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Available online at ScienceDirect Additive Manufacturing 6 (2015) 1–5 Production of high strength Al 85 Nd 8 Ni 5 Co 2 alloy by selective laser melting K.G. Prashanth a,? , H. Shakur Shahabi a , H. V. Uhlenwinkel d , J. Eck a P Joondalup Metallur sit?t Br 5 January 13 Abstract ve composite-lik strength indicate producing high ? K Compression 1. Introduction Aluminum and its alloys are among the most widely used materials for structural and functional applications, due to their high cessability the or obtained spinning, casting bility NC/UFG enthalp thereby [10,11] MG/NC/UFG thermally k.g.prashanth@ifw-dresden.de Selective Laser Melting (SLM) is an additive manufactur- ing technique capable of producing MG/NC/UFG materials [12–14]. It allows the fabrication of complex and intricate geometries with a high degree of accuracy, high design flexibil- http://dx.doi.or 2214-8604/? specific strength, high corrosion resistance and good pro- [1]. The strength of Al-alloys can be increased by formation of amorphous/glassy (MG), nanocrystalline (NC) ultrafine grained (UFG) structures [2–4]. The MGs can be by non-equilibrium processing techniques like melt mechanical alloying, gas atomization or copper mold [5–7]. However, Al-based MGs have poor glass forma- and hence their size is limited, typically <1 mm [8,9]. The alloys exhibit thermal instability due to the excess y associated with the high density of grain boundaries; limiting their high temperature application spectrum . Therefore, in order to utilize the advantages of the structured alloys, there is a strong need to develop stable Al-based alloys without size limitations. ? Corresponding author. Tel.: +49 351 4659 685; fax: +49 351 4659 452. E-mail addresses: kgprashanth@, (K.G. Prashanth). ity along with excellent process capabilities and high material utilization [12–14]. Although reports exist on the production of Al–Si and Al–Si–Mg based alloy systems by SLM, other systems were not explored extensively and systematically. The Al–Zn (7XXX) system in one of the commercial and con- ventional Al-based alloy systems that exhibits high strength. However, there are no reports on the fabrication of 7XXX alloys by SLM, which might be a result of two reasons: (1) they are extremely brittle and may lead to cracking of the samples during the fabrication process and (2) evaporation of Zn (which has low boiling point) during the SLM process, which makes it unsuit- able for the SLM process. Hence, there is a strong need to explore the various unconventional Al-based alloys (MG/NC/UFG) that may exhibit high strength at ambient temperatures and can be successfully fabricated by SLM. Recently, Li et al. have reported the production of amorphous Al 86 Ni 6 Y 4.5 Co 2 La 1.5 by SLM [15,16]. They have performed single line scans at different laser powers on a pre-fabricated porous Al 86 Ni 6 Y 4.5 Co 2 La 1.5 metallic glass perform that lead g/10.1016/j.addma.2015.01.001 2015 Elsevier B.V. All rights reserved. IFW Dresden, Institut für Komplexe Materialien, b School of Engineering, Edith Cowan University, 270 c Metal Extraction Aluminum alloys; Intermetallic compounds; Attar a,b , V.C. Srivastava c , N. Ellendt d , ert a,e , S. Scudino a ostfach 27 01 16, D-01171 Dresden, Germany Drive, Joondalup, Perth, WA 6027, Australia gical Laboratory, Jamshedpur 831007, India emen, D-28359 Bremen, Germany haft, D-01062 Dresden, Germany 2015 January 2015 been prepared by selective laser melting (SLM). The alloy shows a phases dispersed in an Al matrix, which leads to high compressive that SLM is an effective alternative to conventional routes for strength Al-based alloys. test 2 K.G. Prashanth et al. / Additive Manufacturing 6 (2015) 1–5 to the formation of a gradient-like microstructure with crystal- lization of the amorphous phase at some places. They have also shown that the microstructure can be controlled by controlling the laser parameters. However, several cracks and micro-cracks were observed due to high thermal gradient observed during the process scanning process. lo cracking of free these based detailed correlation analyzes stable, This in fracture 2. ter gas-atomized (SLM Yb-Y size shape e pre-requisites cess. are: 50 and ent 1939 673 cracks by PR K The lite package the by electron tingen, spectroscop ried Darmstadt, ～ was measured directly on the specimen using a Fiedler laser- extensometer. The samples were always tested for compression along the build direction, which implies the top of the sample after fabrication is held at the top during the compression test as well. In order to ensure the microstructural stability of the material heat ness UN with were quadratic indentation load 3. Al cross-section. the [14,18] track are ment observ ricating e Fig. cating Such els the full around surement phases track [14] laps along with ( the Ni ing AP AP tern a phases: Ia3d) ( width, ( and which caused high levels of thermal stress. Following this study Li et al. investigated the use of a re- strategy to prevent the macro-cracking during the SLM They have claimed that a high power initial scan fol- wed by a lower power re-scan strategy can be used to avoid Al-based MG/NC alloys. However, even by adoption such a re-melt strategy, completely amorphous and defect Al-based samples were not possible by SLM. However, works throw light on the potential fabrication of Al- NC/UFG materials by SLM and furthermore there are no reports on the fabrication and microstructural property with the mechanical properties. The present work this aspect by focusing on the production of a thermally high strength Al-based NC alloy using the SLM process. is followed by a detailed structural and microstructural vestigation, along with the mechanical properties evaluation, analysis and structure–property correlation. Materials and methods Al 85 Nd 8 Ni 5 Co 2 (at.%) cylindrical specimens (3 mm diame- and 8 mm height) were produced by SLM from spherical powder (GAP) using an SLM 250 HL device Solutions GmbH, Luebeck, Germany) equipped with an AG laser with a maximum power of 400 W and a spot of ～80 H9262m. The gas atomized powder was spherical in with an average particle size of 48 ± 5 H9262m. The powder xhibited excellent flowability, which is one of the important to be used as a raw material for the SLM pro- The parameters used for the fabrication of the specimens power of 320 W for volume and contour, layer thickness of H9262m, stripe hatch with a spacing of ～110 H9262m between them hatch style rotation of 73 ? between the layers. Two differ- scanning speeds were used: 1455 mm/s for the volume and mm/s for the contour. The Al substrate plate was heated to K during the entire SLM process to avoid the formation of in the SLM samples. Structural analysis was performed X-ray diffraction (XRD) using a D3290 PANalytical X’pert O (PANalytical GmbH, Kassel-Waldau, Germany) with Co- H9251 radiation (λ = 0.17889 nm) in Bragg-Brentano configuration. Rietveld method was employed for estimating the crystal- size from the XRD patterns using the WinPlotR software [17]. The density of the consolidated samples was evaluated by Archimedes principle. The microstructure was characterized scanning electron microscopy (SEM) in the back scattered (BSE) mode using a Gemini 1530 microscope (G?t- Germany) equipped with an energy-dispersive X-ray y (EDX) facility. The compression tests were car- out using an Instron 8562 testing system (Instron GmbH, Germany) under quasistatic loading (strain rate 1 × 10 ?4 s ?1 ) in the temperature range 303–573 K. The strain during the high-temperature tests, the specimens were treated under argon atmosphere at 723 K for 4 h. The hard- of the individual phases was determined using an “Asmec AT” nano-indentor (ASMEC GmbH, Radeberg, Germany) a Berkovich shape tip. A total number of 150 indentations performed on a highly polished surface using a typical loading and unloading procedure. In order to have the on a single phase with sub-micron size, a maximum of 2 mN was selected. Results and discussion Fig. 1 shows the SEM and EDX mapping images of the 85 Nd 8 Ni 5 Co 2 as-prepared (AP) SLM sample taken along its The low magnification image (Fig. 1(a)) reveals typical track morphology observed in the SLM specimens . The distance between the tracks is ～100 H9262m with a overlap of ～10 H9262m and no visible porosity and defects seen, which is also corroborated with the density measure- studies, where an average relative density of 99.75% is ed. It is to be noted that the hatch distance used for fab- these samples are 110 H9262m and hence the tracks should xhibit a distance of 110 H9262m. However, it can be observed from 1(a) that the width of the tracks is only ～100 H9262m, indi- that there is a presence of hatch overlaps of ～10 H9262m. strategy of using hatch overlaps reduces the porosity lev- as well as any possible discontinuities in the sample between hatches, there by resulting in a sound sample with near to density. It has to be noted that there is a rotational shift of 15 deg between Fig. 1(a) and (b) as a result of the mea- sequence. The microstructure is non-uniform with the exhibiting a bimodal distribution. This is the result of the overlap – core morphology typical for the SLM samples , consisting of a fine microstructure along the track over- (marked as (1) in Fig. 1(b)) and a coarse microstructure the track cores (marked as (2) in Fig. 1(b)). Bright platelets 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 dark areas are rich in Al and the bright platelets are rich in Nd, and Co. The bright platelets have different contrasts suggest- the presence of four different phases in the Al 85 Nd 8 Ni 5 Co 2 SLM samples. The presence of four different phases in the SLM sample was confirmed by XRD. The diffraction pat- (Fig. 2) displays the existence of H9251-Al (cubic, Fm3m) with crystallite size d = 72 nm, along with the three intermetallic AlNdNi 4 (orthorhombic, Cmc2), Al 4 CoNi 2 (cubic, and AlNd 3 (hexagonal, P63mmc). The AlNdNi 4 platelets d = 29 nm) are 2.12 ± 0.34 H9262m in length and 0.41 ± 0.13 H9262m in whereas the Al 4 CoNi 2 (d = 42 nm) and AlNd 3 platelets d = 35 nm) are 2.32 ± 0.48 H9262m and 6.01 ± 0.74 H9262m in length 1.00 ± 0.08 H9262m and 0.89 ± 0.14 H9262m in width, respectively. K.G. Prashanth et al. / Additive Manufacturing 6 (2015) 1–5 3 Fig. 1. SEM images of the Al 85 Nd 8 Ni 5 Co 2 alloy showing (a) the laser tracks and (b) the intermetallic phases with different contrasts embedded in the Al matrix and (d–g) The XRD patterns of the different samples indicate that the phases present in the GAP, AP SLM and HT SLM samples are the same (Fig. 2). Interestingly, the XRD patterns of the AP SLM and HT SLM are very similar. The crystallite sizes of the phases in Al ple and sho mate plastic lar that ble at Fig. as-prepared 0.97 ples W tion at observ for the tured temperature tered temperature allo ties be respect allo and SLM HT SLM are 84 nm for Al, 34 nm for AlNdNi 4 , 50 nm for 4 CoNi 2 and 38 nm for AlNd 3 , indicating that the SLM sam- is thermally stable with no significant phase transformation grain growth taking place during the heat-treatment process. The AP SLM specimen tested at room temperature (RT) ws a very high yield strength (YS) of 0.94 GPa and an ulti- compressive strength (UCS) of 1.08 GPa along with 2.45% strain (Fig. 3(a and b)). The HT sample shows simi- properties (YS = 0.81 GPa and UCS = 0.97 GPa), suggesting the microstructure of the SLM material is thermally sta- at high temperatures. Such high strength levels are retained elevated temperatures. For example, UCS of 1.05 GPa and 2. XRD patterns (λ = 0.17889 nm) of the gas atomized powder along with and heat-treated SLM samples. possible Therefore, the Al-based technology e allo ( This with preferential stepped present ture. [12] the bimodal microstructure. High magnification SEM micrograph showing (c) the corresponding EDX images. GPa, and strain of 4.5 and 8.5% are observed for the sam- tested at 373 K and 423 K, respectively (Fig. 3(a and b)). ith further increase of the testing temperature, the deforma- exceeds 20%, where the measurements were stopped. Even a testing temperature of 573 K, a high UCS of ～0.50 GPa is ed, which indicates that the present material can be used high strength applications at high operating temperatures. Although, intense scientific research has been focused on development of high-strength amorphous and nanostruc- Al-based alloys, only a few reports deal with the high mechanical properties of these materials. The sin- nano-crystalline Al–Fe material [19] displays the best high properties among the published works on Al-based ys. These results are compared with the mechanical proper- of the present Al–Nd–Ni–Co SLM alloy in Fig. 3(c). It can observed that the present alloy has similar strength levels with to the Al–Fe alloy at all test temperatures. The Al–Fe y was produced by spark plasma sintering, where the size shape of the component is restricted, whereas processing by permits the production of parts having theoretically any geometry with minimized need for post processing. additive manufacturing offers the possibility to tailor shape and corresponding properties of these high-strength parts to meet specific requirements, which renders this 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 Al 85 Nd 8 Ni 5 Co 2 y. The fracture takes place in a stepped morphology Fig. 4(a)), similar to the Al–12Si alloy prepared by SLM [12]. can be ascribed to the bimodal microstructure (Fig. 3(b)) fine grains along the track-overlaps, which may act as a path for crack propagation, leading to the observed morphology. The superior properties observed in the alloy can be attributed to the composite-like microstruc- The fine microstructure, characteristic of the SLM process , leads to the presence of fine H9251-Al phase surrounded by 4 K.G. Prashanth et al. / Additive Manufacturing 6 (2015) 1–5 Fig. 3. (a) Compressive stress–strain curves of the as-prepared (AP) and heat-treated (HT; 4 h at 723 K) SLM samples tested at room temperature (curves 1 and 2) along with the AP SLM samples tested at different temperatures, (b) corresponding mechanical data and (c) comparison of the present results with the nano-crystalline Al–Fe alloy produced by spark plasma sintering [19]. the intermetallic phases. Since the microstructure is developed from a rapid solidification process, it is expected to have an inherently strong bonding between the Al matrix and the inter- metallic reinforcements, aiding an improved and effective load transfer along the interface. This concept leads to the interfacial strengthening mechanism in the present alloy both at RT and at high temperatures. The hardness of the intermetallic phases evaluated from the nano-indentation measurements are 2.84 ± 0.08 GPa for AlNd 3 , 3.74 ± 0.10 GPa for Al 4 CoNi 2 and 5.45 ± 0.09 GPa for AlNdNi 4 ; which is much higher than pure H9251-Al (0.33 ± 0.03 GPa). Hence, during RT deformation, the cracks are expected to develop along the H9251-Al phase and to propagate with further loading. However, the intermetallic platelets act as obstacles which leads to either the arrest or deflection of the cracks, arrest, further deformation proceeds through the initiation of new cracks leading to crack multiplication (Fig. 4(b)). On the other hand, crack deflection suggests that the effective mean crack path is increased leading to appreciable deformation in the material [20]. Generally, at high temperatures the dislocation movement is accelerated and the magnitude of the Peierls stress is drastically reduced leading to low strength of the material [20]. However, in the present alloy, the Al matrix is surrounded by the intermetallic reinforcement, which may confine the dislocation movement along the grain boundaries according to the confinement theory, which leads to further strengthening of the material at elevated temperatures [21,22]. All of the mech- anisms: interfacial strengthening, crack arrest and initiation of new cracks, crack deflection phenomena and confinement phenomena Fig. arrest as observed from Fig. 4(b and c). In case of the crack 4. Fracture surface after compression tests showing (a) stepped morphology, (b) and crack deflection by the intermetallic phases in the Al 85 Nd 8 Ni 5 Co 2 alloy during (Fig. 4(d)) operate simultaneously leading to multiple cracks and (c) crack deflection. (d) Schematic illustrating the crack compression test. K.G. Prashanth et al. / Additive Manufacturing 6 (2015) 1–5 5 superior room temperature as well as high temperature compressive strengths. 4. Summary Highly dense high strength Al 85 Nd 8 Ni 5 Co 2 alloy has been successfully prepared by SLM. The alloy exhibits a com- posite microstructure with the intermetallic phases AlNdNi 4 , Al 4 CoNi 2 and AlNd 3 dispersed in the Al-matrix. The inter- metallic phases are in the form of platelets and their width ranges in the sub-micron regime. The AP SLM and HT SLM samples show a UCS of 1.08 GPa and 0.97 GPa with ～2.5% strain at RT. The high temperature compression tests reveal that the present system can retain their high strength due to the composite-like microstructure and confinement phenomena, where the grain coarsening and the accelerated mobility of the dislocations at high temperatures are retarded by the intermetallic phases. Addi- tionally, interfacial strengthening, the crack arrest and crack deflection mechanisms also contribute to the superior strength observed in the Al 85 Nd 8 Ni 5 Co 2 alloy. The present results indi- cate that SLM is one of the best options to produce high strength, thermally stable, dense and near-net shaped Al-based alloys. References [7] Surreddi KB, Scudino S, Sakaliyska M, Prashanth KG, Sordelet DJ, Eckert J. Crystallization behavior and consolidation of gas- atomized Al 84 Gd 6 Ni 7 Co 3 glassy powder. J Alloys Compd 2010;491: 137–42. [8] Wang Z, Tan J, Sun BA, Scudino S, Prashanth KG, Zhang WW, et al. 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