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Microstructure and mechanical properties of Cu-modified AlSi10Mg fabricated by Laser-Powder Bed Fusion
Materialia 9 (2020) 100590
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Microstructure and mechanical properties of Cu-modified AlSi10Mg fabricated by Laser-Powder Bed Fusion X. Garmendia∗, S. Chalker, M. Bilton, C.J. Sutcliffe, P.R. Chalker School of Engineering, University of Liverpool, Liverpool L69 3GH, UK
a r t i c l e
i n f o
Keywords: Aluminium alloys Laser-Powder Bed Fusion Grain refinement Metal matrix composites (MMCs) Strengthening mechanisms
a b s t r a c t This influence of surface modification of powder feedstocks for additive manufacturing is assessed. Commercial AlSi10Mg powder was coated with 1 wt% copper using a copper formate precursor, which was reduced to metallic copper after vacuum heat treatment. Consolidated samples were produced using Laser-Powder Bed Fusion (L-PBF). Both the untreated and coated powders produced materials which exhibited high preferential crystal orientation, with grains aligned parallel to the build direction. The microstructure of the consolidated materials was characterized by SEM and TEM. The cells consisted of a 𝛼-Al matrix with copper particles decorating the intercellular region and eutectic silicon homogeneously dispersed in the cell boundary. The effect of the T6 heat treatment on the mechanical properties of the L-PBF-processed materials was a decrease in the strength. The specimens fabricated from coated powder exhibited an increased ultimate tensile strength (461 ± 0.2 MPa) and 12% elongation at break compared to the AlSi10Mg parts (346 ± 1.0 MPa) and 10% elongation. The study establishes the feasibility of modifying the composition of L-PBF-processed materials via the coating of powder feedstocks and that the coating additions can be exploited to enhance mechanical properties of parts after heat treatment.
1. Introduction Recently, Laser-Powder Bed Fusion (L-PBF) has become an emerging Additive Manufacturing (AM) technique, due to its ability to produce complex geometry components in a near-net-shape manner. Aluminium and its alloys have been widely used in AM on account of their low density, low cost, high thermal conductivity and good mechanical properties [1]. Amongst these, one of the most widely used alloys is the precipitation – hardenable hypoeutectic AlSi10Mg. Other Al alloys that have been assessed for AM processing are the aluminium AW6061 (Al–Mg–Si–Cu) [2,3] and AW-2139 (Al–Cu–Mg–Ag) [4], however the range of aluminium-based powder feedstocks for laser-powder bed fusion (L-PBF) processes is relatively limited. This is largely because laser melting involves highly focused energy input, resulting in high local-temperature gradients and fast cooling-rates (104 –106 °K/s) [5]. These factors lead to a totally different solidification behaviour and microstructure compared with that observed from casting. For example, the strength of hardenable alloys is defined by the size and distribution of precipitates formed, however the highly localised melting and resolidification that occurs during the additive manufacturing of metallic alloy powders, precludes the formation of precipitates e.g. Mg2 Si, normally found in castings of Al–Si–Mg alloys [6]. Instead, eutectic Si is found segregated at the cell boundaries, in the as-fabricated condition
∗
[7]. Subsequent solution heat treatment, followed by water quenching and artificial aging, causes the silicon to diffuse, forming particles that increase in size with temperature [8]. When the alloy was processed by preheating the build platform at 300 °C and subsequently solution heattreating at 525 °C for 6 h, the grains coarsened with increasing duration of treatment, forming silicon particulates in the 𝛼-Al matrix. Subsequent artificial ageing at 165 °C for 7 h, caused the Si-particles to coarsen and needle-like Mg2 Si precipitates to arise [9]. From high-cycle fatigue measurements, it was concluded that the combination of 300 °C buildplatform heating and T6 heat treatment served to increase the fatigue resistance and mitigate the effects of build-scan differences on fatigue life. Despite the complexity of materials-related issues that can arise from laser powder bed fusion processing, additive manufacturing allows the agile-manufacture of complex components without using sacrificial moulds, making it an attractive technology for aerospace, automotive and biomedical industries, amongst others [10]. To address the paucity of commercial aluminium alloys [11,12], strategies for tailoring the properties of additively manufactured parts has included the use of mixtures of commercial powder feedstocks or the use of powder-additives [13,14]. Recently, gas-atomisation was used to synthesise an AlSi10Mg composite, decorated with TiB2 nano-powder for Laser-Powder Bed Fusion [15]. The addition of 7 wt%, nano-TiB2 was found to increase the laser absorptivity of the AlSi10Mg powder by approximately 50%. The
Corresponding author. E-mail address: [email protected] (X. Garmendia).
https://doi.org/10.1016/j.mtla.2020.100590 Received 15 November 2019; Accepted 10 January 2020 Available online 14 January 2020 2589-1529/© 2020 Acta Materialia Inc. Published by Elsevier Ltd. This is an open access article under the CC BY license. (http://creativecommons.org/licenses/by/4.0/)
X. Garmendia, S. Chalker and M. Bilton et al.
Materialia 9 (2020) 100590
Table 1 Chemical composition of AlSi10Mg powder. Element Nominal composition (wt%)
Min. Max.
Al
Cu
Fe
Mg
Mn
Ni
Pb
Si
Sn
Ti
Zn
Bal Bal.
0 0.05
0 0.25
0.25 0.45
0 0.1
0 0.05
0 0.02
9 11
0 0.02
0 0.15
0 0.1
additively manufactured material exhibited a texture-less fine-grained microstructure, with a uniform distribution of TiB2 along the grain and cell boundaries and nano-Si rod-like precipitates inside the cells. The composite had a high tensile strength of 530 ± 16 MPa, which was attributed to Hall–Petch, load-bearing and Orowan strengthening mechanisms. The enhanced ductility of ~15.5 ± 1.2% was ascribed to grain boundary modification induced by the nano-TiB2 content and increased dislocation plasticity in the bulk of the grains by nano-Si precipitates. Zhang et al. [16] demonstrated the reduction of the hot cracking phenomena on the Al–Cu–Mg (2xxx) alloy, which was achieved by the addition of 2 wt% Zr particles blended with the Al–Cu–Mg feedstock powder. The addition of elemental zirconium was reported to promote the formation of Al3 Zr precipitates and cause a refined grain microstructure, which prevented the formation and propagation of cracks. The Zr-treated AlSi10Mg parts showed an increased yield strength of 446 ± 4.3 MPa, compared with 253 ± 9.8 MPa for the untreated material. The ultimate tensile strength (UTS) of the Zr-treated alloy was 451 ± 3.6 MPa, an increase of ~16% more for the untreated alloy. However, the elongation at break was reduced from 6 ± 1.6% for the starting powder to 2.7 ± 1.1% for the Zr-treated alloy. From the previous research, it is evident that the surface modification of cast alloy powders can be exploited to enhance the properties of AM processed parts. Furthermore, the addition of copper to AlSiMg alloys is beneficial to their performance. For example, the influence of Cu content on ageing behaviour of AlSiMgCu cast alloys has been previously reported [17,18]. In this paper, the effect of copper-enrichment of AlSi10Mg from a powder surface treatment is addressed. The rationale for this work is to establish how copper from the coated Al–Si alloy is incorporated into selective laser melted material and what its influence is on the mechanical properties and microstructure of the consolidated parts. 2. Experimental 2.1. Powder synthesis AlSi10Mg powder, supplied by LPW Technology Ltd., the composition of which is shown in Table 1 and with a size range of 20–60 μm was used as a starting material. The powder was coated with 1 wt% copper using a copper formate – methanol solution. Copper formate was chosen, as it can be thermally decomposed at low temperature to yield elemental copper [19]. Copper formate tetrahydrate (109.8 g, 98%, Alfa Aesar) was dissolved in a mixture of dry methanol (468.2 g, 99.8% anhydrous methanol, Sigma Aldrich) and octylamine (126.3 g, 99%, Aldrich). The octylamine was added to promote solubility. 3 kg of AlSi10Mg powder (hereafter referred to as Al), was immersed in the solution, mixed thoroughly and allowed to dry, by evaporation in an extracted fume hood. Once dry, trays of powder (750 g) were covered in aluminium foil and evacuated for one hour before heating treatment. The powder was slowly heated to drive off the octylamine at 140–150 °C and decompose the formate tetrahydrate complex at 220 °C. The temperature was held at this temperature for one hour before cooling to room temperature under vacuum. The treated powder (hereafter referred to as Al–Cu) was passed through a 75 μm sieve and then heated again to 220 °C in a vacuum oven, to remove any residual volatiles. The material was then sieved a second time. The average copper content of the Al–Cu powder was determined to be 0.97 ± 0.02 wt% by X-ray fluorescence (XRF). The flow characteristics were evaluated for the treated and untreated
Table 2 Summary of flowability and moisture content of the powders. Powder
Flowability (s)
Moisture (ppm)
AlSi10Mg Al–Cu
71.6 ± 0.5 76.4 ± 0.9
118.9 598.5
powders using Hall and Carney funnels following the ASTM B527-15 standard [20]. In both cases, 50 g samples of powders were used and yielded flow times of 76.4 ± 0.9 s for the Al–Cu powder and 71.6 ± 0.5 s for the Al powder. The native oxide present on aluminium powders can form hydroxides (e.g. Al(OH)3 ) which can thermally decompose to form Al2 O3 and water vapour. To assess the influence of the hydroxide, the residual moisture level of the powders was analysed by Karl Fischer Titration method, using a Metrohm 860 KF Thermoprep instrument. Here the term “moisture” relates to the Al(OH)3 content of the powder, as well as any adsorbed water. After completion of the L-PBF-processing, the moisture/hydroxide content of the Al–Cu powder was 598.5 ppm or 0.06%. This was five times higher than that of pristine Al-powder, of 118.9 ppm or 0.012%. The flow and moisture results of the powders are summarised in Table 2. 2.2. Laser-Powder Bed Fusion process The powders were processed using a Renishaw AM400 system equipped with a 400 W SPI fibre laser and a wavelength of 1.06 μm, which is used in continuous mode. Samples were fabricated in a reduced build volume platform of 78 × 78 × 55 mm, and the chamber was flooded with argon to ensure an oxygen level below 1000 ppm. The build platform was not preheated in this study. During the scanning process, an inert gas flow crosses the platform from right to left, removing the unwanted spattered particles from the powder bed. A meander scan strategy was selected to build the parts, which rotates the laser scanning vector 67° after each scan. The main process parameters involved in the laser L-PBF process are the laser power, hatching distance, point distance, exposure time, and layer thickness. Values for these were 200 W, 0.13 mm, 80 μm, and 140 μs, respectively. For this study, the layer thickness was kept constant at 30 μm. The process parameters giving the highest density for uncoated AlSi10Mg were determined first and were used as the starting point for evaluating the conditions for the coated powder. Samples (10 × 10 × 10 mm3 ) were built, mounted in epoxy resin, and polished down to 0.020 μm using colloidal silica suspension. The density was assessed using 20 optical images, which were tiled and image-processed. For the mechanical testing, builds were made (Fig. 1) containing hexagonal section bars for tensile specimens. Three cubes were printed on top of the bars for porosity level evaluation of the XY, XZ and YZ cross-sections. An additional cube (“4”) was used for XRD measurement and the pyramidal structure (“5”) was used for a hardness test. On selected samples, a Solution Heat Treatment (SHT) was applied, followed by artificial aging (AA) and heat treatment (T6). These were performed following the ASTM F3318-18 standard [21]. The solution heat treatment was carried out at 530 °C for 6 h followed by water quenching and subsequent artificial aging at 160 °C for 6 h. 2.3. Materials characterisation The microstructure of the samples was analysed using a JEOL JSM 6610 Scanning Electron Microscope (SEM) equipped with an Oxford
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Materialia 9 (2020) 100590
using beam conditions of 20 kV, 5.5 nA and a 12 mm working distance. X-ray diffraction (XRD) measurements were performed using a Rigaku Miniflex with Cu K𝛼 radiation operated at 30 kV and 15 mA, with a step size of 0.01° (2𝜃) and a scanning speed of 0.03°/min. Consolidated material was analysed using scanning transmission electron microscopy (STEM) on an aberration-corrected JEOL 2100FCs S/TEM operating at 200 kV. The electron-transparent lamellae were prepared using the FEI Helios 600i dual beam focused ion beam (FIB) instrument, with a size thickness of ≈50–100 nm. Specimens were prepared by the lift out method and low energy (2 kV) polishing was applied as a final step to help removing potential Ga-ion induced surface amorphisation. On samples coated with copper, Electron Energy Loss Spectroscopy (EELS) was performed using the JEOL 2100FCs equipped with a GIF Quantum 693 electron energy filter. Images were processed using Gatan Digital Micrograph, applying periodic masks on FFT images. 2.4. Mechanical properties
Fig. 1. View of the build geometry used for the tensile, porosity, XRD and hardness specimens.
INCA X-act energy-dispersive X-ray spectroscopy (EDX). Fractographs of the tensile tested surfaces were obtained using the same microscope. Samples were etched using Keller’s reagent to reveal the microstructure. Texture measurements were performed on a FEI Helios Nanolab 600i Focused Ion Beam (FIB) instrument equipped with an EDAX-EBSD system. After mechanical polishing, samples were prepared using a Gatan Precision Ion Polishing System (PIPS). Samples were polished for 1 h, under a 5 kV Ar ion beam at a 5° incident angle. High resolution electron backscattered diffraction (EBSD) maps were collected using a step size of 80 nm in areas of fine-grained material (10k times magnification), and 250 nm steps were used to map large areas of coarse-grained material at lower (750x magnification). All EBSD maps were recorded
For mechanical properties evaluation, samples were manufactured with their longitudinal axes aligned parallel to the build plate, and three cylindrical “dog-bone” tension specimens were extracted from each build. Specimens built perpendicular to the build plate were not considered in this study. Tensile tests were conducted using 6 mm diameter and 24 mm gauge length specimens, machined in accordance with ASTM standard E8/E8M [22]. An Instron 5984 universal testing machine equipped with a 100 kN load cell was used, applying a displacement rate of 0.005 mm/mm/min up to the yield point and of 0.05 mm/mm/min until fracture. Hardness Vickers measurements were performed using a weight of 500 g in a Wilson VH3100 machine, in accordance with ASTM E92-17 standard [23]. A total of 12 indents were taken from the polished XY cross section. 3. Results 3.1. Al and Al–Cu powders Fig. 2 shows backscattered electron micrographs obtained from (a) the untreated and (b) Cu-treated aluminium powders. Fig. 2. Backscattered SEM images obtained at 20 kV of AlSi10Mg powder, (a) as received; and (b) after treatment with copper formate complex in methanol solution and vacuum-drying at 220 °C; (c) magnified view of copper-coated particle; (d) EDX map of Cu distribution.
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Materialia 9 (2020) 100590
same area as (c), indicating that the bright particulates correlate with the presence of copper. The presence of copper was also demonstrated by XRD, where peaks corresponding to face-centred cubic (fcc) Cu were evident at 2𝜃 = 43.3° and 2𝜃 = 50.43°, as shown in Fig. 3. 3.2. Microstructure and texture
Fig. 3. XRD patterns of Al and Al–Cu powder with copper peaks highlighted.
Morphology of the Al powder was a mixture of spherical particles and elongated particles and satellites. The backscattered image is sensitive to the atomic number of elements in the samples and the bright particulate distribution shows the copper is discontinuous and of various sizes. The variability of the copper distribution is shown more clearly in Fig. 2(c) indicating that the dusting of copper is predominantly less than 2 μm in diameter and irregularly shaped. Fig. 2(d) shows an energy-dispersive X-ray (EDX) map of the Cu k𝛼1 radiation from the
SEM examination of etched XZ planes were performed to observe the cellular morphology evolution within the melt pool. Fig. 4(a) shows a repetitive fish-scale-like pattern, which has been previously reported by Zhou et al. [24]. The depth of the melt pools varies across the section, with values between one and three times the layer thickness. This is in accordance with the partial re-melting of previously solidified layers and following epitaxial growth of grains that is expected during the L-PBF process. Three different types of cells were identified across the melt pool. A fine cellular microstructure was developed close to the melt pool boundary triple points, MPB1, as shown in Fig. 4(b), while a coarser quasi-equiaxed cell type was formed on standard melt pool boundaries, MPB2, illustrated in Fig. 4(c). MPB1 is developed in the heat affected zone (HAZ) that is generated in the areas where the melt pools are overlapped. The bulk region (BR) represented in Fig. 4(d) was characterized by elongated cells with long axes parallel to the build direction. The measurements of the cells yielded a cell width of ≈ 300 nm in the elongated cells of the bulk region, while quasi-equiaxed cells at MPB1 and MPB2 showed diameters of ≈ 400 nm and ≈ 600 nm, respectively. It was observed that the cellular structure located at the edges follows a direction normal to the melt pool boundary. Closer observation of the cell structures using STEM, illustrated in Fig. 5, reveals the elemental composition of the bulk and intercellular regions of the L-PBF-processed Al–Cu. Supersaturated 𝛼-Al phase is Fig. 4. SEM image of etched Al–Cu XZ section at (a) high magnification with melt pool boundary highlighted in red and high magnification images of melt pool boundary 1 (MPB1), melt pool boundary 2 (MPB2) and bulk region (BR) in (b), (c) and (d), respectively.
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Materialia 9 (2020) 100590
Fig. 5. TEM HAADF image of the Al–Cu L-PBF-processed sample and corresponding EDX mapping of the main constituent elements. Fig. 6. EBSD orientation maps of L-PBF-processed (a) Al–Cu and (b) Al samples, with their grain size showed in (c) and (d).
observed on the bulk of the cellular structure in the Al K map. The intercellular region is decorated with homogeneously distributed silicon and localized copper particles. A small proportion of oxygen and iron was also found in the cell boundary. The EBSD orientation maps of the L-PBF-processed Al and Al–Cu and their corresponding grain size distribution are shown in Fig. 6, in order to reveal the overall crystallographic texture of the material. Strong texture was developed in Al and Al–Cu material, with a bimodal grain structure composed of elongated and equiaxed grains respectively. Morphologically, the addition of copper did not result in an increase in small equiaxed grains. However, the Al–Cu material promoted an overall
grain size reduction of ≈ 15 μm. In both cases, the grain size increases towards the centre of the melt pool, with long elongated grains crossing multiple melt pools. These grains showed a ⟨001⟩ crystal growth direction along the building direction, i.e., parallel to the build direction. Some of these columnar grains are slightly tilted towards the ⟨101⟩ direction. In contrast, a proportion of small equiaxed grains showed no evidence of any texture. With respect to the grain sizes, columnar grains range between 40–150 μm in length and 5–15 μm in width, while the equiaxed grains share a diameter of approximately 6 μm. Fig. 7 shows the XRD of the L-PBF-processed parts, which do not exhibit the copper peaks from discrete particles seen in the Cu-treated
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Materialia 9 (2020) 100590
1000
Al-Cu Al
900
Intensity (a.u.)
800 700 600 500 400 300 200 100 0
30
40
50
60
70
80
2 (°) Fig. 7. XRD patterns of L-PBF-processed Al and Al–Cu samples. Fig. 9. STEM-BF image of L-PBF-processed Al–Cu sample.
powder (Fig. 3). However, energy dispersive X-ray analysis of the L-PBF-processed Al–Cu samples indicates the presence of copper, revealing that the copper coating has been incorporated within the alloy. This was also confirmed by the EELS spectrum in Fig. 8, showing an energy loss at 951 eV, corresponding to the L2 edge of copper. The image in Fig. 9 shows evidence of dislocation networks in some of the cells of the L-PBF-processed Al–Cu sample. Discrete copper particles were present on the grain boundaries, generating the pile-up of dislocations. 3.3. Mechanical properties An initial density evaluation on three XY plane cross-sections of Al–Cu cubes fabricated using the optimised parameters of Al material was done. Results showed an average density of 99.1 ± 0.2% and, therefore, the optimised parameters of Al were used in the subsequent Al–Cu builds. Density measurements of the three cross-sections were taken from cubes located in top of the tensile bars, for confirmation purposes, the results of which can be seen in Table 3. It can be observed that the density of the confirmation Al–Cu cubes, which are stacked on top of the bars, decreased compared to the initial value of 99.1 ± 0.2%. The different tensile stress-strain curves for the as-built and heattreated specimens are shown in Fig. 10, with their respective data indicated in Table 4. Generally, the samples behaved in a relatively ductile manner, as can be deduced from the elongation at break and high ultimate tensile strength values. The Al–Cu specimens showed very homogeneous UTS values, while the elongation at break values were
Table 3 Density values from the L-PBF-processed Al and Al–Cu build used for tensile test bars. Cross-section Material
XY
XZ
YZ
Al Al for T6 Al–Cu Al–Cu for T6
98.7 99.1 94.4 97.4
99.5 99.3 95.0 96.5
99.8 99.0 94.2 95.3
more scattered. The as-built L-PBF-processed Al specimens had an ultimate tensile strength and elongation at break values of 410 ± 10.1 MPa and 7 ± 0.5%, respectively. These values exceed that of the majority of cast and wrought alloys [1]. This is generally attributed to the very fine microstructure developed upon the rapid solidification that occurs during the L-PBF process. When the Al specimens were subjected to T6 heat treatment, the material ductility was slightly enhanced, from 7.0 ± 0.5% to 9.2 ± 0.5%. On the other hand, the UTS was reduced by 100 MPa, from 410 ± 10.1 MPa to 310 ± 4.4 MPa. The as-built Al–Cu material displayed a UTS of 461.0 ± 0.2 MPa, 51 MPa higher than that of the Al. Also, it exhibited a higher ductility, with an elongation at break 3% higher than that of the Al. In contrast to the Al specimens, the elongation of Al–Cu specimens decreased from 12 ± 0.5% to 9.7 ± 2.7%
Fig. 8. HAADF-STEM image of (a) Al–Cu sample with the incident beam location highlighted with the red marker and, (b) its corresponding EELS spectrum.
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Materialia 9 (2020) 100590
Table 4 Mechanical properties of L-PBF-processed Al and Al–Cu.
AlSi10MgAs built AlSi10MgT6 Al–CuAs built Al–CuT6
UTS (MPa)
Yield Strength (MPa)
Young Modulus (GPa)
Elongation at break (%)
410 310 461 346
267 244 264 254
70.8 (±2) 72.4 (±2.4) 73 (±1.8) 76.3 (±7)
7 (±0.5) 9.2 (±0.5) 12 (±0.5) 9.7 (±2.7)
(±10.1) (±4.4) (±0.2) (±1)
(±3.7) (±3.1) (±1.8) (±13.74)
for the as-built and heat-treated Al–Cu was significantly higher than that of Al. Indent 9 of the Al–Cu T6 sample was excluded from the data as it was compromised by local porosity, rendering it unrepresentative of the test material. In general, the reduction of hardness after T6 was smaller on the coated powder material. Measurements on as-built consolidated material yielded values of 117.3 ± 3 HV0.5 and 110.3 ± 5.8 HV0.5 for Al and Al–Cu samples, respectively. However, when samples were subjected to T6 heat treatment, the Al–Cu sample showed an average hardness of 101.6 ± 14.2 HV0.5 , while the Al sample exhibited a value of 90.6 ± 2.3 HV0.5 . The hardness results are concordant with the UTS results obtained after tensile testing, excluding the as-built Al–Cu values, which would be expected to be greater than those observed for Al. 4. Discussion 4.1. Microstructure and texture
Fig. 10. Engineering tensile stress-strain curves of (a) as build and (b) heat treated Al and Al–Cu L-PBF-processed samples.
when subjected to heat treatment. The UTS reduction for heat treated Al–Cu was approximately 25%, the same in the case of Al. Tensile tested specimen fractographs, shown in Fig. 11, were used to reveal the brittle or ductile nature of the samples, which is associated with the sample elongation capacity. Macroscopically, the fracture surfaces showed little plastic deformation, however the microstructural appearance was fibrous, which is associated with ductile fracture. No necking was seen macroscopically, which is consistent with the fracture stress being equal to the ultimate tensile strength in Fig. 10. Failure was initiated at a surface defect, such as pore or a crack, which propagated across the fracture plane, perpendicular to the load direction. Some voids with un-melted powder were observed in the fractographs, as well as shallow dimples with a diameter range of ≈ 13–66 μm. Hardness test results showed different micro hardness values (HV0.5 ) for both materials, as shown in Fig. 12. The scatter of hardness values
The processability of metallic powders by Laser-Powder Bed Fusion is influenced by factors such as powder morphology, particle size distribution, flowability, chemical composition, moisture content, and laser absorptivity. The attainment of a homogeneous powder layer for the L-PBF process is determined by the particle size distribution, powder flowability, and packing density. The powder coating method employed here did not significantly modify the flow behaviour of the powders, as it does not alter the powder morphology or particle size distribution. During processing, a layer of powder is deposited over previously solidified material, which is partially re-melted by the laser. The macrostructure of the consolidated material shows fish-scale-like features, remnant of melt pools with heights up to three times the layer thickness. The EBSD results demonstrated the epitaxial growth of fine grains in the direction parallel to the build direction. This resulted in long elongated grains that cross several cells and melt pools. As a consequence, the epitaxially grown grains share their crystal orientation. This general texture of the materials contributed to the undesired anisotropy of the material. Some quasi-equiaxed grains were observed in both materials, while the Al–Cu samples showed an overall grain size reduction. The precipitation sequence of Al–Si–Mg alloys usually starts with the supersaturated solid solution trapping solute atoms in solution within the solid. These solute atoms adopt the form of metastable Guinier–Preston (GP) zones. These are decomposed into needle-like 𝛽’’ phases, which then transforms into the semi-coherent rod-shaped 𝛽’Mg2 Si phase. Finally, the latter is transformed into a stable, incoherent, fcc 𝛽-Mg2 Si phase (in the form of plates) [25]. As previously reported, aluminium cells are the first to solidify and subsequently, the liquid of approximately eutectic composition is contained in the intercellular region [26]. The uncoated samples present an aluminium matrix with fine eutectic silicon populating the intercellular regions. Upon material solidification, copper is pushed towards the intercellular region, were the eutectic Al–Si is formed. This explains why most of the copper is concentrated in the cell boundary, although some copper-enriched precipitates were visible by TEM on the 𝛼-Al matrix. The precipitation sequence in Al–Cu alloys is quite complex, and it occurs in a series of reactions leading to metastable precipitates before the formation of the equilibrium phase 𝜃-Al2 Cu is achieved. Some needle-shaped iron-rich precipitates were also found in the matrix of coated and uncoated material cells, particularly after T6 heat treatment. These precipitates are undesirable, forming elongated brittle 𝛽-Al5 FeSi phases. The heat
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Materialia 9 (2020) 100590
Fig. 11. Fracture surfaces of L-PBFprocessed (a) Al, (b) T6 Al, (c) Al–Cu, (d) T6 Al–Cu.
4.2. Mechanical properties
Fig. 12. Hardness Vickers (HV0.5 ) values for as-built and heat-treated Al and Al–Cu samples.
treatment promotes the formation of silicon platelets, which have a brittle diamond crystal structure. Consequently, these precipitates are detrimental for mechanical properties since they act as sites for crack nucleation under cyclic loading [27].
According to previous studies, there are four main strengthening mechanisms in metal matrix composites: dislocation strengthening, Orowan strengthening, grain boundary reduction strengthening and load-bearing strengthening [15]. The aluminium alloys processed in this study showed UTS values higher than conventional cast A360. This is attributed to the fine grain and eutectic cellular structure developed as a result of the rapid solidification during the L-PBF process. In addition, the dislocation entanglement shown by the TEM contributes to the material strengthening by dislocation motion hindering. The higher porosity level of the L-PBF-processed Al–Cu specimens, may be attributable to unoptimized laser parameters or the evolution of volatile gases resulting from the coating process. Nevertheless, the additional porosity did not appear to detrimentally influence the mechanical properties of the as-built or heat treated material. In fact, both the as-built and T6-treated Al–Cu material exhibited higher UTS and elongation values than the AlSi10Mg counterparts. This increase in the ultimate tensile strength of the Al–Cu specimens do not show the expected correlation with the HV0.5 results. This is attributed to the higher porosity shown by those specimens, which leads to a higher scattering in the hardness results. The probability of an indent reaching a zone with a pore and, therefore, yielding a low hardness value, are higher in the Al–Cu specimens. We attribute, the origin of the UTS enhancement in the Al–Cu case to a combination of Hall–Petch and Orowan strengthening mechanisms. One factor is the grain refinement shown in the EBSD analysis and promoted by the copper addition, which contributes to the Hall–Petch strengthening. A second factor is the copper precipitates
X. Garmendia, S. Chalker and M. Bilton et al.
located both on the intercellular and matrix sites which contribute to the dislocation motion hindering, i.e., Orowan strengthening. The overall reduction in hardness and tensile strength when the specimens were subjected to aging is promoted by the nucleation and growth of the silicon precipitates of brittle nature, in the so-called Ostwald ripening process. The size and volume fraction of those precipitates influence the strength of the material, following the Orowan equation. The bigger the precipitate size is, the lower the strength of the material will be. After the same heat treatment conditions were applied, the Al–Cu specimens yielded a higher ultimate tensile strength. The size distribution of the silicon precipitates was smaller in the heat-treated Al–Cu material, which leads the higher UTS values. 5. Conclusions In this study, AlSi10Mg powder was coated with 1 wt% copper and samples were fabricated using Laser-Powder Bed Fusion. The microstructure and mechanical properties of the uncoated and coated materials were evaluated. The processability of copper coated AlSi10Mg powder has been demonstrated. The copper coated specimens showed an increase of 51 MPa in the ultimate tensile strength when compared to AlSi10Mg. This is attributed to the grain size reduction and consequent Hall–Petch strengthening, together with Oworan strengthening promoted by the copper-enriched precipitates. Microstructurally, the material consists of an 𝛼-Al cell matrix with eutectic Si decorating the cell boundaries. Furthermore, copper is segregated in the cell boundary. Both materials exhibit silicon platelets and needle-like iron-rich precipitates when subjected to peak hardening. However, the copper addition causes a reduction in the size of the silicon precipitates after heat treatment. The study establishes the feasibility of modifying the composition of L-PBF-processed materials via the coating of powder feedstocks and that the coating additions can be exploited to enhance mechanical properties of parts after heat treatment. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgments The authors gratefully acknowledges financial support from Renishaw plc and the Engineering and Physical Sciences Research Council (EPSRC) Centre for Doctoral Training in Additive Manufacturing and 3D Printing. We also acknowledge funding support from the Innovate UK project FLAC (102665) and EPSRC project EP/N017773/1. The authors would also like to thank Dr Karl Dawson for helpful discussions regarding the electron microscopy and texture analysis. References [1] K. Kempen, L. Thijs, J. Van Humbeeck, J.-.P. Kruth, Mechanical properties of AlSi10Mg produced by selective laser melting, Phys. Proc. 39 (2012) 439–446, doi:10.1016/J.PHPRO.2012.10.059. [2] B.A. Fulcher, D.K. Leigh, T.J. Watt, Comparison of AlSi10Mg and Al 6061 processed through DMLS, in: B.A. Fulcher, D.K. Leigh, T.J. Watt (Eds.), Harvest Technologies, Department of Mechanical Engineering, The University of Texas at Austin, Belton, TX 76513, Austin, TX 78712, 2014, pp. 404–419. [3] E. Louvis, P. Fox, C.J. Sutcliffe, Selective laser melting of aluminium components, J. Mater. Process. Technol. 211 (2011) 275–284, doi:10.1016/j.jmatprotec.2010.09.019.
Materialia 9 (2020) 100590 [4] C. Brice, R. Shenoy, M. Kral, K. Buchannan, Precipitation behavior of aluminum alloy 2139 fabricated using additive manufacturing, Mater. Sci. Eng. A 648 (2015) 9–14, doi:10.1016/J.MSEA.2015.08.088. [5] Q. Jia, P. Rometsch, S. Cao, K. Zhang, X. Wu, Towards a high strength aluminium alloy development methodology for selective laser melting, Mater. Des. 174 (2019) 107775, doi:10.1016/J.MATDES.2019.107775. [6] A. Hadadzadeh, B.S. Amirkhiz, M. Mohammadi, Contribution of Mg2Si precipitates to the strength of direct metal laser sintered AlSi10Mg, Mater. Sci. Eng. A. 739 (2019) 295–300, doi:10.1016/J.MSEA.2018.10.055. [7] N.T. Aboulkhair, C. Tuck, I. Ashcroft, I. Maskery, N.M. Everitt, On the precipitation hardening of selective laser melted AlSi10Mg, Metall. Mater. Trans. A 46 (2015) 3337–3341, doi:10.1007/s11661-015-2980-7. [8] D. Herzog, V. Seyda, E. Wycisk, C. Emmelmann, Additive manufacturing of metals, Acta Mater. 117 (2016) 371–392, doi:10.1016/J.ACTAMAT.2016.07.019. [9] E. Brandl, U. Heckenberger, V. Holzinger, D. Buchbinder, Additive manufactured AlSi10Mg samples using selective laser melting (SLM): microstructure, high cycle fatigue, and fracture behavior, Mater. Des. 34 (2012) 159–169, doi:10.1016/J.MATDES.2011.07.067. [10] D. Buchbinder, H. Schleifenbaum, S. Heidrich, W. Meiners, J. Bültmann, High power selective laser melting (HP SLM) of aluminum parts, Phys. Proc. (2011), doi:10.1016/j.phpro.2011.03.035. [11] E.O. Olakanmi, R.F. Cochrane, K.W. Dalgarno, A review on selective laser sintering/melting (SLS/SLM) of aluminium alloy powders: processing, microstructure, and properties, Prog. Mater. Sci. 74 (2015) 401–477, doi:10.1016/J.PMATSCI.2015.03.002. [12] I.E. [Ames L. and I.S.U. Anderson Ames, IA (United States)], E.M.H. [Ames L. and I.S.U. White Ames, IA (United States)], R. [Oak R.N.L. (ORNL) Dehoff Oak Ridge, TN (United States)], Feedstock powder processing research needs for additive manufacturing development, (2018). doi:10.1016/j.cossms.2018.01.002. [13] K. Bartkowiak, S. Ullrich, T. Frick, M. Schmidt, New developments of laser processing aluminium alloys via additive manufacturing technique, Phys. Proc. 12 (2011) 393–401. [14] M.L. Montero, R. Mertens, B. Vrancken, X. Wang, B. Van Hooreweder, J. Kruth, J. Van Humbeeck, Changing the alloy composition of Al7075 for better processability by selective laser melting, J. Mater. Process. Tech. 238 (2016) 437–445, doi:10.1016/j.jmatprotec.2016.08.003. [15] X.P. Li, G. Ji, Z. Chen, A. Addad, Y. Wu, H.W. Wang, J. Vleugels, J. Van Humbeeck, J.P. Kruth, Selective laser melting of nano-TiB2 decorated AlSi10Mg alloy with high fracture strength and ductility, Acta Mater. 129 (2017) 183–193, doi:10.1016/J.ACTAMAT.2017.02.062. [16] H. Zhang, H. Zhu, X. Nie, J. Yin, Z. Hu, X. Zeng, Effect of zirconium addition on crack, microstructure and mechanical behavior of selective laser melted Al–Cu–Mg alloy, Scr. Mater. 134 (2017) 6–10, doi:10.1016/j.scriptamat.2017.02.036. [17] G. Wang, Q. Sun, L. Feng, L. Hui, C. Jing, Influence of Cu content on ageing behavior of AlSiMgCu cast alloys, Mater. Des. 28 (2007) 1001–1005, doi:10.1016/J.MATDES.2005.11.015. [18] Y.J. Li, S. Brusethaug, A. Olsen, Influence of Cu on the mechanical properties and precipitation behavior of AlSi7Mg0.5 alloy during aging treatment, Scr. Mater. 54 (2006) 99–103, doi:10.1016/j.scriptamat.2005.08.044. [19] A.K. Galwey, D. Jamieson, M.E. Brown, Thermal decomposition of three crystalline modifications of anhydrous copper(II) formate, J. Phys. Chem. 78 (1974) 2664– 2670, doi:10.1021/j100619a006. [20] ASTM International, Standard test method for tap density of metal powders and compounds, ASTM Stand. B527 (2015) 15–18, doi:10.1520/B0527-15.2. 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Microstructure and mechanical properties of Cu-modified AlSi10Mg fabricated by Laser-Powder Bed Fusion X. Garmendia∗, S. Chalker, M. Bilton, C.J. Sutcliffe, P.R. Chalker School of Engineering, University of Liverpool, Liverpool L69 3GH, UK
a r t i c l e
i n f o
Keywords: Aluminium alloys Laser-Powder Bed Fusion Grain refinement Metal matrix composites (MMCs) Strengthening mechanisms
a b s t r a c t This influence of surface modification of powder feedstocks for additive manufacturing is assessed. Commercial AlSi10Mg powder was coated with 1 wt% copper using a copper formate precursor, which was reduced to metallic copper after vacuum heat treatment. Consolidated samples were produced using Laser-Powder Bed Fusion (L-PBF). Both the untreated and coated powders produced materials which exhibited high preferential crystal orientation, with grains aligned parallel to the build direction. The microstructure of the consolidated materials was characterized by SEM and TEM. The cells consisted of a 𝛼-Al matrix with copper particles decorating the intercellular region and eutectic silicon homogeneously dispersed in the cell boundary. The effect of the T6 heat treatment on the mechanical properties of the L-PBF-processed materials was a decrease in the strength. The specimens fabricated from coated powder exhibited an increased ultimate tensile strength (461 ± 0.2 MPa) and 12% elongation at break compared to the AlSi10Mg parts (346 ± 1.0 MPa) and 10% elongation. The study establishes the feasibility of modifying the composition of L-PBF-processed materials via the coating of powder feedstocks and that the coating additions can be exploited to enhance mechanical properties of parts after heat treatment.
1. Introduction Recently, Laser-Powder Bed Fusion (L-PBF) has become an emerging Additive Manufacturing (AM) technique, due to its ability to produce complex geometry components in a near-net-shape manner. Aluminium and its alloys have been widely used in AM on account of their low density, low cost, high thermal conductivity and good mechanical properties [1]. Amongst these, one of the most widely used alloys is the precipitation – hardenable hypoeutectic AlSi10Mg. Other Al alloys that have been assessed for AM processing are the aluminium AW6061 (Al–Mg–Si–Cu) [2,3] and AW-2139 (Al–Cu–Mg–Ag) [4], however the range of aluminium-based powder feedstocks for laser-powder bed fusion (L-PBF) processes is relatively limited. This is largely because laser melting involves highly focused energy input, resulting in high local-temperature gradients and fast cooling-rates (104 –106 °K/s) [5]. These factors lead to a totally different solidification behaviour and microstructure compared with that observed from casting. For example, the strength of hardenable alloys is defined by the size and distribution of precipitates formed, however the highly localised melting and resolidification that occurs during the additive manufacturing of metallic alloy powders, precludes the formation of precipitates e.g. Mg2 Si, normally found in castings of Al–Si–Mg alloys [6]. Instead, eutectic Si is found segregated at the cell boundaries, in the as-fabricated condition
∗
[7]. Subsequent solution heat treatment, followed by water quenching and artificial aging, causes the silicon to diffuse, forming particles that increase in size with temperature [8]. When the alloy was processed by preheating the build platform at 300 °C and subsequently solution heattreating at 525 °C for 6 h, the grains coarsened with increasing duration of treatment, forming silicon particulates in the 𝛼-Al matrix. Subsequent artificial ageing at 165 °C for 7 h, caused the Si-particles to coarsen and needle-like Mg2 Si precipitates to arise [9]. From high-cycle fatigue measurements, it was concluded that the combination of 300 °C buildplatform heating and T6 heat treatment served to increase the fatigue resistance and mitigate the effects of build-scan differences on fatigue life. Despite the complexity of materials-related issues that can arise from laser powder bed fusion processing, additive manufacturing allows the agile-manufacture of complex components without using sacrificial moulds, making it an attractive technology for aerospace, automotive and biomedical industries, amongst others [10]. To address the paucity of commercial aluminium alloys [11,12], strategies for tailoring the properties of additively manufactured parts has included the use of mixtures of commercial powder feedstocks or the use of powder-additives [13,14]. Recently, gas-atomisation was used to synthesise an AlSi10Mg composite, decorated with TiB2 nano-powder for Laser-Powder Bed Fusion [15]. The addition of 7 wt%, nano-TiB2 was found to increase the laser absorptivity of the AlSi10Mg powder by approximately 50%. The
Corresponding author. E-mail address: [email protected] (X. Garmendia).
https://doi.org/10.1016/j.mtla.2020.100590 Received 15 November 2019; Accepted 10 January 2020 Available online 14 January 2020 2589-1529/© 2020 Acta Materialia Inc. Published by Elsevier Ltd. This is an open access article under the CC BY license. (http://creativecommons.org/licenses/by/4.0/)
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Table 1 Chemical composition of AlSi10Mg powder. Element Nominal composition (wt%)
Min. Max.
Al
Cu
Fe
Mg
Mn
Ni
Pb
Si
Sn
Ti
Zn
Bal Bal.
0 0.05
0 0.25
0.25 0.45
0 0.1
0 0.05
0 0.02
9 11
0 0.02
0 0.15
0 0.1
additively manufactured material exhibited a texture-less fine-grained microstructure, with a uniform distribution of TiB2 along the grain and cell boundaries and nano-Si rod-like precipitates inside the cells. The composite had a high tensile strength of 530 ± 16 MPa, which was attributed to Hall–Petch, load-bearing and Orowan strengthening mechanisms. The enhanced ductility of ~15.5 ± 1.2% was ascribed to grain boundary modification induced by the nano-TiB2 content and increased dislocation plasticity in the bulk of the grains by nano-Si precipitates. Zhang et al. [16] demonstrated the reduction of the hot cracking phenomena on the Al–Cu–Mg (2xxx) alloy, which was achieved by the addition of 2 wt% Zr particles blended with the Al–Cu–Mg feedstock powder. The addition of elemental zirconium was reported to promote the formation of Al3 Zr precipitates and cause a refined grain microstructure, which prevented the formation and propagation of cracks. The Zr-treated AlSi10Mg parts showed an increased yield strength of 446 ± 4.3 MPa, compared with 253 ± 9.8 MPa for the untreated material. The ultimate tensile strength (UTS) of the Zr-treated alloy was 451 ± 3.6 MPa, an increase of ~16% more for the untreated alloy. However, the elongation at break was reduced from 6 ± 1.6% for the starting powder to 2.7 ± 1.1% for the Zr-treated alloy. From the previous research, it is evident that the surface modification of cast alloy powders can be exploited to enhance the properties of AM processed parts. Furthermore, the addition of copper to AlSiMg alloys is beneficial to their performance. For example, the influence of Cu content on ageing behaviour of AlSiMgCu cast alloys has been previously reported [17,18]. In this paper, the effect of copper-enrichment of AlSi10Mg from a powder surface treatment is addressed. The rationale for this work is to establish how copper from the coated Al–Si alloy is incorporated into selective laser melted material and what its influence is on the mechanical properties and microstructure of the consolidated parts. 2. Experimental 2.1. Powder synthesis AlSi10Mg powder, supplied by LPW Technology Ltd., the composition of which is shown in Table 1 and with a size range of 20–60 μm was used as a starting material. The powder was coated with 1 wt% copper using a copper formate – methanol solution. Copper formate was chosen, as it can be thermally decomposed at low temperature to yield elemental copper [19]. Copper formate tetrahydrate (109.8 g, 98%, Alfa Aesar) was dissolved in a mixture of dry methanol (468.2 g, 99.8% anhydrous methanol, Sigma Aldrich) and octylamine (126.3 g, 99%, Aldrich). The octylamine was added to promote solubility. 3 kg of AlSi10Mg powder (hereafter referred to as Al), was immersed in the solution, mixed thoroughly and allowed to dry, by evaporation in an extracted fume hood. Once dry, trays of powder (750 g) were covered in aluminium foil and evacuated for one hour before heating treatment. The powder was slowly heated to drive off the octylamine at 140–150 °C and decompose the formate tetrahydrate complex at 220 °C. The temperature was held at this temperature for one hour before cooling to room temperature under vacuum. The treated powder (hereafter referred to as Al–Cu) was passed through a 75 μm sieve and then heated again to 220 °C in a vacuum oven, to remove any residual volatiles. The material was then sieved a second time. The average copper content of the Al–Cu powder was determined to be 0.97 ± 0.02 wt% by X-ray fluorescence (XRF). The flow characteristics were evaluated for the treated and untreated
Table 2 Summary of flowability and moisture content of the powders. Powder
Flowability (s)
Moisture (ppm)
AlSi10Mg Al–Cu
71.6 ± 0.5 76.4 ± 0.9
118.9 598.5
powders using Hall and Carney funnels following the ASTM B527-15 standard [20]. In both cases, 50 g samples of powders were used and yielded flow times of 76.4 ± 0.9 s for the Al–Cu powder and 71.6 ± 0.5 s for the Al powder. The native oxide present on aluminium powders can form hydroxides (e.g. Al(OH)3 ) which can thermally decompose to form Al2 O3 and water vapour. To assess the influence of the hydroxide, the residual moisture level of the powders was analysed by Karl Fischer Titration method, using a Metrohm 860 KF Thermoprep instrument. Here the term “moisture” relates to the Al(OH)3 content of the powder, as well as any adsorbed water. After completion of the L-PBF-processing, the moisture/hydroxide content of the Al–Cu powder was 598.5 ppm or 0.06%. This was five times higher than that of pristine Al-powder, of 118.9 ppm or 0.012%. The flow and moisture results of the powders are summarised in Table 2. 2.2. Laser-Powder Bed Fusion process The powders were processed using a Renishaw AM400 system equipped with a 400 W SPI fibre laser and a wavelength of 1.06 μm, which is used in continuous mode. Samples were fabricated in a reduced build volume platform of 78 × 78 × 55 mm, and the chamber was flooded with argon to ensure an oxygen level below 1000 ppm. The build platform was not preheated in this study. During the scanning process, an inert gas flow crosses the platform from right to left, removing the unwanted spattered particles from the powder bed. A meander scan strategy was selected to build the parts, which rotates the laser scanning vector 67° after each scan. The main process parameters involved in the laser L-PBF process are the laser power, hatching distance, point distance, exposure time, and layer thickness. Values for these were 200 W, 0.13 mm, 80 μm, and 140 μs, respectively. For this study, the layer thickness was kept constant at 30 μm. The process parameters giving the highest density for uncoated AlSi10Mg were determined first and were used as the starting point for evaluating the conditions for the coated powder. Samples (10 × 10 × 10 mm3 ) were built, mounted in epoxy resin, and polished down to 0.020 μm using colloidal silica suspension. The density was assessed using 20 optical images, which were tiled and image-processed. For the mechanical testing, builds were made (Fig. 1) containing hexagonal section bars for tensile specimens. Three cubes were printed on top of the bars for porosity level evaluation of the XY, XZ and YZ cross-sections. An additional cube (“4”) was used for XRD measurement and the pyramidal structure (“5”) was used for a hardness test. On selected samples, a Solution Heat Treatment (SHT) was applied, followed by artificial aging (AA) and heat treatment (T6). These were performed following the ASTM F3318-18 standard [21]. The solution heat treatment was carried out at 530 °C for 6 h followed by water quenching and subsequent artificial aging at 160 °C for 6 h. 2.3. Materials characterisation The microstructure of the samples was analysed using a JEOL JSM 6610 Scanning Electron Microscope (SEM) equipped with an Oxford
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using beam conditions of 20 kV, 5.5 nA and a 12 mm working distance. X-ray diffraction (XRD) measurements were performed using a Rigaku Miniflex with Cu K𝛼 radiation operated at 30 kV and 15 mA, with a step size of 0.01° (2𝜃) and a scanning speed of 0.03°/min. Consolidated material was analysed using scanning transmission electron microscopy (STEM) on an aberration-corrected JEOL 2100FCs S/TEM operating at 200 kV. The electron-transparent lamellae were prepared using the FEI Helios 600i dual beam focused ion beam (FIB) instrument, with a size thickness of ≈50–100 nm. Specimens were prepared by the lift out method and low energy (2 kV) polishing was applied as a final step to help removing potential Ga-ion induced surface amorphisation. On samples coated with copper, Electron Energy Loss Spectroscopy (EELS) was performed using the JEOL 2100FCs equipped with a GIF Quantum 693 electron energy filter. Images were processed using Gatan Digital Micrograph, applying periodic masks on FFT images. 2.4. Mechanical properties
Fig. 1. View of the build geometry used for the tensile, porosity, XRD and hardness specimens.
INCA X-act energy-dispersive X-ray spectroscopy (EDX). Fractographs of the tensile tested surfaces were obtained using the same microscope. Samples were etched using Keller’s reagent to reveal the microstructure. Texture measurements were performed on a FEI Helios Nanolab 600i Focused Ion Beam (FIB) instrument equipped with an EDAX-EBSD system. After mechanical polishing, samples were prepared using a Gatan Precision Ion Polishing System (PIPS). Samples were polished for 1 h, under a 5 kV Ar ion beam at a 5° incident angle. High resolution electron backscattered diffraction (EBSD) maps were collected using a step size of 80 nm in areas of fine-grained material (10k times magnification), and 250 nm steps were used to map large areas of coarse-grained material at lower (750x magnification). All EBSD maps were recorded
For mechanical properties evaluation, samples were manufactured with their longitudinal axes aligned parallel to the build plate, and three cylindrical “dog-bone” tension specimens were extracted from each build. Specimens built perpendicular to the build plate were not considered in this study. Tensile tests were conducted using 6 mm diameter and 24 mm gauge length specimens, machined in accordance with ASTM standard E8/E8M [22]. An Instron 5984 universal testing machine equipped with a 100 kN load cell was used, applying a displacement rate of 0.005 mm/mm/min up to the yield point and of 0.05 mm/mm/min until fracture. Hardness Vickers measurements were performed using a weight of 500 g in a Wilson VH3100 machine, in accordance with ASTM E92-17 standard [23]. A total of 12 indents were taken from the polished XY cross section. 3. Results 3.1. Al and Al–Cu powders Fig. 2 shows backscattered electron micrographs obtained from (a) the untreated and (b) Cu-treated aluminium powders. Fig. 2. Backscattered SEM images obtained at 20 kV of AlSi10Mg powder, (a) as received; and (b) after treatment with copper formate complex in methanol solution and vacuum-drying at 220 °C; (c) magnified view of copper-coated particle; (d) EDX map of Cu distribution.
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same area as (c), indicating that the bright particulates correlate with the presence of copper. The presence of copper was also demonstrated by XRD, where peaks corresponding to face-centred cubic (fcc) Cu were evident at 2𝜃 = 43.3° and 2𝜃 = 50.43°, as shown in Fig. 3. 3.2. Microstructure and texture
Fig. 3. XRD patterns of Al and Al–Cu powder with copper peaks highlighted.
Morphology of the Al powder was a mixture of spherical particles and elongated particles and satellites. The backscattered image is sensitive to the atomic number of elements in the samples and the bright particulate distribution shows the copper is discontinuous and of various sizes. The variability of the copper distribution is shown more clearly in Fig. 2(c) indicating that the dusting of copper is predominantly less than 2 μm in diameter and irregularly shaped. Fig. 2(d) shows an energy-dispersive X-ray (EDX) map of the Cu k𝛼1 radiation from the
SEM examination of etched XZ planes were performed to observe the cellular morphology evolution within the melt pool. Fig. 4(a) shows a repetitive fish-scale-like pattern, which has been previously reported by Zhou et al. [24]. The depth of the melt pools varies across the section, with values between one and three times the layer thickness. This is in accordance with the partial re-melting of previously solidified layers and following epitaxial growth of grains that is expected during the L-PBF process. Three different types of cells were identified across the melt pool. A fine cellular microstructure was developed close to the melt pool boundary triple points, MPB1, as shown in Fig. 4(b), while a coarser quasi-equiaxed cell type was formed on standard melt pool boundaries, MPB2, illustrated in Fig. 4(c). MPB1 is developed in the heat affected zone (HAZ) that is generated in the areas where the melt pools are overlapped. The bulk region (BR) represented in Fig. 4(d) was characterized by elongated cells with long axes parallel to the build direction. The measurements of the cells yielded a cell width of ≈ 300 nm in the elongated cells of the bulk region, while quasi-equiaxed cells at MPB1 and MPB2 showed diameters of ≈ 400 nm and ≈ 600 nm, respectively. It was observed that the cellular structure located at the edges follows a direction normal to the melt pool boundary. Closer observation of the cell structures using STEM, illustrated in Fig. 5, reveals the elemental composition of the bulk and intercellular regions of the L-PBF-processed Al–Cu. Supersaturated 𝛼-Al phase is Fig. 4. SEM image of etched Al–Cu XZ section at (a) high magnification with melt pool boundary highlighted in red and high magnification images of melt pool boundary 1 (MPB1), melt pool boundary 2 (MPB2) and bulk region (BR) in (b), (c) and (d), respectively.
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Fig. 5. TEM HAADF image of the Al–Cu L-PBF-processed sample and corresponding EDX mapping of the main constituent elements. Fig. 6. EBSD orientation maps of L-PBF-processed (a) Al–Cu and (b) Al samples, with their grain size showed in (c) and (d).
observed on the bulk of the cellular structure in the Al K map. The intercellular region is decorated with homogeneously distributed silicon and localized copper particles. A small proportion of oxygen and iron was also found in the cell boundary. The EBSD orientation maps of the L-PBF-processed Al and Al–Cu and their corresponding grain size distribution are shown in Fig. 6, in order to reveal the overall crystallographic texture of the material. Strong texture was developed in Al and Al–Cu material, with a bimodal grain structure composed of elongated and equiaxed grains respectively. Morphologically, the addition of copper did not result in an increase in small equiaxed grains. However, the Al–Cu material promoted an overall
grain size reduction of ≈ 15 μm. In both cases, the grain size increases towards the centre of the melt pool, with long elongated grains crossing multiple melt pools. These grains showed a ⟨001⟩ crystal growth direction along the building direction, i.e., parallel to the build direction. Some of these columnar grains are slightly tilted towards the ⟨101⟩ direction. In contrast, a proportion of small equiaxed grains showed no evidence of any texture. With respect to the grain sizes, columnar grains range between 40–150 μm in length and 5–15 μm in width, while the equiaxed grains share a diameter of approximately 6 μm. Fig. 7 shows the XRD of the L-PBF-processed parts, which do not exhibit the copper peaks from discrete particles seen in the Cu-treated
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Materialia 9 (2020) 100590
1000
Al-Cu Al
900
Intensity (a.u.)
800 700 600 500 400 300 200 100 0
30
40
50
60
70
80
2 (°) Fig. 7. XRD patterns of L-PBF-processed Al and Al–Cu samples. Fig. 9. STEM-BF image of L-PBF-processed Al–Cu sample.
powder (Fig. 3). However, energy dispersive X-ray analysis of the L-PBF-processed Al–Cu samples indicates the presence of copper, revealing that the copper coating has been incorporated within the alloy. This was also confirmed by the EELS spectrum in Fig. 8, showing an energy loss at 951 eV, corresponding to the L2 edge of copper. The image in Fig. 9 shows evidence of dislocation networks in some of the cells of the L-PBF-processed Al–Cu sample. Discrete copper particles were present on the grain boundaries, generating the pile-up of dislocations. 3.3. Mechanical properties An initial density evaluation on three XY plane cross-sections of Al–Cu cubes fabricated using the optimised parameters of Al material was done. Results showed an average density of 99.1 ± 0.2% and, therefore, the optimised parameters of Al were used in the subsequent Al–Cu builds. Density measurements of the three cross-sections were taken from cubes located in top of the tensile bars, for confirmation purposes, the results of which can be seen in Table 3. It can be observed that the density of the confirmation Al–Cu cubes, which are stacked on top of the bars, decreased compared to the initial value of 99.1 ± 0.2%. The different tensile stress-strain curves for the as-built and heattreated specimens are shown in Fig. 10, with their respective data indicated in Table 4. Generally, the samples behaved in a relatively ductile manner, as can be deduced from the elongation at break and high ultimate tensile strength values. The Al–Cu specimens showed very homogeneous UTS values, while the elongation at break values were
Table 3 Density values from the L-PBF-processed Al and Al–Cu build used for tensile test bars. Cross-section Material
XY
XZ
YZ
Al Al for T6 Al–Cu Al–Cu for T6
98.7 99.1 94.4 97.4
99.5 99.3 95.0 96.5
99.8 99.0 94.2 95.3
more scattered. The as-built L-PBF-processed Al specimens had an ultimate tensile strength and elongation at break values of 410 ± 10.1 MPa and 7 ± 0.5%, respectively. These values exceed that of the majority of cast and wrought alloys [1]. This is generally attributed to the very fine microstructure developed upon the rapid solidification that occurs during the L-PBF process. When the Al specimens were subjected to T6 heat treatment, the material ductility was slightly enhanced, from 7.0 ± 0.5% to 9.2 ± 0.5%. On the other hand, the UTS was reduced by 100 MPa, from 410 ± 10.1 MPa to 310 ± 4.4 MPa. The as-built Al–Cu material displayed a UTS of 461.0 ± 0.2 MPa, 51 MPa higher than that of the Al. Also, it exhibited a higher ductility, with an elongation at break 3% higher than that of the Al. In contrast to the Al specimens, the elongation of Al–Cu specimens decreased from 12 ± 0.5% to 9.7 ± 2.7%
Fig. 8. HAADF-STEM image of (a) Al–Cu sample with the incident beam location highlighted with the red marker and, (b) its corresponding EELS spectrum.
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Materialia 9 (2020) 100590
Table 4 Mechanical properties of L-PBF-processed Al and Al–Cu.
AlSi10MgAs built AlSi10MgT6 Al–CuAs built Al–CuT6
UTS (MPa)
Yield Strength (MPa)
Young Modulus (GPa)
Elongation at break (%)
410 310 461 346
267 244 264 254
70.8 (±2) 72.4 (±2.4) 73 (±1.8) 76.3 (±7)
7 (±0.5) 9.2 (±0.5) 12 (±0.5) 9.7 (±2.7)
(±10.1) (±4.4) (±0.2) (±1)
(±3.7) (±3.1) (±1.8) (±13.74)
for the as-built and heat-treated Al–Cu was significantly higher than that of Al. Indent 9 of the Al–Cu T6 sample was excluded from the data as it was compromised by local porosity, rendering it unrepresentative of the test material. In general, the reduction of hardness after T6 was smaller on the coated powder material. Measurements on as-built consolidated material yielded values of 117.3 ± 3 HV0.5 and 110.3 ± 5.8 HV0.5 for Al and Al–Cu samples, respectively. However, when samples were subjected to T6 heat treatment, the Al–Cu sample showed an average hardness of 101.6 ± 14.2 HV0.5 , while the Al sample exhibited a value of 90.6 ± 2.3 HV0.5 . The hardness results are concordant with the UTS results obtained after tensile testing, excluding the as-built Al–Cu values, which would be expected to be greater than those observed for Al. 4. Discussion 4.1. Microstructure and texture
Fig. 10. Engineering tensile stress-strain curves of (a) as build and (b) heat treated Al and Al–Cu L-PBF-processed samples.
when subjected to heat treatment. The UTS reduction for heat treated Al–Cu was approximately 25%, the same in the case of Al. Tensile tested specimen fractographs, shown in Fig. 11, were used to reveal the brittle or ductile nature of the samples, which is associated with the sample elongation capacity. Macroscopically, the fracture surfaces showed little plastic deformation, however the microstructural appearance was fibrous, which is associated with ductile fracture. No necking was seen macroscopically, which is consistent with the fracture stress being equal to the ultimate tensile strength in Fig. 10. Failure was initiated at a surface defect, such as pore or a crack, which propagated across the fracture plane, perpendicular to the load direction. Some voids with un-melted powder were observed in the fractographs, as well as shallow dimples with a diameter range of ≈ 13–66 μm. Hardness test results showed different micro hardness values (HV0.5 ) for both materials, as shown in Fig. 12. The scatter of hardness values
The processability of metallic powders by Laser-Powder Bed Fusion is influenced by factors such as powder morphology, particle size distribution, flowability, chemical composition, moisture content, and laser absorptivity. The attainment of a homogeneous powder layer for the L-PBF process is determined by the particle size distribution, powder flowability, and packing density. The powder coating method employed here did not significantly modify the flow behaviour of the powders, as it does not alter the powder morphology or particle size distribution. During processing, a layer of powder is deposited over previously solidified material, which is partially re-melted by the laser. The macrostructure of the consolidated material shows fish-scale-like features, remnant of melt pools with heights up to three times the layer thickness. The EBSD results demonstrated the epitaxial growth of fine grains in the direction parallel to the build direction. This resulted in long elongated grains that cross several cells and melt pools. As a consequence, the epitaxially grown grains share their crystal orientation. This general texture of the materials contributed to the undesired anisotropy of the material. Some quasi-equiaxed grains were observed in both materials, while the Al–Cu samples showed an overall grain size reduction. The precipitation sequence of Al–Si–Mg alloys usually starts with the supersaturated solid solution trapping solute atoms in solution within the solid. These solute atoms adopt the form of metastable Guinier–Preston (GP) zones. These are decomposed into needle-like 𝛽’’ phases, which then transforms into the semi-coherent rod-shaped 𝛽’Mg2 Si phase. Finally, the latter is transformed into a stable, incoherent, fcc 𝛽-Mg2 Si phase (in the form of plates) [25]. As previously reported, aluminium cells are the first to solidify and subsequently, the liquid of approximately eutectic composition is contained in the intercellular region [26]. The uncoated samples present an aluminium matrix with fine eutectic silicon populating the intercellular regions. Upon material solidification, copper is pushed towards the intercellular region, were the eutectic Al–Si is formed. This explains why most of the copper is concentrated in the cell boundary, although some copper-enriched precipitates were visible by TEM on the 𝛼-Al matrix. The precipitation sequence in Al–Cu alloys is quite complex, and it occurs in a series of reactions leading to metastable precipitates before the formation of the equilibrium phase 𝜃-Al2 Cu is achieved. Some needle-shaped iron-rich precipitates were also found in the matrix of coated and uncoated material cells, particularly after T6 heat treatment. These precipitates are undesirable, forming elongated brittle 𝛽-Al5 FeSi phases. The heat
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Materialia 9 (2020) 100590
Fig. 11. Fracture surfaces of L-PBFprocessed (a) Al, (b) T6 Al, (c) Al–Cu, (d) T6 Al–Cu.
4.2. Mechanical properties
Fig. 12. Hardness Vickers (HV0.5 ) values for as-built and heat-treated Al and Al–Cu samples.
treatment promotes the formation of silicon platelets, which have a brittle diamond crystal structure. Consequently, these precipitates are detrimental for mechanical properties since they act as sites for crack nucleation under cyclic loading [27].
According to previous studies, there are four main strengthening mechanisms in metal matrix composites: dislocation strengthening, Orowan strengthening, grain boundary reduction strengthening and load-bearing strengthening [15]. The aluminium alloys processed in this study showed UTS values higher than conventional cast A360. This is attributed to the fine grain and eutectic cellular structure developed as a result of the rapid solidification during the L-PBF process. In addition, the dislocation entanglement shown by the TEM contributes to the material strengthening by dislocation motion hindering. The higher porosity level of the L-PBF-processed Al–Cu specimens, may be attributable to unoptimized laser parameters or the evolution of volatile gases resulting from the coating process. Nevertheless, the additional porosity did not appear to detrimentally influence the mechanical properties of the as-built or heat treated material. In fact, both the as-built and T6-treated Al–Cu material exhibited higher UTS and elongation values than the AlSi10Mg counterparts. This increase in the ultimate tensile strength of the Al–Cu specimens do not show the expected correlation with the HV0.5 results. This is attributed to the higher porosity shown by those specimens, which leads to a higher scattering in the hardness results. The probability of an indent reaching a zone with a pore and, therefore, yielding a low hardness value, are higher in the Al–Cu specimens. We attribute, the origin of the UTS enhancement in the Al–Cu case to a combination of Hall–Petch and Orowan strengthening mechanisms. One factor is the grain refinement shown in the EBSD analysis and promoted by the copper addition, which contributes to the Hall–Petch strengthening. A second factor is the copper precipitates
X. Garmendia, S. Chalker and M. Bilton et al.
located both on the intercellular and matrix sites which contribute to the dislocation motion hindering, i.e., Orowan strengthening. The overall reduction in hardness and tensile strength when the specimens were subjected to aging is promoted by the nucleation and growth of the silicon precipitates of brittle nature, in the so-called Ostwald ripening process. The size and volume fraction of those precipitates influence the strength of the material, following the Orowan equation. The bigger the precipitate size is, the lower the strength of the material will be. After the same heat treatment conditions were applied, the Al–Cu specimens yielded a higher ultimate tensile strength. The size distribution of the silicon precipitates was smaller in the heat-treated Al–Cu material, which leads the higher UTS values. 5. Conclusions In this study, AlSi10Mg powder was coated with 1 wt% copper and samples were fabricated using Laser-Powder Bed Fusion. The microstructure and mechanical properties of the uncoated and coated materials were evaluated. The processability of copper coated AlSi10Mg powder has been demonstrated. The copper coated specimens showed an increase of 51 MPa in the ultimate tensile strength when compared to AlSi10Mg. This is attributed to the grain size reduction and consequent Hall–Petch strengthening, together with Oworan strengthening promoted by the copper-enriched precipitates. Microstructurally, the material consists of an 𝛼-Al cell matrix with eutectic Si decorating the cell boundaries. Furthermore, copper is segregated in the cell boundary. Both materials exhibit silicon platelets and needle-like iron-rich precipitates when subjected to peak hardening. However, the copper addition causes a reduction in the size of the silicon precipitates after heat treatment. The study establishes the feasibility of modifying the composition of L-PBF-processed materials via the coating of powder feedstocks and that the coating additions can be exploited to enhance mechanical properties of parts after heat treatment. Declaration of Competing Interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgments The authors gratefully acknowledges financial support from Renishaw plc and the Engineering and Physical Sciences Research Council (EPSRC) Centre for Doctoral Training in Additive Manufacturing and 3D Printing. We also acknowledge funding support from the Innovate UK project FLAC (102665) and EPSRC project EP/N017773/1. The authors would also like to thank Dr Karl Dawson for helpful discussions regarding the electron microscopy and texture analysis. References [1] K. Kempen, L. Thijs, J. Van Humbeeck, J.-.P. Kruth, Mechanical properties of AlSi10Mg produced by selective laser melting, Phys. Proc. 39 (2012) 439–446, doi:10.1016/J.PHPRO.2012.10.059. [2] B.A. Fulcher, D.K. Leigh, T.J. Watt, Comparison of AlSi10Mg and Al 6061 processed through DMLS, in: B.A. Fulcher, D.K. Leigh, T.J. Watt (Eds.), Harvest Technologies, Department of Mechanical Engineering, The University of Texas at Austin, Belton, TX 76513, Austin, TX 78712, 2014, pp. 404–419. [3] E. Louvis, P. Fox, C.J. Sutcliffe, Selective laser melting of aluminium components, J. Mater. Process. Technol. 211 (2011) 275–284, doi:10.1016/j.jmatprotec.2010.09.019.
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