereby limiting their high temperature application spectrum

10,11]. Therefore, in order to utilize the advantages of the

G/NC/UFG structured alloys, there is a strong need to develophermally stable Al-based alloys without size limitations.

∗ Corresponding author. Tel.: +49 351 4659 685; fax: +49 351 4659 452.E-mail addresses: kgprashanth@gmail.com,

.g.prashanth@ifw-dresden.de (K.G. Prashanth).

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ttp://dx.doi.org/10.1016/j.addma.2015.01.001214-8604/© 2015 Elsevier B.V. All rights reserved.

mpression test

Selective Laser Melting (SLM) is an additive manufactur-ng technique capable of producing MG/NC/UFG materials12–14]. It allows the fabrication of complex and intricateeometries with a high degree of accuracy, high design flexibil-ty along with excellent process capabilities and high materialtilization [12–14]. Although reports exist on the productionf Al–Si and Al–Si–Mg based alloy systems by SLM, otherystems were not explored extensively and systematically. Thel–Zn (7XXX) system in one of the commercial and con-entional Al-based alloy systems that exhibits high strength.owever, there are no reports on the fabrication of 7XXX alloysy SLM, which might be a result of two reasons: (1) they arextremely brittle and may lead to cracking of the samples duringhe fabrication process and (2) evaporation of Zn (which has lowoiling point) during the SLM process, which makes it unsuit-ble for the SLM process. Hence, there is a strong need to explorehe various unconventional Al-based alloys (MG/NC/UFG) that

ay exhibit high strength at ambient temperatures and can beuccessfully fabricated by SLM.

Recently, Li et al. have reported the production of amorphous

l86Ni6Y4.5Co2La1.5 by SLM [15,16]. They have performed

ingle line scans at different laser powers on a pre-fabricatedorous Al86Ni6Y4.5Co2La1.5 metallic glass perform that lead

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K.G. Prashanth et al. / Addi

o the formation of a gradient-like microstructure with crystal-ization of the amorphous phase at some places. They have alsohown that the microstructure can be controlled by controllinghe laser parameters. However, several cracks and micro-cracksere observed due to high thermal gradient observed during therocess which caused high levels of thermal stress.

Following this study Li et al. investigated the use of a re-canning strategy to prevent the macro-cracking during the SLMrocess. They have claimed that a high power initial scan fol-owed by a lower power re-scan strategy can be used to avoidracking Al-based MG/NC alloys. However, even by adoptionf such a re-melt strategy, completely amorphous and defectree Al-based samples were not possible by SLM. However,hese works throw light on the potential fabrication of Al-ased NC/UFG materials by SLM and furthermore there are noetailed reports on the fabrication and microstructural propertyorrelation with the mechanical properties. The present worknalyzes this aspect by focusing on the production of a thermallytable, high strength Al-based NC alloy using the SLM process.his is followed by a detailed structural and microstructural

nvestigation, along with the mechanical properties evaluation,racture analysis and structure–property correlation.

. Materials and methods

Al85Nd8Ni5Co2 (at.%) cylindrical specimens (3 mm diame-er and 8 mm height) were produced by SLM from sphericalas-atomized powder (GAP) using an SLM 250 HL deviceSLM Solutions GmbH, Luebeck, Germany) equipped with anb-YAG laser with a maximum power of 400 W and a spot

ize of ∼80 �m. The gas atomized powder was spherical inhape with an average particle size of 48 ± 5 �m. The powderxhibited excellent flowability, which is one of the importantre-requisites to be used as a raw material for the SLM pro-ess. The parameters used for the fabrication of the specimensre: power of 320 W for volume and contour, layer thickness of0 �m, stripe hatch with a spacing of ∼110 �m between themnd hatch style rotation of 73◦ between the layers. Two differ-nt scanning speeds were used: 1455 mm/s for the volume and939 mm/s for the contour. The Al substrate plate was heated to73 K during the entire SLM process to avoid the formation ofracks in the SLM samples. Structural analysis was performedy X-ray diffraction (XRD) using a D3290 PANalytical X’pertRO (PANalytical GmbH, Kassel-Waldau, Germany) with Co-� radiation (λ = 0.17889 nm) in Bragg-Brentano configuration.he Rietveld method was employed for estimating the crystal-

ite size from the XRD patterns using the WinPlotR softwareackage [17].

The density of the consolidated samples was evaluated byhe Archimedes principle. The microstructure was characterizedy scanning electron microscopy (SEM) in the back scatteredlectron (BSE) mode using a Gemini 1530 microscope (Göt-ingen, Germany) equipped with an energy-dispersive X-ray

pectroscopy (EDX) facility. The compression tests were car-ied out using an Instron 8562 testing system (Instron GmbH,armstadt, Germany) under quasistatic loading (strain rate1 × 10−4 s−1) in the temperature range 303–573 K. The strain

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anufacturing 6 (2015) 1–5

as measured directly on the specimen using a Fiedler laser-xtensometer. The samples were always tested for compressionlong the build direction, which implies the top of the samplefter fabrication is held at the top during the compression tests well. In order to ensure the microstructural stability of theaterial during the high-temperature tests, the specimens were

eat treated under argon atmosphere at 723 K for 4 h. The hard-ess of the individual phases was determined using an “AsmecNAT” nano-indentor (ASMEC GmbH, Radeberg, Germany)ith a Berkovich shape tip. A total number of 150 indentationsere performed on a highly polished surface using a typicaluadratic loading and unloading procedure. In order to have thendentation on a single phase with sub-micron size, a maximumoad of 2 mN was selected.

. Results and discussion

Fig. 1 shows the SEM and EDX mapping images of thel85Nd8Ni5Co2 as-prepared (AP) SLM sample taken along its

ross-section. The low magnification image (Fig. 1(a)) revealshe typical track morphology observed in the SLM specimens14,18]. The distance between the tracks is ∼100 �m with arack overlap of ∼10 �m and no visible porosity and defectsre seen, which is also corroborated with the density measure-ent studies, where an average relative density of 99.75% is

bserved. It is to be noted that the hatch distance used for fab-icating these samples are 110 �m and hence the tracks shouldxhibit a distance of 110 �m. However, it can be observed fromig. 1(a) that the width of the tracks is only ∼100 �m, indi-ating that there is a presence of hatch overlaps of ∼10 �m.uch strategy of using hatch overlaps reduces the porosity lev-ls as well as any possible discontinuities in the sample betweenhe hatches, there by resulting in a sound sample with near toull density. It has to be noted that there is a rotational shift ofround 15 deg between Fig. 1(a) and (b) as a result of the mea-urement sequence. The microstructure is non-uniform with thehases exhibiting a bimodal distribution. This is the result of therack overlap – core morphology typical for the SLM samples14], consisting of a fine microstructure along the track over-aps (marked as (1) in Fig. 1(b)) and a coarse microstructurelong the track cores (marked as (2) in Fig. 1(b)). Bright plateletsith different phase contrast are distributed within a dark matrix

Fig. 1(c)), making it a composite-like microstructure.The elemental mapping EDX images (Fig. 1(d–h)), show that

he dark areas are rich in Al and the bright platelets are rich in Nd,i and Co. The bright platelets have different contrasts suggest-

ng the presence of four different phases in the Al85Nd8Ni5Co2P SLM samples. The presence of four different phases in theP SLM sample was confirmed by XRD. The diffraction pat-

ern (Fig. 2) displays the existence of �-Al (cubic, Fm3m) with crystallite size d = 72 nm, along with the three intermetallichases: AlNdNi4 (orthorhombic, Cmc2), Al4CoNi2 (cubic,a3d) and AlNd3 (hexagonal, P63mmc). The AlNdNi4 platelets

d = 29 nm) are 2.12 ± 0.34 �m in length and 0.41 ± 0.13 �m inidth, whereas the Al4CoNi2 (d = 42 nm) and AlNd3 platelets

d = 35 nm) are 2.32 ± 0.48 �m and 6.01 ± 0.74 �m in lengthnd 1.00 ± 0.08 �m and 0.89 ± 0.14 �m in width, respectively.

K.G. Prashanth et al. / Additive Manufacturing 6 (2015) 1–5 3

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ig. 1. SEM images of the Al85Nd8Ni5Co2 alloy showing (a) the laser tracks anhe intermetallic phases with different contrasts embedded in the Al matrix and

he XRD patterns of the different samples indicate that thehases present in the GAP, AP SLM and HT SLM samples arehe same (Fig. 2). Interestingly, the XRD patterns of the AP SLMnd HT SLM are very similar. The crystallite sizes of the phasesn HT SLM are 84 nm for Al, 34 nm for AlNdNi4, 50 nm forl4CoNi2 and 38 nm for AlNd3, indicating that the SLM sam-le is thermally stable with no significant phase transformationnd grain growth taking place during the heat-treatment process.

The AP SLM specimen tested at room temperature (RT)hows a very high yield strength (YS) of 0.94 GPa and an ulti-ate compressive strength (UCS) of 1.08 GPa along with 2.45%

lastic strain (Fig. 3(a and b)). The HT sample shows simi-ar properties (YS = 0.81 GPa and UCS = 0.97 GPa), suggesting

hat the microstructure of the SLM material is thermally sta-le at high temperatures. Such high strength levels are retainedt elevated temperatures. For example, UCS of 1.05 GPa and

ig. 2. XRD patterns (λ = 0.17889 nm) of the gas atomized powder along withs-prepared and heat-treated SLM samples.

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the bimodal microstructure. High magnification SEM micrograph showing (c) the corresponding EDX images.

.97 GPa, and strain of 4.5 and 8.5% are observed for the sam-les tested at 373 K and 423 K, respectively (Fig. 3(a and b)).ith further increase of the testing temperature, the deforma-

ion exceeds 20%, where the measurements were stopped. Event a testing temperature of 573 K, a high UCS of ∼0.50 GPa isbserved, which indicates that the present material can be usedor high strength applications at high operating temperatures.

Although, intense scientific research has been focused onhe development of high-strength amorphous and nanostruc-ured Al-based alloys, only a few reports deal with the highemperature mechanical properties of these materials. The sin-ered nano-crystalline Al–Fe material [19] displays the best highemperature properties among the published works on Al-basedlloys. These results are compared with the mechanical proper-ies of the present Al–Nd–Ni–Co SLM alloy in Fig. 3(c). It cane observed that the present alloy has similar strength levels withespect to the Al–Fe alloy at all test temperatures. The Al–Felloy was produced by spark plasma sintering, where the sizend shape of the component is restricted, whereas processing byLM permits the production of parts having theoretically anyossible geometry with minimized need for post processing.herefore, additive manufacturing offers the possibility to tailor

he shape and corresponding properties of these high-strengthl-based parts to meet specific requirements, which renders this

echnology unique in comparison to conventional processing.The fracture surface images, as shown in Fig. 4, give further

vidence for the high strength observed in the Al85Nd8Ni5Co2lloy. The fracture takes place in a stepped morphologyFig. 4(a)), similar to the Al–12Si alloy prepared by SLM [12].his can be ascribed to the bimodal microstructure (Fig. 3(b))ith fine grains along the track-overlaps, which may act as areferential path for crack propagation, leading to the observedtepped morphology. The superior properties observed in the

resent alloy can be attributed to the composite-like microstruc-ure. The fine microstructure, characteristic of the SLM process12], leads to the presence of fine �-Al phase surrounded by

4 K.G. Prashanth et al. / Additive Manufacturing 6 (2015) 1–5

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ig. 3. (a) Compressive stress–strain curves of the as-prepared (AP) and heat-tlong with the AP SLM samples tested at different temperatures, (b) correspondil–Fe alloy produced by spark plasma sintering [19].

he intermetallic phases. Since the microstructure is developedrom a rapid solidification process, it is expected to have annherently strong bonding between the Al matrix and the inter-

etallic reinforcements, aiding an improved and effective loadransfer along the interface. This concept leads to the interfacialtrengthening mechanism in the present alloy both at RT and atigh temperatures.

The hardness of the intermetallic phases evaluated fromhe nano-indentation measurements are 2.84 ± 0.08 GPa forlNd3, 3.74 ± 0.10 GPa for Al4CoNi2 and 5.45 ± 0.09 GPa

or AlNdNi4; which is much higher than pure �-Al0.33 ± 0.03 GPa). Hence, during RT deformation, the cracksre expected to develop along the �-Al phase and to propagate

ith further loading. However, the intermetallic platelets act asbstacles which leads to either the arrest or deflection of theracks, as observed from Fig. 4(b and c). In case of the crack

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ig. 4. Fracture surface after compression tests showing (a) stepped morphology, (brrest and crack deflection by the intermetallic phases in the Al85Nd8Ni5Co2 alloy du

(HT; 4 h at 723 K) SLM samples tested at room temperature (curves 1 and 2)chanical data and (c) comparison of the present results with the nano-crystalline

rrest, further deformation proceeds through the initiation ofew cracks leading to crack multiplication (Fig. 4(b)). On thether hand, crack deflection suggests that the effective meanrack path is increased leading to appreciable deformation inhe material [20].

Generally, at high temperatures the dislocation movements accelerated and the magnitude of the Peierls stress israstically reduced leading to low strength of the material [20].owever, in the present alloy, the Al matrix is surroundedy the intermetallic reinforcement, which may confine theislocation movement along the grain boundaries according tohe confinement theory, which leads to further strengthening ofhe material at elevated temperatures [21,22]. All of the mech-

nisms: interfacial strengthening, crack arrest and initiationf new cracks, crack deflection phenomena and confinementhenomena (Fig. 4(d)) operate simultaneously leading to

) multiple cracks and (c) crack deflection. (d) Schematic illustrating the crackring compression test.

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uperior room temperature as well as high temperatureompressive strengths.

. Summary

Highly dense high strength Al85Nd8Ni5Co2 alloy has beenuccessfully prepared by SLM. The alloy exhibits a com-osite microstructure with the intermetallic phases AlNdNi4,l4CoNi2 and AlNd3 dispersed in the Al-matrix. The inter-etallic phases are in the form of platelets and their width ranges

n the sub-micron regime. The AP SLM and HT SLM sampleshow a UCS of 1.08 GPa and 0.97 GPa with ∼2.5% strain at RT.he high temperature compression tests reveal that the presentystem can retain their high strength due to the composite-likeicrostructure and confinement phenomena, where the grain

oarsening and the accelerated mobility of the dislocations atigh temperatures are retarded by the intermetallic phases. Addi-ionally, interfacial strengthening, the crack arrest and crackeflection mechanisms also contribute to the superior strengthbserved in the Al85Nd8Ni5Co2 alloy. The present results indi-ate that SLM is one of the best options to produce high strength,hermally stable, dense and near-net shaped Al-based alloys.

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