properties in structures with a low number of cells, where this buckling tends to play a relevant role in the general behavior of the structure.

* The Gibson-Ashby model predicts both the elastic modulus and the yield stress with good precision and low size effect. To predict the plastic zone properties, such as the tangent modulus, size effect must be taken into account, since considerable change in the exponent of the Gibson-Ashby formula is observed for the different boundary condition models.

Ongoing research will focus on correlating these numerical results with experiments to validate the adopted methodology. In addition, for impact and energy absorption applications, the evaluation of the behavior of structures with different numbers of cells and at stages of larger deformation or densification should be considered.

## Acknowledgements

This work is funded by WBI/AGCID RI02 (DIE23-0001) bilateral project between the University of Liège and the Universidad de La Frontera. As research director of F.R.S.-FNRS, A.M. Habraken thanks the Fund for Scientific Research for financial support.

## References

[1] Valle R, Pincheira G, Tuninetti V, et al (2022) Evaluation of the Orthotropic Behavior in an Auxetic Structure Based on a Novel Design Parameter of a Square Cell with Re-Entrant Struts. Polymers (Basel) 14:4325. https://doi.org/10.3390/polym14204325

[2] Rilling S, Ríos I, Gómez Á, et al (2023) Optimized infill density through topological optimization increases strength of additively manufactured porous polylactic acid. Int J Adv Manuf Technol. https://doi.org/10.1007/s00170-023-12554-z

[3] Almesmari A, Alagha AN, Naji MM, et al (2023) Recent Advancements in Design Optimization of Lattice-Structured Materials. Adv Eng Mater 25:. https://doi.org/10.1002/adem.202201780

[4] Valle R, Pincheira G, Tuninetti V, et al (2022) Design and Characterization of Asymmetric Cell Structure of Auxetic Material for Predictable Directional Mechanical Response. Materials (Basel) 15:. https://doi.org/10.3390/ma15051841

[5] Maconachie T, Leary M, Lozanovski B, et al (2019) SLM lattice structures: Properties, performance, applications and challenges. Mater Des 183:108137. https://doi.org/10.1016/j.matdes.2019.108137

[6] Favre J, Lohmuller P, Piotrowski B, et al (2018) A continuous crystallographic approach to generate cubic lattices and its effect on relative stiffness of architectured materials. Addit Manuf 21:359-368. https://doi.org/10.1016/j.addma.2018.02.020

[7] Wang C, Zhu J, Wu M, et al (2021) Multi-scale design and optimization for solid-lattice hybrid structures and their application to aerospace vehicle components. Chinese J Aeronaut 34:386-398. https://doi.org/10.1016/j.cja.2020.08.015

[8] Syrlybayev D, Perveen A, Talamona D (2023) Experimental investigation of mechanical properties and energy absorption capabilities of hybrid lattice structures manufactured using fused filament fabrication. Int J Adv Manuf Technol 125:2833-2850. https://doi.org/10.1007/s00170-023-10922-3

[9] Zhang K, Shi D, Wang W, Wang Q (2017) Mechanical characterization of hybrid lattice-to-steel joint with pyramidal CFRP truss for marine application. Compos Struct 160:1198-1204. https://doi.org/10.1016/j.compstruct.2016.11.005
[10] Pasini C, Pandini S, Ramorino G, Sartore L (2024) Journal of the Mechanical Behavior of Biomedical Materials Tailoring the properties of composite scaffolds with a 3D-Printed lattice core and a bioactive hydrogel shell for tissue engineering. J Mech Behav Biomed Mater 150:106305. https://doi.org/10.1016/j.jmbbm.2023.106305

[11] Xu F, Zhang X, Zhang H (2018) A review on functionally graded structures and materials for energy absorption. Eng Struct 171:309-325. https://doi.org/10.1016/j.engstruct.2018.05.094

[12] Pirinu A, Primo T, Del Prete A, et al (2023) Mechanical behaviour of AlSi10Mg lattice structures manufactured by the Selective Laser Melting (SLM). Int J Adv Manuf Technol 124:1651-1680. https://doi.org/10.1007/s00170-022-10390-1

[13] Wang Z, Zhou Y, Wang X, Wei K (2021) International Journal of Mechanical Sciences Compression behavior of strut-reinforced hierarchical lattice — Experiment and simulation. Int J Mech Sci 210:106749

[14] Ashby MF (2006) The properties of foams and lattices. Philos Trans R Soc A Math Phys Eng Sci 364:15-30. https://doi.org/10.1098/rsta.2005.1678

[15] Rodrigo C, Xu S, Durandet Y, et al (2023) Quasi-static and dynamic compression of additively manufactured functionally graded lattices: Experiments and simulations. Eng Struct 284:115909. https://doi.org/10.1016/j.engstruct.2023.115909

[16] Gibson LJ AM (1997) Cellular Solids: Structure and Properties, 2nd ed

[17] Zadpoor AA (2019) Mechanical performance of additively manufactured meta-biomaterials. Acta Biomater 85:41-59. https://doi.org/10.1016/j.actbio.2018.12.038

[18] Bruson D, Galati M, Calignano F, Iuliano L (2023) Mechanical Characterisation and Simulation of the Tensile Behaviour of Polymeric Additively Manufactured Lattice Structures. Exp Mech. https://doi.org/10.1007/s11340-023-00976-5

[19] Köhnen P, Haase C, Bültmann J, et al (2018) Mechanical properties and deformation behavior of additively manufactured lattice structures of stainless steel. Mater Des 145:205-217. https://doi.org/10.1016/j.matdes.2018.02.062

[20] Maskery I, Aboulkhair NT, Aremu AO, et al (2017) Compressive failure modes and energy absorption in additively manufactured double gyroid lattices. Addit Manuf 16:24-29. https://doi.org/10.1016/j.addma.2017.04.003
[10] Pasini C, Pandini S, Ramorino G, Sartore L (2024) Journal of the Mechanical Behavior of Biomedical Materials Tailoring the properties of composite scaffolds with a 3D-Printed lattice core and a bioactive hydrogel shell for tissue engineering. J Mech Behav Biomed Mater 150:106305. https://doi.org/10.1016/j.jmbbm.2023.106305

[11] Xu F, Zhang X, Zhang H (2018) A review on functionally graded structures and materials for energy absorption. Eng Struct 171:309–325. https://doi.org/10.1016/j.engstruct.2018.05.094

[12] Pirinu A, Primo T, Del Prete A, et al (2023) Mechanical behaviour of AlSi10Mg lattice structures manufactured by the Selective Laser Melting (SLM). Int J Adv Manuf Technol 124:1651–1680. https://doi.org/10.1007/s00170-022-10390-1

[13] Wang Z, Zhou Y, Wang X, Wei K (2021) International Journal of Mechanical Sciences Compression behavior of strut-reinforced hierarchical lattice — Experiment and simulation. Int J Mech Sci 210:106749

[14] Ashby MF (2006) The properties of foams and lattices. Philos Trans R Soc A Math Phys Eng Sci 364:15–30. https://doi.org/10.1098/rsta.2005.1678

[15] Rodrigo C, Xu S, Durandet Y, et al (2023) Quasi-static and dynamic compression of additively manufactured functionally graded lattices: Experiments and simulations. Eng Struct 284:115909. https://doi.org/10.1016/j.engstruct.2023.115909

[16] Gibson LJ AM (1997) Cellular Solids: Structure and Properties, 2nd ed

[17] Zadpoor AA (2019) Mechanical performance of additively manufactured meta-biomaterials. Acta Biomater 85:41–59. https://doi.org/10.1016/j.actbio.2018.12.038

[18] Bruson D, Galati M, Calignano F, Iuliano L (2023) Mechanical Characterisation and Simulation of the Tensile Behaviour of Polymeric Additively Manufactured Lattice Structures. Exp Mech. https://doi.org/10.1007/s11340-023-00976-5

[19] Köhnen P, Haase C, Bültmann J, et al (2018) Mechanical properties and deformation behavior of additively manufactured lattice structures of stainless steel. Mater Des 145:205–217. https://doi.org/10.1016/j.matdes.2018.02.062

[20] Maskery I, Aboulkhair NT, Aremu AO, et al (2017) Compressive failure modes and energy absorption in additively manufactured double gyroid lattices. Addit Manuf 16:24–29. https://doi.org/10.1016/j.addma.2017.04.003
# On the experimentation of the laser-powder bed fusion process for high-strength aluminum alloys

EL HASSANIN Andrea<fnref n="1" />,ª*, PERNA Alessia Serena<fnref n="1" />,ᵇ, SICIGNANO Nicola<fnref n="1" />,<fnref n="2" />,ᶜ, BORRELLI Domenico Borrelli<fnref n="1" />,<fnref n="2" />,ᵈ, CARAVIELLO Antonio<fnref n="1" />,<fnref n="2" />,ᵉ and ASTARITA Antonello<fnref n="1" />,*f

**Keywords:** Laser-Powder Bed Fusion, Aluminum Alloys, Material Characterization

**Abstract.** Additive Manufacturing (AM), a cutting-edge technique developed for the manufacture of prototypes that aims to produce large-scale products tailored for numerous industrial applications, has found increasing space in recent years. In this context, Laser-Powder Bed Fusion (L-PBF) technology represents one of the most intriguing ones considering the large number of processable and under-development alloys. By combining the inherent benefits of this technology, it is then possible to produce components with improved mechanical characteristics and low weight. Nevertheless, improved properties such as thermal and electrical conductivity, strength, and corrosion resistance can be achieved. In this scenario, high-performance aluminum-based alloys, tailored for the metal AM have raised increasing interest in the industrial framework. In this work, the LPBF processing of an Al-Mg-Sc-Zr was investigated. The tensile characteristics, the microhardness, the density and the microstructure of L-PBF specimens were analyzed.

**Introduction**

The use of powder-based Additive Manufacturing (AM) technologies is transforming the process of creating tailored parts. In this field, Laser-Powder Bed Fusion (L-PBF) is particularly notable as the most cutting-edge and dependable technique for fabricating metal objects [1,2]. In the past few years, there was a significant amount of investigation into the application of Laser-Powder Bed Fusion (LPBF) technology for the AM of metal parts [3,4]. This research explored the use of various pure metals and alloys, the latter including steel, aluminum, titanium, and nickel alloys [5-7]. The primary focus of this research has been two-fold: the creation of complex-shaped parts and the development of materials that possess enhanced mechanical properties and customized physical characteristics to meet the specific requirements of high-performance applications. Among the different types of materials employed in traditional manufacturing techniques, only a limited few have been tailored for additive manufacturing to enhance the manufacturing process. In the most competitive industrial sectors, like the aviation sector, it is imperative to employ a metal alloy with an optimal blend of mechanical properties, including high strength and stiffness but also good formability. Hence, accurately evaluating the underlying metallurgical mechanisms is critical in attaining these desired qualities. In this scenario, the aerospace industry makes extensive use of high-strength Al alloys to create structural components [8]. These alloys offer exceptional strength-to-weight and stiffness-to-weight ratios, which, when coupled with the advantages of AM techniques, result in minimized material wastage and energy usage throughout the fabrication process [9].

<figure><img src="image_226.png" /></figure>

<footnote n="1">¹Department of Chemical, Materials and Production Engineering, University of Naples Federico II, Piazzale V. Tecchio 80, 80125 Naples, Italy</footnote>

<footnote n="2">²Sòphia High Tech S.r.l, Via Malatesta 30 A, 80049, Naples, Italy</footnote>

<footnote>ªandrea.elhassanin@unina.it, ᵇalessiaserena.perna@unina.it,</footnote>

<footnote>ᶜnicola.sicignano@sophiahightech.com, ᵈdomenico.borrelli@sophiahightech.com,</footnote>

<footnote>ᵉantonio.caraviello@sophiahightech.com, ᶜantonello.astarita@unina.it</footnote>
# On the experimentation of the laser-powder bed fusion process for high-strength aluminum alloys

EL HASSANIN Andrea¹,ª*, PERNA Alessia Serena¹,ᵇ, SICIGNANO Nicola¹,²,ᶜ, BORRELLI Domenico Borrelli¹,²,ᵈ, CARAVIELLO Antonio¹,²,ᵉ and ASTARITA Antonello¹,*f

¹Department of Chemical, Materials and Production Engineering, University of Naples Federico II, Piazzale V. Tecchio 80, 80125 Naples, Italy

²Sòphia High Tech S.r.l, Via Malatesta 30 A, 80049, Naples, Italy

ªandrea.elhassanin@unina.it, ᵇalessiaserena.perna@unina.it,

ᶜnicola.sicignano@sophiahightech.com, ᵈdomenico.borrelli@sophiahightech.com,

ᵉantonio.caraviello@sophiahightech.com, ᶜantonello.astarita@unina.it

**Keywords:** Laser-Powder Bed Fusion, Aluminum Alloys, Material Characterization

**Abstract.** Additive Manufacturing (AM), a cutting-edge technique developed for the manufacture of prototypes that aims to produce large-scale products tailored for numerous industrial applications, has found increasing space in recent years. In this context, Laser-Powder Bed Fusion (L-PBF) technology represents one of the most intriguing ones considering the large number of processable and under-development alloys. By combining the inherent benefits of this technology, it is then possible to produce components with improved mechanical characteristics and low weight. Nevertheless, improved properties such as thermal and electrical conductivity, strength, and corrosion resistance can be achieved. In this scenario, high-performance aluminum-based alloys, tailored for the metal AM have raised increasing interest in the industrial framework. In this work, the LPBF processing of an Al-Mg-Sc-Zr was investigated. The tensile characteristics, the microhardness, the density and the microstructure of L-PBF specimens were analyzed.

**Introduction**

The use of powder-based Additive Manufacturing (AM) technologies is transforming the process of creating tailored parts. In this field, Laser-Powder Bed Fusion (L-PBF) is particularly notable as the most cutting-edge and dependable technique for fabricating metal objects [1,2]. In the past few years, there was a significant amount of investigation into the application of Laser-Powder Bed Fusion (LPBF) technology for the AM of metal parts [3,4]. This research explored the use of various pure metals and alloys, the latter including steel, aluminum, titanium, and nickel alloys [5-7]. The primary focus of this research has been two-fold: the creation of complex-shaped parts and the development of materials that possess enhanced mechanical properties and customized physical characteristics to meet the specific requirements of high-performance applications. Among the different types of materials employed in traditional manufacturing techniques, only a limited few have been tailored for additive manufacturing to enhance the manufacturing process. In the most competitive industrial sectors, like the aviation sector, it is imperative to employ a metal alloy with an optimal blend of mechanical properties, including high strength and stiffness but also good formability. Hence, accurately evaluating the underlying metallurgical mechanisms is critical in attaining these desired qualities. In this scenario, the aerospace industry makes extensive use of high-strength Al alloys to create structural components [8]. These alloys offer exceptional strength-to-weight and stiffness-to-weight ratios, which, when coupled with the advantages of AM techniques, result in minimized material wastage and energy usage throughout the fabrication process [9].

<figure><img src="image_230.png" /></figure>
As a result, the final aircraft exhibits enhanced fuel efficiency, thus reducing its environmental impact during operation. To address these challenges, researchers are conducting more studies to expand the range of usable alloys and developing methods to detect and rectify the different types of defects that may arise during L-PBF [10]. The focus of this study was to investigate the processability window through L-PBF of an Al-Mg-Sc-Zr alloy. This work studied the effect of the main L-PBF process parameters, namely laser power and scanning speed, on the final processability through L-PBF of the Al-Mg-Sc-Zr alloy, whose mechanical properties are enhanced both through solid dispersion and precipitation hardening [11]. To evaluate the mechanical properties, tensile tests were performed on the printed specimens. The tensile tests provided insights into the strength, ductility, and deformation behavior of the printed alloy. Vickers microhardness measurements were carried out to assess the hardness distribution and variations within the printed specimens. Density measurements were conducted to determine the porosity level in the printed specimens, as porosity can significantly influence the mechanical properties of the alloy. Microstructural analysis was performed using various techniques, such as optical microscopy and SEM-EDS. This analysis aimed to understand the sensitivity of the chosen alloy against variations of L-PBF processing conditions. Overall, this study provided a comprehensive understanding of the potential of L-PBF for printing Al-Mg-Sc-Zr alloys in the as-built condition: despite it is a age-hardenable alloy, knowing the starting point before heat treatment is necessary also to appreciate the mechanical properties improvement due to the treatment itself. The findings could provide insights for the process development of the selected alloy, given its high feasibility for the production of complex shaped parts with end use in fields where low weight and high performance are required.

## Materials and Methods

L-PBF specimens. For the experimental campaign, powders of Al-Mg-Sc-Zr aluminum alloy were used, supplied as custom powders by Metals 4 Printing [12]. To evaluate the actual composition of the powders, SEM-EDS analysis was carried out using the Oxford Instrument Swift ED 3000 Probe microscope coupled with an Hitachi TM3000 SEM, the latter also used to analyze the powders morphology illustrated in Fig. 1. Moreover, according to the ASTM B822 standard, the following particle size distribution was obtained: D₁₀ = 10.7 µm; D₅₀ = 30.2 µm; D₉₀ = 77.9 µm. Concerning the parts to be printed, it was decided to print 18 cubic specimens (2 specimens per set of process parameters) and 18 dog bone specimens with the same parameters, whose placement inside the process chamber followed the scheme illustrated in Fig. 2. The process parameters values were decided in order to investigate three levels of Volumetric Energy Density (VED, see Table 1), obtained through different combinations of laser power and scanning speed: in this way, the experiments enabled to analyze the feasibility of the process as a function of the most used parameters group, i.e. VED, as well as to check how the different “paths” leading to the same VED values affect the processed material [13], according to the data reported in Table 2. The specimens were printed with a Concept Laser Cusing M2 L-PBF machine, equipped with an yttrium laser source with a maximum power of 400 W and a focus diameter of 70 µm. The specimens were all built with an horizontal positioning on the build platform (see Fig. 1) in order to reduce the feedstock amount required. The material was processed through an island scan strategy, as similarly to other works described elsewhere [14].
As a result, the final aircraft exhibits enhanced fuel efficiency, thus reducing its environmental impact during operation. To address these challenges, researchers are conducting more studies to expand the range of usable alloys and developing methods to detect and rectify the different types of defects that may arise during L-PBF [10]. The focus of this study was to investigate the processability window through L-PBF of an Al-Mg-Sc-Zr alloy. This work studied the effect of the main L-PBF process parameters, namely laser power and scanning speed, on the final processability through L-PBF of the Al-Mg-Sc-Zr alloy, whose mechanical properties are enhanced both through solid dispersion and precipitation hardening [11]. To evaluate the mechanical properties, tensile tests were performed on the printed specimens. The tensile tests provided insights into the strength, ductility, and deformation behavior of the printed alloy. Vickers microhardness measurements were carried out to assess the hardness distribution and variations within the printed specimens. Density measurements were conducted to determine the porosity level in the printed specimens, as porosity can significantly influence the mechanical properties of the alloy. Microstructural analysis was performed using various techniques, such as optical microscopy and SEM-EDS. This analysis aimed to understand the sensitivity of the chosen alloy against variations of L-PBF processing conditions. Overall, this study provided a comprehensive understanding of the potential of L-PBF for printing Al-Mg-Sc-Zr alloys in the as-built condition: despite it is a age-hardenable alloy, knowing the starting point before heat treatment is necessary also to appreciate the mechanical properties improvement due to the treatment itself. The findings could provide insights for the process development of the selected alloy, given its high feasibility for the production of complex shaped parts with end use in fields where low weight and high performance are required.

## Materials and Methods

L-PBF specimens. For the experimental campaign, powders of Al-Mg-Sc-Zr aluminum alloy were used, supplied as custom powders by Metals 4 Printing [12]. To evaluate the actual composition of the powders, SEM-EDS analysis was carried out using the Oxford Instrument Swift ED 3000 Probe microscope coupled with an Hitachi TM3000 SEM, the latter also used to analyze the powders morphology illustrated in Fig. 1. Moreover, according to the ASTM B822 standard, the following particle size distribution was obtained: D₁₀ = 10.7 µm; D₅₀ = 30.2 µm; D₉₀ = 77.9 µm. Concerning the parts to be printed, it was decided to print 18 cubic specimens (2 specimens per set of process parameters) and 18 dog bone specimens with the same parameters, whose placement inside the process chamber followed the scheme illustrated in Fig. 2. The process parameters values were decided in order to investigate three levels of Volumetric Energy Density (VED, see Table 1), obtained through different combinations of laser power and scanning speed: in this way, the experiments enabled to analyze the feasibility of the process as a function of the most used parameters group, i.e. VED, as well as to check how the different “paths” leading to the same VED values affect the processed material [13], according to the data reported in Table 2. The specimens were printed with a Concept Laser Cusing M2 L-PBF machine, equipped with an yttrium laser source with a maximum power of 400 W and a focus diameter of 70 µm. The specimens were all built with an horizontal positioning on the build platform (see Fig. 1) in order to reduce the feedstock amount required. The material was processed through an island scan strategy, as similarly to other works described elsewhere [14].
<figure><img src="image_227.png" /><figcaption>Fig. 1 – Al-Mg-Sc-Zr powders used in this work.</figcaption></figure>

<table><caption>Table 1 – Chemical composition of the Al-Mg-Sc-Zr powder feedstock.</caption><thead><tr><th>Element</th><th>Al</th><th>Mg</th><th>Sc</th><th>Zr</th></tr></thead><tbody><tr><td>(wt%)</td><td>Bal.</td><td>5.1±0.1</td><td>0.4 ± 0.1</td><td>0.7 ± 0.2</td></tr></tbody></table>

<table><caption>Table 2 - Process parameters for printing the test specimens of the experimental campaign.</caption><thead><tr><th>ID</th><th>Layer Thickness (mm)</th><th>Laser Power (W)</th><th>Scan Speed (mm/s)</th><th>Hatch Distance (mm)</th><th>VED (J/mm³)</th></tr></thead><tbody><tr><td>1</td><td>0.03</td><td>200</td><td>847</td><td>0.105</td><td>75</td></tr><tr><td>2</td><td>0.03</td><td>200</td><td>635</td><td>0.105</td><td>100</td></tr><tr><td>3</td><td>0.03</td><td>200</td><td>508</td><td>0.105</td><td>125</td></tr><tr><td>4</td><td>0.03</td><td>240</td><td>1016</td><td>0.105</td><td>75</td></tr><tr><td>5</td><td>0.03</td><td>240</td><td>762</td><td>0.105</td><td>100</td></tr><tr><td>6</td><td>0.03</td><td>240</td><td>610</td><td>0.105</td><td>125</td></tr><tr><td>7</td><td>0.03</td><td>260</td><td>1101</td><td>0.105</td><td>75</td></tr><tr><td>8</td><td>0.03</td><td>260</td><td>825</td><td>0.105</td><td>100</td></tr><tr><td>9</td><td>0.03</td><td>260</td><td>660</td><td>0.105</td><td>125</td></tr></tbody></table>
<figure><img src="image_231.png" /><figcaption>Fig. 1 – Al-Mg-Sc-Zr powders used in this work.</figcaption></figure>

<table><caption>Table 1 – Chemical composition of the Al-Mg-Sc-Zr powder feedstock.</caption><thead><tr><th>Element</th><th>Al</th><th>Mg</th><th>Sc</th><th>Zr</th></tr></thead><tbody><tr><td>(wt%)</td><td>Bal.</td><td>5.1±0.1</td><td>0.4 ± 0.1</td><td>0.7 ± 0.2</td></tr></tbody></table>

<table><caption>Table 2 - Process parameters for printing the test specimens of the experimental campaign.</caption><thead><tr><th>ID</th><th>Layer Thickness (mm)</th><th>Laser Power (W)</th><th>Scan Speed (mm/s)</th><th>Hatch Distance (mm)</th><th>VED (J/mm³)</th></tr></thead><tbody><tr><td>1</td><td>0.03</td><td>200</td><td>847</td><td>0.105</td><td>75</td></tr><tr><td>2</td><td>0.03</td><td>200</td><td>635</td><td>0.105</td><td>100</td></tr><tr><td>3</td><td>0.03</td><td>200</td><td>508</td><td>0.105</td><td>125</td></tr><tr><td>4</td><td>0.03</td><td>240</td><td>1016</td><td>0.105</td><td>75</td></tr><tr><td>5</td><td>0.03</td><td>240</td><td>762</td><td>0.105</td><td>100</td></tr><tr><td>6</td><td>0.03</td><td>240</td><td>610</td><td>0.105</td><td>125</td></tr><tr><td>7</td><td>0.03</td><td>260</td><td>1101</td><td>0.105</td><td>75</td></tr><tr><td>8</td><td>0.03</td><td>260</td><td>825</td><td>0.105</td><td>100</td></tr><tr><td>9</td><td>0.03</td><td>260</td><td>660</td><td>0.105</td><td>125</td></tr></tbody></table>
<figure><img src="image_228.png" /><figcaption>Fig. 2 – Layout of the L-PBF process for the selected specimens and conditions.</figcaption></figure>

Characterization procedure. The produced specimens were characterized in terms of density, mechanical properties and microstructure. To this aim, cubic specimens were used for density, Vickers microhardness and microstructure analyses, while dog bone specimens were used for tensile tests. Density measurements were carried out in accordance with ASTM B962 standard, using a Gibertini Eternity 2000 hydrostatic weighing scale after removing the surface roughness of the specimens through grinding. Prior to the Vickers microhardness tests and microstructure analysis, the samples were prepared according to the standard metallographic procedure and polished until the surface quality was 1 µm [15], considering a plane parallel to the build direction of the specimens. The Vickers microhardness was analyzed, in accordance with the ASTM E92-17, through a square-based pyramid indenter with an angle of 148°, performing at least three indentations on each specimen, with a load of 200g and a dwell time of 15s. Tensile tests were performed according to the ASTM E8M standard and considering the reduced specimen dimensions. The tests were carried out using a Galdabini Quasar 50 universal testing machine, equipped with a 50 kN load cell. Microstructure analysis was carried out after etching the polished specimens with the Keller's reagent by swabbing for at least 30s. Optical microscopy was carried out to take the macrographs of the specimens processed under different conditions at 100x magnification.

## Results and discussion

Density and mechanical properties. Fig. 3 illustrates the L-PBF specimens produced according to the experimental plan described in the previous section (only 1 replicate is shown). The visual inspection indicated that the specimens were sound and without significant defects. However, tensile specimens exhibited a significant distortion caused by the printing process, considering the remarkable thermal conductivity and thermal expansion coefficient of the selected alloy in combination with the horizontal positioning and the absence of stress relief heat treatment post-build.
<figure><img src="image_232.png" /><figcaption>Fig. 2 – Layout of the L-PBF process for the selected specimens and conditions.</figcaption></figure>

Characterization procedure. The produced specimens were characterized in terms of density, mechanical properties and microstructure. To this aim, cubic specimens were used for density, Vickers microhardness and microstructure analyses, while dog bone specimens were used for tensile tests. Density measurements were carried out in accordance with ASTM B962 standard, using a Gibertini Eternity 2000 hydrostatic weighing scale after removing the surface roughness of the specimens through grinding. Prior to the Vickers microhardness tests and microstructure analysis, the samples were prepared according to the standard metallographic procedure and polished until the surface quality was 1 µm [15], considering a plane parallel to the build direction of the specimens. The Vickers microhardness was analyzed, in accordance with the ASTM E92-17, through a square-based pyramid indenter with an angle of 148°, performing at least three indentations on each specimen, with a load of 200g and a dwell time of 15s. Tensile tests were performed according to the ASTM E8M standard and considering the reduced specimen dimensions. The tests were carried out using a Galdabini Quasar 50 universal testing machine, equipped with a 50 kN load cell. Microstructure analysis was carried out after etching the polished specimens with the Keller's reagent by swabbing for at least 30s. Optical microscopy was carried out to take the macrographs of the specimens processed under different conditions at 100x magnification.

## Results and discussion

Density and mechanical properties. Fig. 3 illustrates the L-PBF specimens produced according to the experimental plan described in the previous section (only 1 replicate is shown). The visual inspection indicated that the specimens were sound and without significant defects. However, tensile specimens exhibited a significant distortion caused by the printing process, considering the remarkable thermal conductivity and thermal expansion coefficient of the selected alloy in combination with the horizontal positioning and the absence of stress relief heat treatment post-build.
<figure><img src="image_229.png" /><figcaption>Fig. 3 – L-PBF specimens produced in this work.</figcaption></figure>

On the other hand, the relative density measures results, illustrated in Fig. 4, suggested that the selected alloy was quite sensitive against the investigated processing conditions: more specifically, it was found that the highest densities were achieved for the processing conditions related to the lower values of power and scanning speed, regardless of the specific VED value, reaching a maximum value of approximately 97%. This result suggests that the chosen alloy needed a longer interaction time with the laser to enable a consistent melting without a remarkable presence of defects due to lack of fusion [16].

<figure><img src="image_230.png" /><figcaption>Fig. 4 – Relative density plot vs varied L-PBF process parameters.</figcaption></figure>

Concerning the mechanical properties, Table 3 resumes the collected data of Vickers microhardness and tensile properties, reporting for the sake of brevity only the experimental conditions that defined the investigated processing window as well as the central value. Another reason for this choice is that the experimental results were not suggesting remarkable variations for close processing conditions, coherently with the density experimental outcomes. This essay was confirmed also by the results shown in Table 3, that can be synthesized as follows: i) the
<figure><img src="image_233.png" /><figcaption>Fig. 3 – L-PBF specimens produced in this work.</figcaption></figure>

On the other hand, the relative density measures results, illustrated in Fig. 4, suggested that the selected alloy was quite sensitive against the investigated processing conditions: more specifically, it was found that the highest densities were achieved for the processing conditions related to the lower values of power and scanning speed, regardless of the specific VED value, reaching a maximum value of approximately 97%. This result suggests that the chosen alloy needed a longer interaction time with the laser to enable a consistent melting without a remarkable presence of defects due to lack of fusion [16].

<figure><img src="image_234.png" /><figcaption>Fig. 4 – Relative density plot vs varied L-PBF process parameters.</figcaption></figure>

Concerning the mechanical properties, Table 3 resumes the collected data of Vickers microhardness and tensile properties, reporting for the sake of brevity only the experimental conditions that defined the investigated processing window as well as the central value. Another reason for this choice is that the experimental results were not suggesting remarkable variations for close processing conditions, coherently with the density experimental outcomes. This essay was confirmed also by the results shown in Table 3, that can be synthesized as follows: i) the
microhardness was practically not sensitive to the processing conditions, as the material experienced a steep cooling rate in any case leading to a quenched-like condition after L-PBF; ii) the mechanical strength was affected by the L-PBF conditions, especially in terms of yield strength and elongation at break due to the presence of defects; iii) even if slight variations were observed, the tensile strength and Young's modulus were less sensitive against the L-PBF processing conditions with respect to the other tensile properties.

<table><caption>Table 3 – Vickers microhardness and mechanical properties of the L-PBF produced specimens that defined the extremes and the middle of the investigated processing window.</caption><thead><tr><th>ID</th><th>Vickers<br>microhardness<br>(HV)</th><th>Yield strength<br>(MPa)</th><th>Tensile<br>strength<br>(MPa)</th><th>Elongation at<br>break<br>(%)</th><th>Young's<br>modulus<br>(GPa)</th></tr></thead><tbody><tr><td>1</td><td>107±2.9</td><td>190±20</td><td>299±30</td><td>8.8±2</td><td>69.5±10</td></tr><tr><td>5</td><td>106±4.4</td><td>241±25</td><td>291±31</td><td>18.6±3.3</td><td>75.4±7</td></tr><tr><td>9</td><td>105±6.2</td><td>137±18</td><td>282±45</td><td>17.2±1.8</td><td>71.1±9</td></tr></tbody></table>

Microstructure and chemical composition. Fig. 5 illustrates the comparison between the cross-sectional views of the specimens ID # 1, 5 and 9 obtained according to the metallographic procedure described in the previous section. The first notable outcome was that the characteristic element that describes the microstructure of the L-PBF processed material, i.e. the molten pool, was not significantly affected in terms of shape as a function of varied processing conditions. This means that the selected processed parameters have not significantly changed the intensity of the laser-powder interaction, leading in any case to molten pools with a with greater than the depth, suggesting therefore the conduction welding mode [17]. Moreover, coherently with the density measurements, the produced specimens presented a notable number of defects, whose irregular shape suggested that the lack of fusion occurred in any case [18]. Despite this latter result, the optical macrographs did not suggest any correlation between the defects observed and the relative density quantified as a function of the process parameters.

<figure><img src="image_231.png" /><figcaption>Fig. 5 – Optical macrographs of the L-PBF specimens that defined the extremes and the middle of the investigated processing window (magnification 100x).</figcaption></figure>

For what concerns the influence of the selected processing conditions on the chemical composition of the alloy, the EDS results reported in Table 4 highlight that the original was slightly altered by the L-PBF process. However, the reported results could be affected by the specific area considered for the EDS analysis: as the L-PBF process leads to uncontrolled formation and
microhardness was practically not sensitive to the processing conditions, as the material experienced a steep cooling rate in any case leading to a quenched-like condition after L-PBF; ii) the mechanical strength was affected by the L-PBF conditions, especially in terms of yield strength and elongation at break due to the presence of defects; iii) even if slight variations were observed, the tensile strength and Young's modulus were less sensitive against the L-PBF processing conditions with respect to the other tensile properties.

<table><caption>Table 3 – Vickers microhardness and mechanical properties of the L-PBF produced specimens that defined the extremes and the middle of the investigated processing window.</caption><thead><tr><th>ID</th><th>Vickers<br>microhardness<br>(HV)</th><th>Yield strength<br>(MPa)</th><th>Tensile<br>strength<br>(MPa)</th><th>Elongation at<br>break<br>(%)</th><th>Young's<br>modulus<br>(GPa)</th></tr></thead><tbody><tr><td>1</td><td>107±2.9</td><td>190±20</td><td>299±30</td><td>8.8±2</td><td>69.5±10</td></tr><tr><td>5</td><td>106±4.4</td><td>241±25</td><td>291±31</td><td>18.6±3.3</td><td>75.4±7</td></tr><tr><td>9</td><td>105±6.2</td><td>137±18</td><td>282±45</td><td>17.2±1.8</td><td>71.1±9</td></tr></tbody></table>

Microstructure and chemical composition. Fig. 5 illustrates the comparison between the cross-sectional views of the specimens ID # 1, 5 and 9 obtained according to the metallographic procedure described in the previous section. The first notable outcome was that the characteristic element that describes the microstructure of the L-PBF processed material, i.e. the molten pool, was not significantly affected in terms of shape as a function of varied processing conditions. This means that the selected processed parameters have not significantly changed the intensity of the laser-powder interaction, leading in any case to molten pools with a with greater than the depth, suggesting therefore the conduction welding mode [17]. Moreover, coherently with the density measurements, the produced specimens presented a notable number of defects, whose irregular shape suggested that the lack of fusion occurred in any case [18]. Despite this latter result, the optical macrographs did not suggest any correlation between the defects observed and the relative density quantified as a function of the process parameters.

<figure><img src="image_235.png" /><figcaption>Fig. 5 – Optical macrographs of the L-PBF specimens that defined the extremes and the middle of the investigated processing window (magnification 100x).</figcaption></figure>

For what concerns the influence of the selected processing conditions on the chemical composition of the alloy, the EDS results reported in Table 4 highlight that the original was slightly altered by the L-PBF process. However, the reported results could be affected by the specific area considered for the EDS analysis: as the L-PBF process leads to uncontrolled formation and
distribution of intermetallic phases as well as to segregation of elements, it is convincing that these mechanisms make the EDS less robust against the region of interest.

<table><caption>Table 4 – Chemical composition of the produced L-PBF specimens that defined the extremes and the middle of the investigated processing window.</caption><thead><tr><th>Element (wt%)</th><th>Al</th><th>Mg</th><th>Sc</th><th>Zr</th></tr></thead><tbody><tr><td>ID 1</td><td>Bal.</td><td>4.2±0.1</td><td>0.4 ± 0.1</td><td>0.4 ± 0.2</td></tr><tr><td>ID 5</td><td>Bal.</td><td>4.8±0.1</td><td>0.4 ± 0.2</td><td>0.2 ± 0.1</td></tr><tr><td>ID 9</td><td>Bal.</td><td>4.1±0.2</td><td>0.3 ± 0.2</td><td>0.3 ± 0.2</td></tr></tbody></table>

## Conclusions

This work deals with the experimental investigation of the L-PBF processing of an Al-Mg-Sc-Zr alloy, exploring a wide frame of process parameters combination, followed by a characterization of main mechanical and microstructural properties of the parts in the as-built state. The emerged experimental outcomes could be synthesized as follows:

* Every combination of L-PBF processing parameters was feasible for the selected alloy, thus avoiding print interruptions and macroscopic flaws.

* The relative density was highly sensitive to the processing conditions investigated: it was found that higher values, up to approx. 98%, could be achieved when adopting lower laser power and scanning speed values. Vice versa, the lowest relative density of approx. the 88% was measured when the combination of the highest laser power and lowest scan speed was used.

* In general, the Vickers microhardness and the mechanical properties were not particularly affected by the different L-PBF process parameters, except for the yield strength and elongation at break whose values were dependent on the densification of the alloy.

* The microstructure of the processed alloy presented very similar characteristics, including the molten pools shape and the typical lack of fusion defects. The same consideration could be done for the chemical composition, for which a not negligible sensitivity of the EDS techniques might have introduced an uncertainty.

## References

[1] A. Katz-Demyanetz, V. V. Popov, A. Kovalevsky, D. Safranchik, A. Koptyug, Powder-bed additive manufacturing for aerospace application: Techniques, metallic and metal/ceramic composite materials and trends, Manuf Rev (Les Ulis) 6 (2019) 5. https://doi.org/10.1051/mfreview/2019003

[2] W. Abd-Elaziem, S. Elkatatny, A.-E. Abd-Elaziem, M. Khedr, M.A. Abd El-baky, M.A. Hassan, M. Abu-Okail, M. Mohammed, A. Järvenpää, T. Allam, A. Hamada, On the current research progress of metallic materials fabricated by laser powder bed fusion process: a review, Journal of Materials Research and Technology 20 (2022) 681-707. https://doi.org/10.1016/j.jmrt.2022.07.085

[3] A.G. Colomo, D. Wood, F. Martina, S.W. Williams, A comparison framework to support the selection of the best additive manufacturing process for specific aerospace applications, International Journal of Rapid Manufacturing 9 (2020) 194. https://doi.org/10.1504/IJRAPIDM.2020.107736
distribution of intermetallic phases as well as to segregation of elements, it is convincing that these mechanisms make the EDS less robust against the region of interest.

<table><caption>Table 4 – Chemical composition of the produced L-PBF specimens that defined the extremes and the middle of the investigated processing window.</caption><thead><tr><th>Element (wt%)</th><th>Al</th><th>Mg</th><th>Sc</th><th>Zr</th></tr></thead><tbody><tr><td>ID 1</td><td>Bal.</td><td>4.2±0.1</td><td>0.4 ± 0.1</td><td>0.4 ± 0.2</td></tr><tr><td>ID 5</td><td>Bal.</td><td>4.8±0.1</td><td>0.4 ± 0.2</td><td>0.2 ± 0.1</td></tr><tr><td>ID 9</td><td>Bal.</td><td>4.1±0.2</td><td>0.3 ± 0.2</td><td>0.3 ± 0.2</td></tr></tbody></table>

## Conclusions

This work deals with the experimental investigation of the L-PBF processing of an Al-Mg-Sc-Zr alloy, exploring a wide frame of process parameters combination, followed by a characterization of main mechanical and microstructural properties of the parts in the as-built state. The emerged experimental outcomes could be synthesized as follows:

* Every combination of L-PBF processing parameters was feasible for the selected alloy, thus avoiding print interruptions and macroscopic flaws.

* The relative density was highly sensitive to the processing conditions investigated: it was found that higher values, up to approx. 98%, could be achieved when adopting lower laser power and scanning speed values. Vice versa, the lowest relative density of approx. the 88% was measured when the combination of the highest laser power and lowest scan speed was used.

* In general, the Vickers microhardness and the mechanical properties were not particularly affected by the different L-PBF process parameters, except for the yield strength and elongation at break whose values were dependent on the densification of the alloy.

* The microstructure of the processed alloy presented very similar characteristics, including the molten pools shape and the typical lack of fusion defects. The same consideration could be done for the chemical composition, for which a not negligible sensitivity of the EDS techniques might have introduced an uncertainty.

## References

[1] A. Katz-Demyanetz, V. V. Popov, A. Kovalevsky, D. Safranchik, A. Koptyug, Powder-bed additive manufacturing for aerospace application: Techniques, metallic and metal/ceramic composite materials and trends, Manuf Rev (Les Ulis) 6 (2019) 5. https://doi.org/10.1051/mfreview/2019003

[2] W. Abd-Elaziem, S. Elkatatny, A.-E. Abd-Elaziem, M. Khedr, M.A. Abd El-baky, M.A. Hassan, M. Abu-Okail, M. Mohammed, A. Järvenpää, T. Allam, A. Hamada, On the current research progress of metallic materials fabricated by laser powder bed fusion process: a review, Journal of Materials Research and Technology 20 (2022) 681–707. https://doi.org/10.1016/j.jmrt.2022.07.085

[3] A.G. Colomo, D. Wood, F. Martina, S.W. Williams, A comparison framework to support the selection of the best additive manufacturing process for specific aerospace applications, International Journal of Rapid Manufacturing 9 (2020) 194. https://doi.org/10.1504/IJRAPIDM.2020.107736
[4] I. Tolosa, F. Garciandía, F. Zubiri, F. Zapirain, A. Esnaola, Study of mechanical properties of AISI 316 stainless steel processed by “selective laser melting”, following different manufacturing strategies, The International Journal of Advanced Manufacturing Technology 51 (2010) 639–647.
https://doi.org/10.1007/s00170-010-2631-5

[5] J.L. Leirmo, High Strength Aluminium Alloys in Laser-Based Powder Bed Fusion – a Review, Procedia CIRP 104 (2021) 1747–1752. https://doi.org/10.1016/j.procir.2021.11.294

[6] V.B. Vukkum, R.K. Gupta, Review on corrosion performance of laser powder-bed fusion printed 316L stainless steel: Effect of processing parameters, manufacturing defects, post-processing, feedstock, and microstructure, Mater Des 221 (2022) 110874.
https://doi.org/10.1016/j.matdes.2022.110874

[7] T.C. Dzogbewu, D.J. de Beer, Additive manufacturing of NiTi shape memory alloy and its industrial applications, Heliyon 10 (2024) e23369. https://doi.org/10.1016/j.heliyon.2023.e23369

[8] P. Rambabu, N. Eswara Prasad, V. V. Kutumbarao, R.J.H. Wanhill, Aluminium Alloys for Aerospace Applications, in: 2017: pp. 29–52. https://doi.org/10.1007/978-981-10-2134-3_2

[9] M. Javaid, A. Haleem, R.P. Singh, R. Suman, S. Rab, Role of additive manufacturing applications towards environmental sustainability, Advanced Industrial and Engineering Polymer Research 4 (2021) 312–322. https://doi.org/10.1016/j.aiepr.2021.07.005

[10] S.R. Narasimharaju, W. Zeng, T.L. See, Z. Zhu, P. Scott, X. Jiang, S. Lou, A comprehensive review on laser powder bed fusion of steels: Processing, microstructure, defects and control methods, mechanical properties, current challenges and future trends, J Manuf Process 75 (2022) 375–414. https://doi.org/10.1016/j.jmapro.2021.12.033

[11] K.L. Kendig, D.B. Miracle, Strengthening mechanisms of an Al-Mg-Sc-Zr alloy, Acta Mater 50 (2002) 4165–4175. https://doi.org/10.1016/S1359-6454(02)00258-6

[12] Information on https://www.metals4printing.com/home-en/, (n.d.).

[13] A. El Hassanin, A.T. Silvestri, F. Napolitano, D. Borrelli, A. Caraviello, A. Astarita, Investigation of the laser-related parameters in Laser-Powder Bed Fusion of Inconel 718-Cu blends at fixed Volumetric Energy Density, J Manuf Process 99 (2023) 456–468.
https://doi.org/10.1016/j.jmapro.2023.05.068

[14] A. El Hassanin, A.T. Silvestri, F. Napolitano, F. Scherillo, A. Caraviello, D. Borrelli, A. Astarita, Laser-powder bed fusion of pre-mixed Inconel 718-Cu powders: An experimental study, J Manuf Process 71 (2021) 329–344. https://doi.org/10.1016/j.jmapro.2021.09.028

[15] Kay Geels, Metallographic and Materialographic Specimen Preparation, Light Microscopy, Image Analysis, and Hardness Testing, n.d. https://books.google.it/books?id=oaehZy3Vo1kC

[16] A. du Plessis, Effects of process parameters on porosity in laser powder bed fusion revealed by X-ray tomography, Addit Manuf 30 (2019) 100871.
https://doi.org/10.1016/j.addma.2019.100871

[17] T. Ingrassia, V. Nigrelli, V. Ricotta, C. Tartamella, Process parameters influence in additive manufacturing, in: 2017: pp. 261–270. https://doi.org/10.1007/978-3-319-45781-9_27

[18] W. Wang, J. Ning, S.Y. Liang, Prediction of lack-of-fusion porosity in laser powder-bed fusion considering boundary conditions and sensitivity to laser power absorption, The International Journal of Advanced Manufacturing Technology 112 (2021) 61–70.
https://doi.org/10.1007/s00170-020-06224-7
[4] I. Tolosa, F. Garciandía, F. Zubiri, F. Zapirain, A. Esnaola, Study of mechanical properties of AISI 316 stainless steel processed by “selective laser melting”, following different manufacturing strategies, The International Journal of Advanced Manufacturing Technology 51 (2010) 639–647.
https://doi.org/10.1007/s00170-010-2631-5

[5] J.L. Leirmo, High Strength Aluminium Alloys in Laser-Based Powder Bed Fusion – a Review, Procedia CIRP 104 (2021) 1747–1752. https://doi.org/10.1016/j.procir.2021.11.294

[6] V.B. Vukkum, R.K. Gupta, Review on corrosion performance of laser powder-bed fusion printed 316L stainless steel: Effect of processing parameters, manufacturing defects, post-processing, feedstock, and microstructure, Mater Des 221 (2022) 110874.
https://doi.org/10.1016/j.matdes.2022.110874

[7] T.C. Dzogbewu, D.J. de Beer, Additive manufacturing of NiTi shape memory alloy and its industrial applications, Heliyon 10 (2024) e23369. https://doi.org/10.1016/j.heliyon.2023.e23369

[8] P. Rambabu, N. Eswara Prasad, V. V. Kutumbarao, R.J.H. Wanhill, Aluminium Alloys for Aerospace Applications, in: 2017: pp. 29–52. https://doi.org/10.1007/978-981-10-2134-3_2

[9] M. Javaid, A. Haleem, R.P. Singh, R. Suman, S. Rab, Role of additive manufacturing applications towards environmental sustainability, Advanced Industrial and Engineering Polymer Research 4 (2021) 312–322. https://doi.org/10.1016/j.aiepr.2021.07.005

[10] S.R. Narasimharaju, W. Zeng, T.L. See, Z. Zhu, P. Scott, X. Jiang, S. Lou, A comprehensive review on laser powder bed fusion of steels: Processing, microstructure, defects and control methods, mechanical properties, current challenges and future trends, J Manuf Process 75 (2022) 375–414. https://doi.org/10.1016/j.jmapro.2021.12.033

[11] K.L. Kendig, D.B. Miracle, Strengthening mechanisms of an Al-Mg-Sc-Zr alloy, Acta Mater 50 (2002) 4165–4175. https://doi.org/10.1016/S1359-6454(02)00258-6

[12] Information on https://www.metals4printing.com/home-en/, (n.d.).

[13] A. El Hassanin, A.T. Silvestri, F. Napolitano, D. Borrelli, A. Caraviello, A. Astarita, Investigation of the laser-related parameters in Laser-Powder Bed Fusion of Inconel 718-Cu blends at fixed Volumetric Energy Density, J Manuf Process 99 (2023) 456–468.
https://doi.org/10.1016/j.jmapro.2023.05.068

[14] A. El Hassanin, A.T. Silvestri, F. Napolitano, F. Scherillo, A. Caraviello, D. Borrelli, A. Astarita, Laser-powder bed fusion of pre-mixed Inconel 718-Cu powders: An experimental study, J Manuf Process 71 (2021) 329–344. https://doi.org/10.1016/j.jmapro.2021.09.028

[15] Kay Geels, Metallographic and Materialographic Specimen Preparation, Light Microscopy, Image Analysis, and Hardness Testing, n.d. https://books.google.it/books?id=oaehZy3Vo1kC

[16] A. du Plessis, Effects of process parameters on porosity in laser powder bed fusion revealed by X-ray tomography, Addit Manuf 30 (2019) 100871.
https://doi.org/10.1016/j.addma.2019.100871

[17] T. Ingrassia, V. Nigrelli, V. Ricotta, C. Tartamella, Process parameters influence in additive manufacturing, in: 2017: pp. 261–270. https://doi.org/10.1007/978-3-319-45781-9_27

[18] W. Wang, J. Ning, S.Y. Liang, Prediction of lack-of-fusion porosity in laser powder-bed fusion considering boundary conditions and sensitivity to laser power absorption, The International Journal of Advanced Manufacturing Technology 112 (2021) 61–70.
https://doi.org/10.1007/s00170-020-06224-7
# Corrective capabilities of different rescanning strategies to restore microstructure and density of initially porous 316L laser powder bed fusion

LE SAINT Théo¹,ª*, KELLER Clément¹,ᵇ, ABROUG Foued¹,ᶜ and ARNAUD Lionel¹,ᵈ

¹Laboratoire Génie de Production (LGP), Université de Technologie de Tarbes, France

ªtheo.le_saint@uttop.fr, ᵇclement.keller@uttop.fr, ᶜfoued.abroug@uttop.fr,

ᵈlionel.arnaud@uttop.fr

**Keywords:** LPBF, Rescanning, 316L, Defects, Microstructure, Characterization

**Abstract.** Mechanical properties of Laser Power Bed Fusion (LPBF) parts, and particularly fatigue properties, are heavily affected by defects including surface roughness or porosity. To mitigate the occurrence of these defects, in-situ and on-line corrective measures can be implemented to the fabrication process, among them, rescanning, which consists in remelting an already solidified layer. Initially porous LPBF samples were created and then rescanned using different scanning parameters and strategies. Results show that it is possible to regain part's health, compared to a standardly processed one, in terms of density, hardness and even improved roughness. This remelting process is known to refine microstructure of fabricated materials as well as reduce surface roughness and porosity without requiring further post-processing steps. Therefore, employing rescanning as a corrective technique appears to be a promising approach for rectifying detected defects during the fabrication process. The objective of this study is to assess the corrective capabilities of different rescanning strategies to restore the microstructure of an initially porous 316L LPBF simulating a defected part. This study shows that various rescanning strategies allow for densification of initially porous material from 98.83 ± 0.20 % to 99.75 ± 0.09 %, as well as lateral surface roughness reduction from Ra 20.2 ± 5.2 µm to Ra 12.7 ± 0.1 µm and microhardness increase from 243 ± 5 HV₀.₅ to 253 ± 3 HV₀.₅.

**Introduction**

Laser Powder Bed Fusion (LPBF) is a metallic additive manufacturing process, that consists in selectively melt thin layers of metallic powder on top of each other [1]. This process enables the manufacturing of geometrically complex parts, allowing for weight reduction in structures [2]. The thermal gradients and cooling rate occurring during fabrication produce a complex and often fine microstructure within the materials, leading to enhanced mechanical properties compared to their conventionally produced counterparts, particularly a higher yield strength, ultimate tensile strength and better ductility [3,4]. However, the mechanical properties, and especially the fatigue life, of LPBF parts are heavily affected by the presence of defects. This encompasses surface roughness, residual stress and porosity within the material such as lack of fusion, gas induced porosity or keyholes [5].

Hence, there is a need for corrective actions to mitigate defects, which can be in the form of post-processing steps. Grinding, polishing, milling or media blasting allow for surface roughness reduction [6] but are not applicable for complex geometry. Hot Isostatic Pressing can reduce porosity but coarsen the fine LPBF microstructure [7] and does not entirely eliminate fully enclosed porosity [8].

Another approach is in-situ monitoring of the fabrication through real-time vision, spectral, acoustic or thermal signal acquisition and treatment to either modify the process parameters on the spot with a close-loop control framework [9] or trigger a corrective response such as complete

<figure><img src="image_232.png" /></figure>
# Corrective capabilities of different rescanning strategies to restore microstructure and density of initially porous 316L laser powder bed fusion

LE SAINT Théo¹,ᵃ*, KELLER Clément¹,ᵇ, ABROUG Foued¹,ᶜ and ARNAUD Lionel¹,ᵈ

¹Laboratoire Génie de Production (LGP), Université de Technologie de Tarbes, France

ᵃtheo.le_saint@uttop.fr, ᵇclement.keller@uttop.fr, ᶜfoued.abroug@uttop.fr,

ᵈlionel.arnaud@uttop.fr

**Keywords:** LPBF, Rescanning, 316L, Defects, Microstructure, Characterization

**Abstract.** Mechanical properties of Laser Power Bed Fusion (LPBF) parts, and particularly fatigue properties, are heavily affected by defects including surface roughness or porosity. To mitigate the occurrence of these defects, in-situ and on-line corrective measures can be implemented to the fabrication process, among them, rescanning, which consists in remelting an already solidified layer. Initially porous LPBF samples were created and then rescanned using different scanning parameters and strategies. Results show that it is possible to regain part's health, compared to a standardly processed one, in terms of density, hardness and even improved roughness. This remelting process is known to refine microstructure of fabricated materials as well as reduce surface roughness and porosity without requiring further post-processing steps. Therefore, employing rescanning as a corrective technique appears to be a promising approach for rectifying detected defects during the fabrication process. The objective of this study is to assess the corrective capabilities of different rescanning strategies to restore the microstructure of an initially porous 316L LPBF simulating a defected part. This study shows that various rescanning strategies allow for densification of initially porous material from 98.83 ± 0.20 % to 99.75 ± 0.09 %, as well as lateral surface roughness reduction from Ra 20.2 ± 5.2 µm to Ra 12.7 ± 0.1 µm and microhardness increase from 243 ± 5 HV₀.₅ to 253 ± 3 HV₀.₅.

**Introduction**

Laser Powder Bed Fusion (LPBF) is a metallic additive manufacturing process, that consists in selectively melt thin layers of metallic powder on top of each other [1]. This process enables the manufacturing of geometrically complex parts, allowing for weight reduction in structures [2]. The thermal gradients and cooling rate occurring during fabrication produce a complex and often fine microstructure within the materials, leading to enhanced mechanical properties compared to their conventionally produced counterparts, particularly a higher yield strength, ultimate tensile strength and better ductility [3,4]. However, the mechanical properties, and especially the fatigue life, of LPBF parts are heavily affected by the presence of defects. This encompasses surface roughness, residual stress and porosity within the material such as lack of fusion, gas induced porosity or keyholes [5].

Hence, there is a need for corrective actions to mitigate defects, which can be in the form of post-processing steps. Grinding, polishing, milling or media blasting allow for surface roughness reduction [6] but are not applicable for complex geometry. Hot Isostatic Pressing can reduce porosity but coarsen the fine LPBF microstructure [7] and does not entirely eliminate fully enclosed porosity [8].

Another approach is in-situ monitoring of the fabrication through real-time vision, spectral, acoustic or thermal signal acquisition and treatment to either modify the process parameters on the spot with a close-loop control framework [9] or trigger a corrective response such as complete

<figure><img src="image_236.png" /></figure>
layer removal [10] or rescanning. This latter process consists in remelting an already solidified layer using the already existing laser, thus requiring no modification of the machine. Rescanning is known to reduce surface roughness [11-13], refine the microstructure features such as grain or sub-grain cell size [14,15], increase tensile properties [16,17] and on top, densify the parts [11,12,14,18]. Some studies have explored rescanning as a corrective action [18,19], nevertheless, the corrective capabilities of various rescanning strategies remain unexplored.

In this study, rescanning as an in-situ corrective response to defects is tested. Initially porous 316L samples are produced, simulating defected parts, and different rescanning strategies and powers are applied. The main objectives of this study are to identify the rescanning parameters best suited to regain the material density and to characterize the effect on the defect population. The impact of this corrective rescan on the surface roughness and microhardness of the sample is also assessed.

## Materials and Methods

All the specimens in this study have been fabricated with AISI 316L (LaserForm 316L (B)) produced by 3D Systems. This powder has a granulometric distribution $D_{10}$, $D_{50}$ and $D_{90}$ of 9.8, 20.1 and 38.1µm respectively. All the samples were produced using a 3D Systems ProX DMP 300B LPBF machine, the standard process parameters for this powder are listed in Table 1.

<table><caption>Table 1: Standard process parameters for 316L on ProX DMP 300.</caption><thead><tr><th>P [W]</th><th>V [mm/s]</th><th>Lt [µm]</th><th>H [µm]</th><th>Ev [J/mm³]</th><th>Rotation angle [°]</th><th>Scanning pattern</th><th>Contour</th></tr></thead><tbody><tr><td>215</td><td>1800</td><td>40</td><td>50</td><td>59.72</td><td>67</td><td>Bi-directional</td><td>n.a.</td></tr></tbody></table>

To assess the corrective capabilities of rescanning, 10×10×10mm³ cubes are fabricated with an initial low power scan to favor the apparition of natural defect and then different rescanning powers and strategies are applied. The initial scan is at 80% of the standard power, corresponding to 172W and the subsequent rescanning powers tested are 80%, 100%, 120% and 150% of the standard scanning power, respectively corresponding to 172W, 215W, 258W and 323W. The rescanning strategy is either SR: Superimposed Rescanning or PR: Perpendicular Rescanning, respectively 0° and 90° between the scan and the rescan, as illustrated in Fig. 1a. The nomenclature of the sample is detailed in Fig 1b. In addition, 2 samples are added, called 80 and 100, with only an initial scan of respectively 80% and 100% of scanning power, 80 being the reference sample with no correction rescanning applied and 100 being the standard sample as it is manufactured with standard parameters and thus acting as the correction aim of this study.

<figure><img src="image_233.png" /><figcaption>Figure 1: a) Schematics of the rescanning strategies tested in this study and b) nomenclature of the samples.</figcaption></figure>

All the samples were automatically polished using SiC paper up to 4000 grit. Image analysis was employed to determine the density of each sample. The Keyence VHX-5000 optical microscope captured assembly images of the entire samples at ×100 magnification. Each assembly
layer removal [10] or rescanning. This latter process consists in remelting an already solidified layer using the already existing laser, thus requiring no modification of the machine. Rescanning is known to reduce surface roughness [11-13], refine the microstructure features such as grain or sub-grain cell size [14,15], increase tensile properties [16,17] and on top, densify the parts [11,12,14,18]. Some studies have explored rescanning as a corrective action [18,19], nevertheless, the corrective capabilities of various rescanning strategies remain unexplored.

In this study, rescanning as an in-situ corrective response to defects is tested. Initially porous 316L samples are produced, simulating defected parts, and different rescanning strategies and powers are applied. The main objectives of this study are to identify the rescanning parameters best suited to regain the material density and to characterize the effect on the defect population. The impact of this corrective rescan on the surface roughness and microhardness of the sample is also assessed.

## Materials and Methods

All the specimens in this study have been fabricated with AISI 316L (LaserForm 316L (B)) produced by 3D Systems. This powder has a granulometric distribution $D_{10}$, $D_{50}$ and $D_{90}$ of 9.8, 20.1 and 38.1µm respectively. All the samples were produced using a 3D Systems ProX DMP 300B LPBF machine, the standard process parameters for this powder are listed in Table 1.

<table><caption>Table 1: Standard process parameters for 316L on ProX DMP 300.</caption><thead><tr><th>P [W]</th><th>V [mm/s]</th><th>Lt [µm]</th><th>H [µm]</th><th>Ev [J/mm³]</th><th>Rotation angle [°]</th><th>Scanning pattern</th><th>Contour</th></tr></thead><tbody><tr><td>215</td><td>1800</td><td>40</td><td>50</td><td>59.72</td><td>67</td><td>Bi-directional</td><td>n.a.</td></tr></tbody></table>

To assess the corrective capabilities of rescanning, 10×10×10mm³ cubes are fabricated with an initial low power scan to favor the apparition of natural defect and then different rescanning powers and strategies are applied. The initial scan is at 80% of the standard power, corresponding to 172W and the subsequent rescanning powers tested are 80%, 100%, 120% and 150% of the standard scanning power, respectively corresponding to 172W, 215W, 258W and 323W. The rescanning strategy is either SR: Superimposed Rescanning or PR: Perpendicular Rescanning, respectively 0° and 90° between the scan and the rescan, as illustrated in Fig. 1a. The nomenclature of the sample is detailed in Fig 1b. In addition, 2 samples are added, called 80 and 100, with only an initial scan of respectively 80% and 100% of scanning power, 80 being the reference sample with no correction rescanning applied and 100 being the standard sample as it is manufactured with standard parameters and thus acting as the correction aim of this study.

<figure><img src="image_237.png" /><figcaption>Figure 1: a) Schematics of the rescanning strategies tested in this study and b) nomenclature of the samples.</figcaption></figure>

All the samples were automatically polished using SiC paper up to 4000 grit. Image analysis was employed to determine the density of each sample. The Keyence VHX-5000 optical microscope captured assembly images of the entire samples at ×100 magnification. Each assembly
measured 13.1×11.6 mm, with a resolution of 2.16 µm/pixel, effectively covering the entirety of each sample. Subsequently, the micrographs were segmented using Matlab, allowing for density calculation and defect population characterization, such as defect size distribution. For those data treatments, only defects with an area above 10 pixels are treated.

The lateral surface roughness of the samples was assessed using micrographs of the radial sections of the samples at ×100 magnifications. The images were binarized and the profiles were defined as the sum of the most radially extreme pixels. Allowing to calculate the arithmetic mean Ra and the maximum peak to valley Rz of each sample. For each sample, Ra and Rz values were calculated over two profiles and of all the samples the shortest profile was 4.9 mm long.

To assess the impact of rescanning on mechanical properties, microhardness testing was carried out using a Zwick Roell DuraScan 70. For each sample, 60 HV0.5 indentations were performed in bulk zones of the samples with 0.5 mm spacing over x and y between indents.

## Results and discussions

The density of the samples for different rescanning powers and strategies is shown in Fig. 2a. The density exhibited by reference sample 80 (98.83 ± 0.20 %) is rather low compared to the density of the standard sample 100 (99.97 ± 0.01 %). The density evolution with rescanning power is nearly the same for both rescanning strategies. An initial density decrease is noted at 80% of rescanning power with a subsequent regain in density for 100% with a value of 99.75 ± 0.09 % for 80_SR100 (SR: Superimposed Rescanning). Above 100% of rescanning laser power, only a slight increase is noted for 80_SR150 with a density of 99.86 ± 0.02 %.

<figure><img src="image_234.png" /><figcaption>Figure 2: a) Density of samples versus rescanning power for different rescanning strategies, b) micrograph of 80_SR150.</figcaption></figure>

For rescanning powers of 120% and 150%, the edge part of the samples seems denser than the bulk part, as shown in Fig. 2b. This might be explained by the high thermal energy input of above 100% of rescanning power and a reduced thermal conductivity in edge zones compared to bulk zones due to the presence of powder.

When it comes to defect size, the distributions of the maximum Feret diameter of the porosities of the reference sample 80, 80_SR100 and the sample fabricated with standard process parameters 100 are shown in Fig 3, which also shows the total number of defects detected and the maximum Feret diameter of the biggest defect of each sample. Superimposed rescanning with 100% of rescanning power allowed for a decrease in the size of the biggest defect from 350 µm down to 229 µm. This defect size, although reduced, is still significantly bigger than the biggest defect of the standard sample 100, which exposes a biggest defect of 48 µm.
measured 13.1×11.6 mm, with a resolution of 2.16 µm/pixel, effectively covering the entirety of each sample. Subsequently, the micrographs were segmented using Matlab, allowing for density calculation and defect population characterization, such as defect size distribution. For those data treatments, only defects with an area above 10 pixels are treated.

The lateral surface roughness of the samples was assessed using micrographs of the radial sections of the samples at ×100 magnifications. The images were binarized and the profiles were defined as the sum of the most radially extreme pixels. Allowing to calculate the arithmetic mean Ra and the maximum peak to valley Rz of each sample. For each sample, Ra and Rz values were calculated over two profiles and of all the samples the shortest profile was 4.9 mm long.

To assess the impact of rescanning on mechanical properties, microhardness testing was carried out using a Zwick Roell DuraScan 70. For each sample, 60 HV0.5 indentations were performed in bulk zones of the samples with 0.5 mm spacing over x and y between indents.

## Results and discussions

The density of the samples for different rescanning powers and strategies is shown in Fig. 2a. The density exhibited by reference sample 80 (98.83 ± 0.20 %) is rather low compared to the density of the standard sample 100 (99.97 ± 0.01 %). The density evolution with rescanning power is nearly the same for both rescanning strategies. An initial density decrease is noted at 80% of rescanning power with a subsequent regain in density for 100% with a value of 99.75 ± 0.09 % for 80_SR100 (SR: Superimposed Rescanning). Above 100% of rescanning laser power, only a slight increase is noted for 80_SR150 with a density of 99.86 ± 0.02 %.

<figure><img src="image_238.png" /><figcaption>Figure 2: a) Density of samples versus rescanning power for different rescanning strategies, b) micrograph of 80_SR150.</figcaption></figure>

For rescanning powers of 120% and 150%, the edge part of the samples seems denser than the bulk part, as shown in Fig. 2b. This might be explained by the high thermal energy input of above 100% of rescanning power and a reduced thermal conductivity in edge zones compared to bulk zones due to the presence of powder.

When it comes to defect size, the distributions of the maximum Feret diameter of the porosities of the reference sample 80, 80_SR100 and the sample fabricated with standard process parameters 100 are shown in Fig 3, which also shows the total number of defects detected and the maximum Feret diameter of the biggest defect of each sample. Superimposed rescanning with 100% of rescanning power allowed for a decrease in the size of the biggest defect from 350 µm down to 229 µm. This defect size, although reduced, is still significantly bigger than the biggest defect of the standard sample 100, which exposes a biggest defect of 48 µm.
<figure><img src="image_235.png" /><figcaption>Figure 3: Maximum Feret diameter distribution, number of defects and maximum Feret diameter of the biggest defect of the defect population of a) 80, b) 80_SR100 and c) 100.</figcaption></figure>

Rescanning seems to also significantly decrease the number of defects. For the sample 80_SR100, the number of defects is 3.8 times lower than the reference sample 80. The number of defects present in 80_SR100 remains however 2.3 times higher than the standard sample 100. It can also be noted that the shape of the distribution of defect size is modified between 80 and 80_SR100. Particularly, the right side of the 80_SR100 distribution is lowered, revealing a greater impact of rescanning on the biggest defects of the whole population.

Regarding the surface roughness, the evolution of the lateral values of this parameter, Ra and Rz are shown in respectively Fig. 4a and Fig. 4b, for various rescanning powers and strategies. The reference sample 80 reveals higher lateral surface roughness values of Ra 20.2 ± 5.2 µm and Rz 190 ± 56 µm than the standard sample 100 with values of Ra and Rz of respectively 16.3 ± 1.7 µm and 161 ± 39 µm. With both rescanning strategies at 80% of rescanning power, the surface roughness is even higher. For all samples rescanned with 100% of rescanning power or above, except for the sample 80_SR120, Ra and Rz values are reduced compared to the reference sample 80, with a minimum reached at Ra 12.7 ± 0.1 µm and Rz 108 ± 6 µm for 80_PR120 (PR: Perpendicular Rescanning). It is worth mentioning that 80_SR150, 80_PR120 and 80_PR150 expose lateral roughness values even below the standard sample 100 ranges, meaning that rescanning might also be used to improve lateral surface roughness.

<figure><img src="image_236.png" /><figcaption>Figure 4: Lateral surface roughness a) Ra and b) Rz of samples versus rescanning power for different rescanning strategies.</figcaption></figure>
<figure><img src="image_239.png" /><figcaption>Figure 3: Maximum Feret diameter distribution, number of defects and maximum Feret diameter of the biggest defect of the defect population of a) 80, b) 80_SR100 and c) 100.</figcaption></figure>

Rescanning seems to also significantly decrease the number of defects. For the sample 80_SR100, the number of defects is 3.8 times lower than the reference sample 80. The number of defects present in 80_SR100 remains however 2.3 times higher than the standard sample 100. It can also be noted that the shape of the distribution of defect size is modified between 80 and 80_SR100. Particularly, the right side of the 80_SR100 distribution is lowered, revealing a greater impact of rescanning on the biggest defects of the whole population.

Regarding the surface roughness, the evolution of the lateral values of this parameter, Ra and Rz are shown in respectively Fig. 4a and Fig. 4b, for various rescanning powers and strategies. The reference sample 80 reveals higher lateral surface roughness values of Ra 20.2 ± 5.2 µm and Rz 190 ± 56 µm than the standard sample 100 with values of Ra and Rz of respectively 16.3 ± 1.7 µm and 161 ± 39 µm. With both rescanning strategies at 80% of rescanning power, the surface roughness is even higher. For all samples rescanned with 100% of rescanning power or above, except for the sample 80_SR120, Ra and Rz values are reduced compared to the reference sample 80, with a minimum reached at Ra 12.7 ± 0.1 µm and Rz 108 ± 6 µm for 80_PR120 (PR: Perpendicular Rescanning). It is worth mentioning that 80_SR150, 80_PR120 and 80_PR150 expose lateral roughness values even below the standard sample 100 ranges, meaning that rescanning might also be used to improve lateral surface roughness.

<figure><img src="image_240.png" /><figcaption>Figure 4: Lateral surface roughness a) Ra and b) Rz of samples versus rescanning power for different rescanning strategies.</figcaption></figure>
Finally, to assess the regain in mechanical properties along with densification, the samples are submitted to HV0.5 microhardness testing. The reference sample 80 exhibits a microhardness of 243 ± 5 HV₀.₅, lower than for the standard sample 100 with a value of 254 ± 2 HV₀.₅. Both rescanning strategies show a microhardness maximum for 100% of rescanning power with value for 80_SR100 and 80_PR100 respectively being 253 ± 3 HV₀.₅ and 252 ± 1 HV₀.₅. Those values are fairly equivalent to the one of the reference 100 sample. All other rescanned samples exhibit microhardness lower than the reference.

<figure><img src="image_237.png" /><figcaption>Figure 5: HV0.5 microhardness versus rescanning power for different rescanning strategies.</figcaption></figure>

Considering all the characterization performed in this study, it seems that laser power during rescanning must be equivalent to the one of the first laser scan to recover as much as possible the material properties of defected layers in LPBF process.

## Conclusions

Results show that rescanning as a defect correction measure is promising. The main findings of the study are as follows:

- Most of the density, surface roughness and microhardness are restored with superimposed rescanning at 100% of scanning power compared to the standard material. Particularly, between samples 80 and 80_SR100, the density is increased from 98.83 ± 0.20 % to 99.75 ± 0.09 % as well as the microhardness from 243 ± 5 HV₀.₅ to 253 ± 3 HV₀.₅ and the surface roughness is reduced from of Ra 20.2 ± 5.2 µm and Rz 190 ± 56 µm to Ra 18.8 ± 2.6 µm and Rz 153 ± 31 µm.

- Superimposed rescanning at 150% leads to even better density of 99.86 ± 0.02 % compared to superimposed rescanning at 100% as well as reduction in lateral surface roughness compared to the standard material. Indeed, 80_SR150 exposes lateral roughness of Ra 13.2 ± 1.9 µm and Rz 116 ± 2 µm, lower than the standard material 100 with Ra 16.3 ± 1.7 µm and 161 ± 39 µm. However, the microhardness is no fully restored.

- Superimposed or perpendicular rescanning strategies show similar results, with superimposed strategy being slightly better to restore materials properties.

## References

[1] W. E. Frazier, « Metal Additive Manufacturing: A Review », J. of Materi Eng and Perform, vol. 23, nº 6, p. 1917-1928, juin 2014. https://doi.org/10.1007/s11665-014-0958-z
Finally, to assess the regain in mechanical properties along with densification, the samples are submitted to HV0.5 microhardness testing. The reference sample 80 exhibits a microhardness of 243 ± 5 HV₀.₅, lower than for the standard sample 100 with a value of 254 ± 2 HV₀.₅. Both rescanning strategies show a microhardness maximum for 100% of rescanning power with value for 80_SR100 and 80_PR100 respectively being 253 ± 3 HV₀.₅ and 252 ± 1 HV₀.₅. Those values are fairly equivalent to the one of the reference 100 sample. All other rescanned samples exhibit microhardness lower than the reference.

<figure><img src="image_241.png" /><figcaption>Figure 5: HV0.5 microhardness versus rescanning power for different rescanning strategies.</figcaption></figure>

Considering all the characterization performed in this study, it seems that laser power during rescanning must be equivalent to the one of the first laser scan to recover as much as possible the material properties of defected layers in LPBF process.

## Conclusions

Results show that rescanning as a defect correction measure is promising. The main findings of the study are as follows:

- Most of the density, surface roughness and microhardness are restored with superimposed rescanning at 100% of scanning power compared to the standard material. Particularly, between samples 80 and 80_SR100, the density is increased from 98.83 ± 0.20 % to 99.75 ± 0.09 % as well as the microhardness from 243 ± 5 HV₀.₅ to 253 ± 3 HV₀.₅ and the surface roughness is reduced from of Ra 20.2 ± 5.2 µm and Rz 190 ± 56 µm to Ra 18.8 ± 2.6 µm and Rz 153 ± 31 µm.

- Superimposed rescanning at 150% leads to even better density of 99.86 ± 0.02 % compared to superimposed rescanning at 100% as well as reduction in lateral surface roughness compared to the standard material. Indeed, 80_SR150 exposes lateral roughness of Ra 13.2 ± 1.9 µm and Rz 116 ± 2 µm, lower than the standard material 100 with Ra 16.3 ± 1.7 µm and 161 ± 39 µm. However, the microhardness is no fully restored.

- Superimposed or perpendicular rescanning strategies show similar results, with superimposed strategy being slightly better to restore materials properties.

## References

[1] W. E. Frazier, « Metal Additive Manufacturing: A Review », J. of Materi Eng and Perform, vol. 23, nº 6, p. 1917-1928, juin 2014. https://doi.org/10.1007/s11665-014-0958-z
[2] G. Schuh, G. Bergweiler, K. Lichtenthäler, F. Fiedler, et S. de la Puente Rebollo, « Topology Optimisation and Metal Based Additive Manufacturing of Welding Jig Elements », Procedia CIRP, vol. 93, p. 62-67, janv. 2020. https://doi.org/10.1016/j.procir.2020.04.066

[3] Y. M. Wang et al., « Additively manufactured hierarchical stainless steels with high strength and ductility », Nature Mater, vol. 17, n° 1, Art. n° 1, janv. 2018. https://doi.org/10.1038/nmat5021

[4] L. Liu et al., « Dislocation network in additive manufactured steel breaks strength-ductility trade-off », Materials Today, vol. 21, n° 4, p. 354-361, mai 2018. https://doi.org/10.1016/j.mattod.2017.11.004

[5] S. Chowdhury et al., « Laser powder bed fusion: a state-of-the-art review of the technology, materials, properties & defects, and numerical modelling », Journal of Materials Research and Technology, vol. 20, p. 2109-2172, sept. 2022. https://doi.org/10.1016/j.jmrt.2022.07.121

[6] H. M. Khan, Y. Karabulut, O. Kitay, Y. Kaynak, et I. S. Jawahir, « Influence of the post-processing operations on surface integrity of metal components produced by laser powder bed fusion additive manufacturing: a review », Machining Science and Technology, vol. 25, n° 1, p. 118-176, déc. 2020. https://doi.org/10.1080/10910344.2020.1855649

[7] I. S. Grech, J. H. Sullivan, R. J. Lancaster, J. Plummer, et N. P. Lavery, « The optimisation of hot isostatic pressing treatments for enhanced mechanical and corrosion performance of stainless steel 316L produced by laser powder bed fusion », Additive Manufacturing, vol. 58, p. 103072, oct. 2022. https://doi.org/10.1016/j.addma.2022.103072

[8] S. Tammas-Williams, P. J. Withers, I. Todd, et P. B. Prangnell, « Porosity regrowth during heat treatment of hot isostatically pressed additively manufactured titanium components », Scripta Materialia, vol. 122, p. 72-76, sept. 2016. https://doi.org/10.1016/j.scriptamat.2016.05.002

[9] Y. Cai, J. Xiong, H. Chen, et G. Zhang, « A review of in-situ monitoring and process control system in metal-based laser additive manufacturing », Journal of Manufacturing Systems, vol. 70, p. 309-326, oct. 2023. https://doi.org/10.1016/j.jmsy.2023.07.018

[10] B. M. Colosimo, E. Grossi, F. Caltanissetta, et M. Grasso, « Penelope: A Novel Prototype for In Situ Defect Removal in LPBF », JOM, vol. 72, n° 3, p. 1332-1339, mars 2020. https://doi.org/10.1007/s11837-019-03964-0

[11] B. Liu, B.-Q. Li, et Z. Li, « Selective laser remelting of an additive layer manufacturing process on AlSi10Mg », Results in Physics, vol. 12, p. 982-988, mars 2019. https://doi.org/10.1016/j.rinp.2018.12.018

[12] W. Yu, S. L. Sing, C. K. Chua, et X. Tian, « Influence of re-melting on surface roughness and porosity of AlSi10Mg parts fabricated by selective laser melting », Journal of Alloys and Compounds, vol. 792, p. 574-581, juill. 2019. https://doi.org/10.1016/j.jallcom.2019.04.017

[13] J.-P. Kruth, M. Badrossamay, E. Yasa, J. Deckers, L. Thijs, et J. Humbeeck, « Part and material properties in selective laser melting of metals », 16th International Symposium on Electromachining, ISEM 2010, janv. 2010.

[14] A. Liang, K. S. Pey, T. Polcar, et A. R. Hamilton, « Effects of rescanning parameters on densification and microstructural refinement of 316L stainless steel fabricated by laser powder bed fusion », Journal of Materials Processing Technology, vol. 302, p. 117493, avr. 2022. https://doi.org/10.1016/j.jmatprotec.2022.117493

[15] S. Griffiths, M. D. Rossell, J. Croteau, N. Q. Vo, D. C. Dunand, et C. Leinenbach, « Effect of laser rescanning on the grain microstructure of a selective laser melted Al-Mg-Zr alloy »,
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[3] Y. M. Wang et al., « Additively manufactured hierarchical stainless steels with high strength and ductility », Nature Mater, vol. 17, n° 1, Art. n° 1, janv. 2018. https://doi.org/10.1038/nmat5021

[4] L. Liu et al., « Dislocation network in additive manufactured steel breaks strength-ductility trade-off », Materials Today, vol. 21, n° 4, p. 354-361, mai 2018. https://doi.org/10.1016/j.mattod.2017.11.004

[5] S. Chowdhury et al., « Laser powder bed fusion: a state-of-the-art review of the technology, materials, properties & defects, and numerical modelling », Journal of Materials Research and Technology, vol. 20, p. 2109-2172, sept. 2022. https://doi.org/10.1016/j.jmrt.2022.07.121

[6] H. M. Khan, Y. Karabulut, O. Kitay, Y. Kaynak, et I. S. Jawahir, « Influence of the post-processing operations on surface integrity of metal components produced by laser powder bed fusion additive manufacturing: a review », Machining Science and Technology, vol. 25, n° 1, p. 118-176, déc. 2020. https://doi.org/10.1080/10910344.2020.1855649

[7] I. S. Grech, J. H. Sullivan, R. J. Lancaster, J. Plummer, et N. P. Lavery, « The optimisation of hot isostatic pressing treatments for enhanced mechanical and corrosion performance of stainless steel 316L produced by laser powder bed fusion », Additive Manufacturing, vol. 58, p. 103072, oct. 2022. https://doi.org/10.1016/j.addma.2022.103072

[8] S. Tammas-Williams, P. J. Withers, I. Todd, et P. B. Prangnell, « Porosity regrowth during heat treatment of hot isostatically pressed additively manufactured titanium components », Scripta Materialia, vol. 122, p. 72-76, sept. 2016. https://doi.org/10.1016/j.scriptamat.2016.05.002

[9] Y. Cai, J. Xiong, H. Chen, et G. Zhang, « A review of in-situ monitoring and process control system in metal-based laser additive manufacturing », Journal of Manufacturing Systems, vol. 70, p. 309-326, oct. 2023. https://doi.org/10.1016/j.jmsy.2023.07.018

[10] B. M. Colosimo, E. Grossi, F. Caltanissetta, et M. Grasso, « Penelope: A Novel Prototype for In Situ Defect Removal in LPBF », JOM, vol. 72, n° 3, p. 1332-1339, mars 2020. https://doi.org/10.1007/s11837-019-03964-0

[11] B. Liu, B.-Q. Li, et Z. Li, « Selective laser remelting of an additive layer manufacturing process on AlSi10Mg », Results in Physics, vol. 12, p. 982-988, mars 2019. https://doi.org/10.1016/j.rinp.2018.12.018

[12] W. Yu, S. L. Sing, C. K. Chua, et X. Tian, « Influence of re-melting on surface roughness and porosity of AlSi10Mg parts fabricated by selective laser melting », Journal of Alloys and Compounds, vol. 792, p. 574-581, juill. 2019. https://doi.org/10.1016/j.jallcom.2019.04.017

[13] J.-P. Kruth, M. Badrossamay, E. Yasa, J. Deckers, L. Thijs, et J. Humbeeck, « Part and material properties in selective laser melting of metals », 16th International Symposium on Electromachining, ISEM 2010, janv. 2010.

[14] A. Liang, K. S. Pey, T. Polcar, et A. R. Hamilton, « Effects of rescanning parameters on densification and microstructural refinement of 316L stainless steel fabricated by laser powder bed fusion », Journal of Materials Processing Technology, vol. 302, p. 117493, avr. 2022. https://doi.org/10.1016/j.jmatprotec.2022.117493

[15] S. Griffiths, M. D. Rossell, J. Croteau, N. Q. Vo, D. C. Dunand, et C. Leinenbach, « Effect of laser rescanning on the grain microstructure of a selective laser melted Al-Mg-Zr alloy »,
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[17] T. Han et al., « Refined microstructure and enhanced mechanical properties of AlCrFe2Ni2 medium entropy alloy produced via laser remelting », *Journal of Materials Science & Technology*, vol. 99, p. 18-27, févr. 2022. https://doi.org/10.1016/j.jmst.2021.05.033

[18] E. Yasa et J.-P. Kruth, « Microstructural investigation of Selective Laser Melting 316L stainless steel parts exposed to laser re-melting », *Procedia Engineering*, vol. 19, p. 389-395, janv. 2011. https://doi.org/10.1016/j.proeng.2011.11.130

[19] J. Mireles, S. Ridwan, P. Morton, A. Hinojos, et R. Wicker, « Analysis and correction of defects within parts fabricated using powder bed fusion technology », *Surface Topography: Metrology and Properties*, vol. 3, p. 034002, août 2015. https://doi.org/10.1088/2051-672X/3/3/034002

[20] B. Shen et al., « Influence of laser post-processing on pore evolution of Ti-6Al-4V alloy by laser powder bed fusion », *Journal of Alloys and Compounds*, vol. 818, p. 152845, mars 2020. https://doi.org/10.1016/j.jallcom.2019.152845
# Investigating the effect of the addition of WC into NiTi for stent application

AGARWAL Neha¹,²,a,*, PERUMAL Gopinath¹,b, OBEIDI Muhannad Ahmed¹,²,c, BRABAZON Dermot¹,²,d

¹School of Mechanical and Manufacturing Engineering, Dublin City University, Dublin 9, Ireland

²I-Form Advanced Manufacturing Research Centre, Dublin City University, Dublin 9, Ireland

a.neha.agarwal2@mail.dcu.ie, b.gopinath.perumal@dcu.ie, c.muhannad.ahmedobeidi@dcu.ie,

d.dermot.brabazon@dcu.ie

**Keywords:** Powder Alloying, Radiopacity, Compaction, Sintering, Hardness

**Abstract.** Nitinol is used for self-expandable stents used in biomedical applications considering their excellent properties of being superelastic and biocompatible. Minimally invasive implants need to be securely positioned in the human body with the property of being visible with the help of X-rays. Limited radiopacity of nitinol implants requires additional markers on the implants to facilitate accurate placement of the implant. This study explores the path of adding different concentrations of tungsten carbide (WC) to nickel-rich nitinol powder. Tests were conducted by compacting the mixed powders with a hydraulic press and then furnace sintering the green samples, with varying pressures and temperatures, to produce the solid parts. The mechanical properties of density and hardness of the sintered samples were measured and compared. An increase in hardness from 76 HV to 150 HV indicated that better mechanical properties could be achieved. Fully dense samples without any pores were obtained. These results show potential for the introduction of WC to nitinol for additively manufactured parts.

**Introduction**

Nitinol is a shape-memory alloy made of nickel and titanium, both of which are present in near-equal amounts. Recent advances in biomedical applications have been seen to be utilized as catheter tubes, guide wires, stone retrieval baskets, filters, needles, and stents [1]-[4]. Due to its excellent properties of shape memory alloy effect, superelasticity, high corrosion resistance, good biocompatibility, and good ductility, it is a desirable choice for implant applications [5]-[7]. However, nitinol poses the challenge of being radiolucent, which induces the challenge of the precise placement of the device in the human body. To combat this challenge and enhance other properties, the alloying of nitinol was evaluated. Among many other options erbium (Er) [8], tantalum (Ta) [9], and tungsten carbide (WC) [10] are often combined with other metals to enhance the mechanical properties of the primary material. WC possesses many important features that make it desirable for use in stents such as its high hardness, resistance to corrosion, high density (when compared to NiTi), and biocompatibility. As reported by Mussatto [11], adding WC into stainless steel during the laser bed fusion process resulted in increased strength of the parts produced with a stronger crystallographic texture.

In this present study, we evaluate the possibility of the addition of WC in nickel-rich nitinol to investigate the hardness and density of the new composition. The samples have been produced by the compaction process of the powder using a hydraulic press with subsequent heat sintering to form solid circular disks. The addition of WC into NiTi for enhancing the mechanical property of hardness was evaluated. The addition of different materials in NiTi opens the possibility of assessing several different applications with desired functional properties. Small varying mass

<figure><img src="image_238.png" /></figure>
# Investigating the effect of the addition of WC into NiTi for stent application

AGARWAL Neha¹,²,a,*, PERUMAL Gopinath¹,b, OBEIDI Muhannad Ahmed¹,²,c, BRABAZON Dermot¹,²,d

¹School of Mechanical and Manufacturing Engineering, Dublin City University, Dublin 9, Ireland

²I-Form Advanced Manufacturing Research Centre, Dublin City University, Dublin 9, Ireland

a.neha.agarwal2@mail.dcu.ie, b.gopinath.perumal@dcu.ie, c.muhannad.ahmedobeidi@dcu.ie,

d.dermot.brabazon@dcu.ie

**Keywords:** Powder Alloying, Radiopacity, Compaction, Sintering, Hardness

**Abstract.** Nitinol is used for self-expandable stents used in biomedical applications considering their excellent properties of being superelastic and biocompatible. Minimally invasive implants need to be securely positioned in the human body with the property of being visible with the help of X-rays. Limited radiopacity of nitinol implants requires additional markers on the implants to facilitate accurate placement of the implant. This study explores the path of adding different concentrations of tungsten carbide (WC) to nickel-rich nitinol powder. Tests were conducted by compacting the mixed powders with a hydraulic press and then furnace sintering the green samples, with varying pressures and temperatures, to produce the solid parts. The mechanical properties of density and hardness of the sintered samples were measured and compared. An increase in hardness from 76 HV to 150 HV indicated that better mechanical properties could be achieved. Fully dense samples without any pores were obtained. These results show potential for the introduction of WC to nitinol for additively manufactured parts.

**Introduction**

Nitinol is a shape-memory alloy made of nickel and titanium, both of which are present in near-equal amounts. Recent advances in biomedical applications have been seen to be utilized as catheter tubes, guide wires, stone retrieval baskets, filters, needles, and stents [1]-[4]. Due to its excellent properties of shape memory alloy effect, superelasticity, high corrosion resistance, good biocompatibility, and good ductility, it is a desirable choice for implant applications [5]-[7]. However, nitinol poses the challenge of being radiolucent, which induces the challenge of the precise placement of the device in the human body. To combat this challenge and enhance other properties, the alloying of nitinol was evaluated. Among many other options erbium (Er) [8], tantalum (Ta) [9], and tungsten carbide (WC) [10] are often combined with other metals to enhance the mechanical properties of the primary material. WC possesses many important features that make it desirable for use in stents such as its high hardness, resistance to corrosion, high density (when compared to NiTi), and biocompatibility. As reported by Mussatto [11], adding WC into stainless steel during the laser bed fusion process resulted in increased strength of the parts produced with a stronger crystallographic texture.

In this present study, we evaluate the possibility of the addition of WC in nickel-rich nitinol to investigate the hardness and density of the new composition. The samples have been produced by the compaction process of the powder using a hydraulic press with subsequent heat sintering to form solid circular disks. The addition of WC into NiTi for enhancing the mechanical property of hardness was evaluated. The addition of different materials in NiTi opens the possibility of assessing several different applications with desired functional properties. Small varying mass

<figure><img src="image_242.png" /></figure>
percentages of WC were added to NiTi to enhance the strength of the combined powder with a view to not significantly alter the shape memory properties of the produced NiTi composite part.

## Materials and Methods

### Powder Mixing and Compaction

Pre-alloyed nickel titanium powder was sourced from Fort Wayne Metals, Castlebar, Ireland with an elemental composition of 52.39 at.% Ni and 47.61 at.% Ti with a purity of 99.999%. The particle size distribution of NiTi powder is 12.3 µm (D₁₀), 27.4 µm (D₅₀) and 47.6 µm (D₉₀). A hydraulic press was used for the compaction process of the NiTi powder. The compaction parameters applied were the compaction pressure (ton), holding time (min), and varying powder mass to achieve highly compact powder to produce disc samples geometry of 20 mm diameter. Although the best results were achieved from the compaction of the pre-alloyed NiTi powder at 2.4 MPa, a holding time of 15 mins, and a NiTi powder mass of 2.4g; however, it was still not sufficiently compacted enough to continue onto the sintering process as shown in Figure 1.

<figure><img src="image_239.png" /><figcaption>Figure 1: Pre-alloyed NiTi powder compacted at 22 tons, holding time of 15 mins and mass of 2.4 g.</figcaption></figure>

Due to the failure of the achieving fully compact discs, separate nickel and titanium powders sourced from Fort Wayne Metals, Castlebar, Ireland was used with a purity of 99.996%. A Resodyn LabRam1 acoustic mixer was used to thoroughly mix the sample powder. A three-step mixing process was used with each step lasting 30 seconds with increased acceleration gravity (g) levels of 40g, 60g, and 70g in each step. These newly mixed samples of NiTi were then tested for compaction under different parametersand instant improvements to the samples post compaction were obtained. However, there were microcracks observed on the samples. A preliminary pilot test to investigate the optimum set of mixing ratios, it was found that the addition of water soluble 2% PVA (Polyvinyl Alcohol) to the powder mixture and compaction at a load of 25 tons, holding time of 15 mins, and the mass of the powder at 6.5g to achieve a significant and improved compaction results followed by efficient removal of the parts from the die as shown in Figure 2.

<figure><img src="image_240.png" /><figcaption>Figure 2: Self-mixed NiTi powder was compacted using the parameters of load 25 tons, hold time of 15 mins, mass of powder 6.5g and PVA added at 2% of the mass of the NiTi powder.</figcaption></figure>
percentages of WC were added to NiTi to enhance the strength of the combined powder with a view to not significantly alter the shape memory properties of the produced NiTi composite part.

## Materials and Methods

### Powder Mixing and Compaction

Pre-alloyed nickel titanium powder was sourced from Fort Wayne Metals, Castlebar, Ireland with an elemental composition of 52.39 at.% Ni and 47.61 at.% Ti with a purity of 99.999%. The particle size distribution of NiTi powder is 12.3 µm (D₁₀), 27.4 µm (D₅₀) and 47.6 µm (D₉₀). A hydraulic press was used for the compaction process of the NiTi powder. The compaction parameters applied were the compaction pressure (ton), holding time (min), and varying powder mass to achieve highly compact powder to produce disc samples geometry of 20 mm diameter. Although the best results were achieved from the compaction of the pre-alloyed NiTi powder at 2.4 MPa, a holding time of 15 mins, and a NiTi powder mass of 2.4g; however, it was still not sufficiently compacted enough to continue onto the sintering process as shown in Figure 1.

<figure><img src="image_243.png" /><figcaption>Figure 1: Pre-alloyed NiTi powder compacted at 22 tons, holding time of 15 mins and mass of 2.4 g.</figcaption></figure>

Due to the failure of the achieving fully compact discs, separate nickel and titanium powders sourced from Fort Wayne Metals, Castlebar, Ireland was used with a purity of 99.996%. A Resodyn LabRam1 acoustic mixer was used to thoroughly mix the sample powder. A three-step mixing process was used with each step lasting 30 seconds with increased acceleration gravity (g) levels of 40g, 60g, and 70g in each step. These newly mixed samples of NiTi were then tested for compaction under different parametersand instant improvements to the samples post compaction were obtained. However, there were microcracks observed on the samples. A preliminary pilot test to investigate the optimum set of mixing ratios, it was found that the addition of water soluble 2% PVA (Polyvinyl Alcohol) to the powder mixture and compaction at a load of 25 tons, holding time of 15 mins, and the mass of the powder at 6.5g to achieve a significant and improved compaction results followed by efficient removal of the parts from the die as shown in Figure 2.

<figure><img src="image_244.png" /><figcaption>Figure 2: Self-mixed NiTi powder was compacted using the parameters of load 25 tons, hold time of 15 mins, mass of powder 6.5g and PVA added at 2% of the mass of the NiTi powder.</figcaption></figure>
Once the parameters that are best suited for the compaction of NiTi was found, WC was mixed into the powdered samples at increasing mass percentages of 5%, 10%, 15% and 20% of the total mass of the powder mixed according to Table 1. These self-mixed NiTi powders and WC were then compacted under the previously found parameters for compaction of the self-mixed NiTi powder.

<table><caption>Table 1: Varying masses of WC added to self-mixed NiTi powders for compaction.</caption><thead><tr><th colspan="3">Mass division of WC and NiTi</th></tr><tr><th>WC %</th><th>NiTi (g)</th><th>WC (g)</th></tr></thead><tbody><tr><td>0</td><td>40</td><td>0</td></tr><tr><td>5</td><td>38</td><td>2</td></tr><tr><td>10</td><td>36</td><td>4</td></tr><tr><td>15</td><td>34</td><td>6</td></tr><tr><td>20</td><td>32</td><td>8</td></tr></tbody></table>

## Sintering

Sintering was done to each of the compacted samples with the use of a heat furnace. Due to the high melting point of NiTi, heat sintering method was employed to keep the purity of the samples post processing. Pre-testing was conducted for the heat sintering of the compacted samples by holding the samples inside the heat surface for 3 hours with a temperature 1200 °C (as this was between the ideal sintering temperature between nickel (1091 °C) and titanium (1251 °C)). However, this resulted in the samples tested completely melted and incapable of any use. The samples were then sintered in steps until the heat furnace reached a temperature of 1000 °C holding it for 3 hours before the heat furnace was cooled down in the same steps by which it was heated. Figure 3 shows the sintering process used for the samples. The samples post sintering under these steps were seen to produce ideal samples.

<figure><img src="image_241.png" /><figcaption>Figure 3: Sintering steps conducted until the heat furnace reached 1000 °C with holding time for 3 hours before it was cooled down in steps.</figcaption></figure>

## Metallographic preparation of the samples

All the sintered samples were mounted in resin for ease of use during the grinding and polishing process of the surface. The samples were grinded using abrasive paper that went up in grit size value from P80 to P4000. Once all samples were grinded diamond polishing was done to the
Once the parameters that are best suited for the compaction of NiTi was found, WC was mixed into the powdered samples at increasing mass percentages of 5%, 10%, 15% and 20% of the total mass of the powder mixed according to Table 1. These self-mixed NiTi powders and WC were then compacted under the previously found parameters for compaction of the self-mixed NiTi powder.

<table><caption>Table 1: Varying masses of WC added to self-mixed NiTi powders for compaction.</caption><thead><tr><th colspan="3">Mass division of WC and NiTi</th></tr><tr><th>WC %</th><th>NiTi (g)</th><th>WC (g)</th></tr></thead><tbody><tr><td>0</td><td>40</td><td>0</td></tr><tr><td>5</td><td>38</td><td>2</td></tr><tr><td>10</td><td>36</td><td>4</td></tr><tr><td>15</td><td>34</td><td>6</td></tr><tr><td>20</td><td>32</td><td>8</td></tr></tbody></table>

## Sintering

Sintering was done to each of the compacted samples with the use of a heat furnace. Due to the high melting point of NiTi, heat sintering method was employed to keep the purity of the samples post processing. Pre-testing was conducted for the heat sintering of the compacted samples by holding the samples inside the heat surface for 3 hours with a temperature 1200 °C (as this was between the ideal sintering temperature between nickel (1091 °C) and titanium (1251 °C)). However, this resulted in the samples tested completely melted and incapable of any use. The samples were then sintered in steps until the heat furnace reached a temperature of 1000 °C holding it for 3 hours before the heat furnace was cooled down in the same steps by which it was heated. Figure 3 shows the sintering process used for the samples. The samples post sintering under these steps were seen to produce ideal samples.

<figure><img src="image_245.png" /><figcaption>Figure 3: Sintering steps conducted until the heat furnace reached 1000 °C with holding time for 3 hours before it was cooled down in steps.</figcaption></figure>

## Metallographic preparation of the samples

All the sintered samples were mounted in resin for ease of use during the grinding and polishing process of the surface. The samples were grinded using abrasive paper that went up in grit size value from P80 to P4000. Once all samples were grinded diamond polishing was done to the
surface of the samples to achieve a smooth mirror-like finish on all samples to make them available for testing as shown in Figure 4. The samples were polished using diamond suspensions in decreasing number of particle sizes from 9µm, 6µm, 3µm and 1µm.

<figure><img src="image_242.png" /><figcaption>Figure 4: Polished sample surface of NiTi + 10% WC.</figcaption></figure>

## Characterization Techniques

Morphological characterization was conducted using a scanning electron microscope (SEM), Zeiss EVO LS-15, on all powders which comprised of pre-alloyed NiTi, self-mixed NiTi, and NiTi with varying amounts of WC. This inspection was conducted to analyze the even mixing of all powders before compaction. Energy dispersive X-ray spectroscopy (EDS), INCA, Oxford Instruments EDX system was used in combination with SEM to carry out an elemental analysis on all the samples being used. This technique helped to understand the composition of all powdered and sintered samples by showing the average mass percentage value of each element present within each sample.

As there were many defects found within the topography of the sintered and polished samples it was almost impossible to see the composition of the material using optical microscopy, thus backscattered images were used to get the structural composition of the polished samples. This was done with the addition of the backscatter sensor to the SEM directly above the sample which detects backscattered electrons from the sample onto the sensor and produces high-quality images of the microstructure of the sample a small distance inside the sample being viewed. EDS was then used to map out the different elements present within the sample shown by each color contrast specific to a different element.

Vickers Hardness test was carried out on each of the samples to obtain a value for hardness on each of these samples. A small indentation is made on each of the samples with a load of 0.98 N with a holding time of 20 secs. The average Vickers hardness was reported after 5 measurements were done on each sample. Density measurements were carried out on all the sintered samples using the Archimedes principle. The measurements were carried out using ethanol as fluid. Three measurements were conducted, and the average density has been reported.

## Results and Discussion

### Morphological characteristics of the powder samples

The SEM metallography images show clear structural difference between the NiTi particles and WC particles as shown in Figure 5 and Figure 6, respectively. The NiTi powder showed particles were circular and round consistently throughout the sample with little to no impurities or defects to the particles. On the other hand, WC was observed to have a much smaller particle size (ranging from 1-5 µm) and was very random in shape which is a commonly supplied form for this material [12].
surface of the samples to achieve a smooth mirror-like finish on all samples to make them available for testing as shown in Figure 4. The samples were polished using diamond suspensions in decreasing number of particle sizes from 9µm, 6µm, 3µm and 1µm.

<figure><img src="image_246.png" /><figcaption>Figure 4: Polished sample surface of NiTi + 10% WC.</figcaption></figure>

## Characterization Techniques

Morphological characterization was conducted using a scanning electron microscope (SEM), Zeiss EVO LS-15, on all powders which comprised of pre-alloyed NiTi, self-mixed NiTi, and NiTi with varying amounts of WC. This inspection was conducted to analyze the even mixing of all powders before compaction. Energy dispersive X-ray spectroscopy (EDS), INCA, Oxford Instruments EDX system was used in combination with SEM to carry out an elemental analysis on all the samples being used. This technique helped to understand the composition of all powdered and sintered samples by showing the average mass percentage value of each element present within each sample.

As there were many defects found within the topography of the sintered and polished samples it was almost impossible to see the composition of the material using optical microscopy, thus backscattered images were used to get the structural composition of the polished samples. This was done with the addition of the backscatter sensor to the SEM directly above the sample which detects backscattered electrons from the sample onto the sensor and produces high-quality images of the microstructure of the sample a small distance inside the sample being viewed. EDS was then used to map out the different elements present within the sample shown by each color contrast specific to a different element.

Vickers Hardness test was carried out on each of the samples to obtain a value for hardness on each of these samples. A small indentation is made on each of the samples with a load of 0.98 N with a holding time of 20 secs. The average Vickers hardness was reported after 5 measurements were done on each sample. Density measurements were carried out on all the sintered samples using the Archimedes principle. The measurements were carried out using ethanol as fluid. Three measurements were conducted, and the average density has been reported.

## Results and Discussion

### Morphological characteristics of the powder samples

The SEM metallography images show clear structural difference between the NiTi particles and WC particles as shown in Figure 5 and Figure 6, respectively. The NiTi powder showed particles were circular and round consistently throughout the sample with little to no impurities or defects to the particles. On the other hand, WC was observed to have a much smaller particle size (ranging from 1-5 µm) and was very random in shape which is a commonly supplied form for this material [12].
<figure><img src="image_243.png" /><figcaption>Figure 5: SEM image of self-mixed NiTi powder.</figcaption></figure>

<figure><img src="image_244.png" /><figcaption>Figure 6: SEM image of WC powder.</figcaption></figure>

<figure><img src="image_245.png" /><figcaption>Figure 7: A constant and uniform mixing of both powders were observed of mixed NiTi + 10% WC powder.</figcaption></figure>

As observed in Figure 7, all powders were found to be uniformly mixed with expected amounts of both NiTi and WC in all areas of the sample. EDS was conducted on all samples, and it was found that there was a steady increase in the WC in the average atomic mass percentage of the WC as more WC was added.

Backscattered scanning electron microscope images helped analyze the inner microstructure of each sample. The three distinct color contrasts were observed in the backscattered images of each of the samples as depicted by Figure 8.
<figure><img src="image_247.png" /><figcaption>Figure 5: SEM image of self-mixed NiTi powder.</figcaption></figure>

<figure><img src="image_248.png" /><figcaption>Figure 6: SEM image of WC powder.</figcaption></figure>

<figure><img src="image_249.png" /><figcaption>Figure 7: A constant and uniform mixing of both powders were observed of mixed NiTi + 10% WC powder.</figcaption></figure>

As observed in Figure 7, all powders were found to be uniformly mixed with expected amounts of both NiTi and WC in all areas of the sample. EDS was conducted on all samples, and it was found that there was a steady increase in the WC in the average atomic mass percentage of the WC as more WC was added.

Backscattered scanning electron microscope images helped analyze the inner microstructure of each sample. The three distinct color contrasts were observed in the backscattered images of each of the samples as depicted by Figure 8.
<figure><img src="image_246.png" /><figcaption>Figure 8: Backscattered scanning electron microscope image of NiTi + 10% WC.</figcaption></figure>

### Effect of addition of WC on the density of the samples

Due to the higher melting point of WC than that of NiTi, the WC particles would not form an alloy with NiTi after the sintering process. However, WC particles still retain all their properties post sintering which will increase mechanical properties such as hardness and density while being encapsulated in the ductile and biocompatible environment of the NiTi.

It was observed that there was a linear increase in the density of the samples with an increase in the WC percentage as shown in Figure 9. This is primarily because WC is a very dense material and as WC was not melted during the sintering process it still retains all its properties. Along with the reduction in the number of pores present throughout the sample, it also increased the density of the samples.

<figure><img src="image_247.png" /><figcaption>Figure 9: Density of the sintered samples after the compaction process with varying WC amounts.</figcaption></figure>

### Effect of addition of WC on the hardness values of the samples

A steady linear increase in the Vickers Hardness values was observed as the mass percentage of the WC powder present within the samples was increased as shown in Figure 10. From analysis of all samples, it was seen that the sintered samples containing only NiTi had a lot of pores which resulted in the reduction of the hardness value of the sample. These pores were formed because of either the parameters used for the sintering process or due to the evaporation of the PVA from inside the sample leaving void spaces. The number of pores visible in the samples gradually decreased as the amount of WC increased. This is because the areas in which the pores were formed
<figure><img src="image_250.png" /><figcaption>Figure 8: Backscattered scanning electron microscope image of NiTi + 10% WC.</figcaption></figure>

### Effect of addition of WC on the density of the samples

Due to the higher melting point of WC than that of NiTi, the WC particles would not form an alloy with NiTi after the sintering process. However, WC particles still retain all their properties post sintering which will increase mechanical properties such as hardness and density while being encapsulated in the ductile and biocompatible environment of the NiTi.

It was observed that there was a linear increase in the density of the samples with an increase in the WC percentage as shown in Figure 9. This is primarily because WC is a very dense material and as WC was not melted during the sintering process it still retains all its properties. Along with the reduction in the number of pores present throughout the sample, it also increased the density of the samples.

<figure><img src="image_251.png" /><figcaption>Figure 9: Density of the sintered samples after the compaction process with varying WC amounts.</figcaption></figure>

### Effect of addition of WC on the hardness values of the samples

A steady linear increase in the Vickers Hardness values was observed as the mass percentage of the WC powder present within the samples was increased as shown in Figure 10. From analysis of all samples, it was seen that the sintered samples containing only NiTi had a lot of pores which resulted in the reduction of the hardness value of the sample. These pores were formed because of either the parameters used for the sintering process or due to the evaporation of the PVA from inside the sample leaving void spaces. The number of pores visible in the samples gradually decreased as the amount of WC increased. This is because the areas in which the pores were formed
were filled with the WC powder which resulted in significantly fewer void spaces throughout the sample thereby increasing both the hardness and strength of the material.

<figure><img src="image_248.png" /><figcaption>Figure 10: Vickers Hardness values was observed for all samples with varying WC %.</figcaption></figure>

## Conclusion

In this work, the advantages of adding tungsten carbide (WC) in different amounts to nickel-rich nitinol were investigated. Due to its extraordinarily high melting point, it was found that the WC powder did not entirely melt in any of the samples. The evaporation of PVA left gaps in the sample, but this powder-filled them and provided structural support for the NiTi resulting in fully dense solid samples. The sample's hardness increased with increasing amounts of WC. These are all extremely crucial elements to consider when creating any stent and will lower risk factors like restenosis and stent distortion. Further work will be carried out to evaluate the influence of addition of WC to NiTi on radiopacity and biocompatibility of the produced samples. As NiTi is still used as the primary element for the fabrication of the stent, it still inhibits all important properties such as shape memory, corrosion resistance, and biocompatibility along with the mechanical advantages of WC, making the addition of WC to NiTi stents a potential advantage for the fabrication of stents.

## References:

[1] A. R. Pelton, D. Stöckel, and T. W. Duerig, "Medical Uses of Nitinol," *Mater. Sci. Forum*, vol. 327-328, pp. 63-70, Jan. 2000. https://doi.org/10.4028/www.scientific.net/MSF.327-328.63

[2] Y. Liu, J. Van Humbeeck, R. Stalmans, and L. Delaey, "Some aspects of the properties of NiTi shape memory alloy," *J. Alloys Compd.*, vol. 247, no. 1-2, pp. 115-121, 1997. https://doi.org/10.1016/S0925-8388(96)02572-8

[3] A. R. Pelton, S. M. Russell, and J. DiCello, "The physical metallurgy of nitinol for medical applications," *JOM*, vol. 55, no. 5, pp. 33-37, May 2003. https://doi.org/10.1007/s11837-003-0243-3

[4] J. Ryhänen, "Biocompatibility of Nitinol," *Minim. Invasive Ther. Allied Technol.*, vol. 9, no. 2, pp. 99-105, 2000. https://doi.org/10.3109/13645700009063056

[5] C. Constant, S. Nichols, É. Wagnac, Y. Petit, A. Desrochers, and V. Braïlovski, "Biocompatibility and mechanical stability of Nitinol as biomaterial for intra-articular prosthetic devices," *Materialia*, vol. 9, p. 100567, Mar. 2020. https://doi.org/10.1016/j.mtla.2019.100567
were filled with the WC powder which resulted in significantly fewer void spaces throughout the sample thereby increasing both the hardness and strength of the material.

<figure><img src="image_252.png" /><figcaption>Figure 10: Vickers Hardness values was observed for all samples with varying WC %.</figcaption></figure>

## Conclusion

In this work, the advantages of adding tungsten carbide (WC) in different amounts to nickel-rich nitinol were investigated. Due to its extraordinarily high melting point, it was found that the WC powder did not entirely melt in any of the samples. The evaporation of PVA left gaps in the sample, but this powder-filled them and provided structural support for the NiTi resulting in fully dense solid samples. The sample's hardness increased with increasing amounts of WC. These are all extremely crucial elements to consider when creating any stent and will lower risk factors like restenosis and stent distortion. Further work will be carried out to evaluate the influence of addition of WC to NiTi on radiopacity and biocompatibility of the produced samples. As NiTi is still used as the primary element for the fabrication of the stent, it still inhibits all important properties such as shape memory, corrosion resistance, and biocompatibility along with the mechanical advantages of WC, making the addition of WC to NiTi stents a potential advantage for the fabrication of stents.

## References:

[1] A. R. Pelton, D. Stöckel, and T. W. Duerig, “Medical Uses of Nitinol,” *Mater. Sci. Forum*, vol. 327–328, pp. 63–70, Jan. 2000. https://doi.org/10.4028/www.scientific.net/MSF.327-328.63

[2] Y. Liu, J. Van Humbeeck, R. Stalmans, and L. Delaey, “Some aspects of the properties of NiTi shape memory alloy,” *J. Alloys Compd.*, vol. 247, no. 1–2, pp. 115–121, 1997. https://doi.org/10.1016/S0925-8388(96)02572-8

[3] A. R. Pelton, S. M. Russell, and J. DiCello, “The physical metallurgy of nitinol for medical applications,” *JOM*, vol. 55, no. 5, pp. 33–37, May 2003. https://doi.org/10.1007/s11837-003-0243-3

[4] J. Ryhänen, “Biocompatibility of Nitinol,” *Minim. Invasive Ther. Allied Technol.*, vol. 9, no. 2, pp. 99–105, 2000. https://doi.org/10.3109/13645700009063056

[5] C. Constant, S. Nichols, É. Wagnac, Y. Petit, A. Desrochers, and V. Braïlovski, “Biocompatibility and mechanical stability of Nitinol as biomaterial for intra-articular prosthetic devices,” *Materialia*, vol. 9, p. 100567, Mar. 2020. https://doi.org/10.1016/j.mtla.2019.100567
[6] M. H. Elahinia, M. Hashemi, M. Tabesh, and S. B. Bhaduri, "Manufacturing and processing of NiTi implants: A review," *Prog. Mater. Sci.*, vol. 57, no. 5, pp. 911-946, Jun. 2012. https://doi.org/10.1016/j.pmatsci.2011.11.001

[7] M. Es-Souni, M. Es-Souni, and H. Fischer-Brandies, "Assessing the biocompatibility of NiTi shape memory alloys used for medical applications," *Anal. Bioanal. Chem.*, vol. 381, no. 3, pp. 557-567, Feb. 2005. https://doi.org/10.1007/s00216-004-2888-3

[8] A. Tuissi et al., "Radiopaque Shape Memory Alloys: NiTi-Er with Stable Superelasticity," *Shape Mem. Superelasticity*, vol. 2, no. 2, pp. 196-203, 2016. https://doi.org/10.1007/s40830-016-0066-z

[9] C. Park, S. Kim, H. E. Kim, and T. S. Jang, "Mechanically stable tantalum coating on a nano-roughened NiTi stent for enhanced radiopacity and biocompatibility," *Surf. Coatings Technol.*, vol. 305, pp. 139-145, 2016. https://doi.org/10.1016/j.surfcoat.2016.08.014

[10] D. M. Goldstein, W. J. Buehler, R. C. Wiley, "Effects of alloying upon certain properties of 55.1 NITINOL", NOLTR 64-235, August 1965

[11] A. Mussatto et al., "Laser-powder bed fusion in-process dispersion of reinforcing ceramic nanoparticles onto powder beds via colloid nebulisation," *Mater. Chem. Phys.*, vol. 287, no. February, p. 126245, 2022. https://doi.org/10.1016/j.matchemphys.2022.126245

[12] Z. Yang, J. Fan, Y. Liu, J. Nie, Z. Yang, and Y. Kang, "Effect of the particle size and matrix strength on strengthening and damage process of the particle reinforced metal matrix composites," *Materials (Basel)*, vol. 14, no. 3, pp. 1-12, 2021. https://doi.org/10.3390/ma14030675
[6] M. H. Elahinia, M. Hashemi, M. Tabesh, and S. B. Bhaduri, “Manufacturing and processing of NiTi implants: A review,” *Prog. Mater. Sci.*, vol. 57, no. 5, pp. 911–946, Jun. 2012. https://doi.org/10.1016/j.pmatsci.2011.11.001

[7] M. Es-Souni, M. Es-Souni, and H. Fischer-Brandies, “Assessing the biocompatibility of NiTi shape memory alloys used for medical applications,” *Anal. Bioanal. Chem.*, vol. 381, no. 3, pp. 557–567, Feb. 2005. https://doi.org/10.1007/s00216-004-2888-3

[8] A. Tuissi et al., “Radiopaque Shape Memory Alloys: NiTi–Er with Stable Superelasticity,” *Shape Mem. Superelasticity*, vol. 2, no. 2, pp. 196–203, 2016. https://doi.org/10.1007/s40830-016-0066-z

[9] C. Park, S. Kim, H. E. Kim, and T. S. Jang, “Mechanically stable tantalum coating on a nano-roughened NiTi stent for enhanced radiopacity and biocompatibility,” *Surf. Coatings Technol.*, vol. 305, pp. 139–145, 2016. https://doi.org/10.1016/j.surfcoat.2016.08.014

[10] D. M. Goldstein, W. J. Buehler, R. C. Wiley, “Effects of alloying upon certain properties of 55.1 NITINOL”, NOLTR 64-235, August 1965

[11] A. Mussatto et al., “Laser-powder bed fusion in-process dispersion of reinforcing ceramic nanoparticles onto powder beds via colloid nebulisation,” *Mater. Chem. Phys.*, vol. 287, no. February, p. 126245, 2022. https://doi.org/10.1016/j.matchemphys.2022.126245

[12] Z. Yang, J. Fan, Y. Liu, J. Nie, Z. Yang, and Y. Kang, “Effect of the particle size and matrix strength on strengthening and damage process of the particle reinforced metal matrix composites,” *Materials (Basel)*, vol. 14, no. 3, pp. 1–12, 2021. https://doi.org/10.3390/ma14030675
# Laser metal deposition of NbTaTiV refractory high entropy alloy

BARTH Eric¹, a * and HOR Anis¹, b

¹Espace Clément Ader, 3 Rue Caroline Aigle d'Aigle, 31400 Toulouse, France

aeric.barth@isae-supaero.fr, banis.hor@isae-supaero.fr

**Keywords:** Additive Manufacturing, Laser Metal Deposition, High Temperature, Microstructure, Complex Concentrated Alloys

**Abstract.** A mixture of Ti, V, Nb and Ta elemental powders was used to print an equimolar Refractory High Entropy Alloy NbTaTiV. Single beads and wall structures were printed on Ti-6Al-4V substrates. The samples were then cut to analyze their cross-section and composition. The samples present a large proportion of non-melted Ta particles, due to the low max power of the used laser (500W) and of the high melting-point of Ta. It was found that using a two-step deposition process (deposition, followed by a powder-less re-melting pass) could allow to homogenize the samples and melt said Ta particles. Moreover, the authors report difficulty to control the composition of the samples due to the very different physical properties of the different elements composing the alloy, with a preferential deposition of Ta.

**Introduction**

The concept “High Entropy Alloys (HEA)” was first theorized by Yeh et al. [1] and Cantor et al. [2] in 2004, as a novel way to design alloys by mixing elements in equimolar or close to equimolar proportions. Since then, a sub-category was introduced by Senkov et al. [3]: Refractory High Entropy Alloys (RHEA), made of refractory elements such as Ta, W and Mo. These alloys have immediately attracted high interest, due to their exceptional resistance at extreme temperatures (over 1 600°C) [4], [5]. However, several drawbacks hinder their potential applications. Most notably, as of now, the majority of RHEA is extremely brittle at room temperature [3], [6]. Moreover, their high melting points and complex composition makes it challenging to obtain homogeneous ingots through conventional casting processes [7]. Therefore, using a process with a small size melt-pool and a fast cooling process could help to obtain homogeneous samples. Hence the recent interest in building these materials through additive manufacturing processes, more specifically powder based deposition processes, such as Laser Metal Deposition (LMD). On top of inducing the small melt pools and fast cooling rates that can help to obtain homogeneous samples, this process can be used to build sample from a mixture of elemental powders, whose chemical composition can be modified *in-situ* to generate a composition gradient. This gradient can be used in functional parts, to deposit the highly expensive elements only where they are needed, but also to allow for a high-throughput analysis of all composition opened by the HEA conception paradigm shift. However, LMD requires a minimal level of ductility, as the fast cooling rates can generate residual stresses, which can result in high cracks density if the material is too brittle [6].

In 2018, Lee et al. proposed a new BCC RHEA, with both high resistance at elevated temperatures and good ductility at room temperatures [8], an equimolar mixture of Ti, V, Nb and Ta. This alloy system could be a promising candidate, as its ductility would allow it to be manufactured by additive processes. However, the large differences in the different elements' properties, such as melting temperature (1668°C for Ti and 3020°C for Ta) and density (4.51g/cm³ for Ti and 16.4g/cm³ for Ta) can lead to new challenges, even more so if one uses a mixture of pure elemental powder instead of pre-alloyed ones. These differences can lead to heterogeneities and to difficulties to achieve the desired sample's composition [6], [9].

<figure><img src="image_249.png" /></figure>
# Laser metal deposition of NbTaTiV refractory high entropy alloy

BARTH Eric¹, a * and HOR Anis¹, b

¹Espace Clément Ader, 3 Rue Caroline Aigle d'Aigle, 31400 Toulouse, France

aeric.barth@isae-supaero.fr, banis.hor@isae-supaero.fr

**Keywords:** Additive Manufacturing, Laser Metal Deposition, High Temperature, Microstructure, Complex Concentrated Alloys

**Abstract.** A mixture of Ti, V, Nb and Ta elemental powders was used to print an equimolar Refractory High Entropy Alloy NbTaTiV. Single beads and wall structures were printed on Ti-6Al-4V substrates. The samples were then cut to analyze their cross-section and composition. The samples present a large proportion of non-melted Ta particles, due to the low max power of the used laser (500W) and of the high melting-point of Ta. It was found that using a two-step deposition process (deposition, followed by a powder-less re-melting pass) could allow to homogenize the samples and melt said Ta particles. Moreover, the authors report difficulty to control the composition of the samples due to the very different physical properties of the different elements composing the alloy, with a preferential deposition of Ta.

**Introduction**

The concept “High Entropy Alloys (HEA)” was first theorized by Yeh et al. [1] and Cantor et al. [2] in 2004, as a novel way to design alloys by mixing elements in equimolar or close to equimolar proportions. Since then, a sub-category was introduced by Senkov et al. [3]: Refractory High Entropy Alloys (RHEA), made of refractory elements such as Ta, W and Mo. These alloys have immediately attracted high interest, due to their exceptional resistance at extreme temperatures (over 1 600°C) [4], [5]. However, several drawbacks hinder their potential applications. Most notably, as of now, the majority of RHEA is extremely brittle at room temperature [3], [6]. Moreover, their high melting points and complex composition makes it challenging to obtain homogeneous ingots through conventional casting processes [7]. Therefore, using a process with a small size melt-pool and a fast cooling process could help to obtain homogeneous samples. Hence the recent interest in building these materials through additive manufacturing processes, more specifically powder based deposition processes, such as Laser Metal Deposition (LMD). On top of inducing the small melt pools and fast cooling rates that can help to obtain homogeneous samples, this process can be used to build sample from a mixture of elemental powders, whose chemical composition can be modified *in-situ* to generate a composition gradient. This gradient can be used in functional parts, to deposit the highly expensive elements only where they are needed, but also to allow for a high-throughput analysis of all composition opened by the HEA conception paradigm shift. However, LMD requires a minimal level of ductility, as the fast cooling rates can generate residual stresses, which can result in high cracks density if the material is too brittle [6].

In 2018, Lee et al. proposed a new BCC RHEA, with both high resistance at elevated temperatures and good ductility at room temperatures [8], an equimolar mixture of Ti, V, Nb and Ta. This alloy system could be a promising candidate, as its ductility would allow it to be manufactured by additive processes. However, the large differences in the different elements' properties, such as melting temperature (1 668°C for Ti and 3 020°C for Ta) and density (4.51g/cm³ for Ti and 16.4g/cm³ for Ta) can lead to new challenges, even more so if one uses a mixture of pure elemental powder instead of pre-alloyed ones. These differences can lead to heterogeneities and to difficulties to achieve the desired sample's composition [6], [9].

<figure><img src="image_253.png" /></figure>
This paper, therefore, offers to study the manufacturability of the NbTaTiV RHEA by LMD.

## Material and method

Samples were manufactured using a mixture of high purity (99,9 %at) Ti, V, Nb and Ta elemental powders, manufactured by Metal Powder Tech International, with a 45-90 µm particle size. All powders were found to have good sphericity, except Ti powders, within which a number of particles had highly irregular shapes. The powders dimensions were within the desired margins, except for Ta which was found to have a large fraction of smaller particles.

In order to consider heterogeneous deposition efficiencies of the different elements, a non-equimolar Ti₂₅V₂₀Nb₂₀Ta₃₅ powder chemical composition was selected on the basis of the literature review [10], [11]. The mixing process was performed manually in a controlled environment.

The powder mixture was then poured in a powder hopper connected to a commercial BeAM Mobile, an LPBF additive manufacturing machine, equipped with a 5-axis deposition system, a 500W Nd:YAG laser and a sealed chamber, used in this study to manufacture the samples under Ar atmosphere.

20mm-long single tracks and 4-layers wall samples were then printed in an inert atmosphere (<1000ppm O₂) on a Ti-6Al-4V substrate using variable laser powers (P), printhead speed (V) and powder mass-flow (ṁ). The upper and lower limits of these parameters are reported in Table 1.

<table><caption>Table 1: Upper and lower limits of the manufacturing parameters</caption><thead><tr><td>Laser Power [W]</td><td>Printhead speed [mm/min]</td><td>Powder mass-flow [g/min]</td></tr></thead><tbody><tr><td>400 - 500</td><td>400 - 1200</td><td>3.5 - 7.3</td></tr></tbody></table>

As suggested by Dobbelstein et al. [6], some samples were manufactured using a 2-step deposition process. First, a conventional deposition step, followed by a second re-melting pass, without adding more powder (Fig. 1). This process is meant to ensure the melting of the high melting-temperature elements.

<figure><img src="image_250.png" /></figure>

<figure><img src="image_251.png" /><figcaption>Fig. 1: Illustration of the two deposition strategies used (a): 1-step strategy and (b): 2-step strategy, with a re-melting step</figcaption></figure>

Each single-track sample was made 3 consecutive times to assess the repeatability of the process. In order to analyze the composition of the powder exiting the printhead, powder was deposited without turning on the laser on steel substrate covered with glue, which will thereafter be referred to as “Sticky Bands”. This was made to retrieve a single layer of powder particle and avoid errors in composition analysis linked with density-based stratification of the various elements.

The printed samples were then cut and polished using SiC abrasive papers and diamond suspensions. The cross-sections were observed using a Zeiss EVO-15 scanning electron microscope (SEM) equipped with an Energy Dispersive X-Ray (EDX) for composition analysis.
This paper, therefore, offers to study the manufacturability of the NbTaTiV RHEA by LMD.

## Material and method

Samples were manufactured using a mixture of high purity (99,9 %at) Ti, V, Nb and Ta elemental powders, manufactured by Metal Powder Tech International, with a 45-90 µm particle size. All powders were found to have good sphericity, except Ti powders, within which a number of particles had highly irregular shapes. The powders dimensions were within the desired margins, except for Ta which was found to have a large fraction of smaller particles.

In order to consider heterogeneous deposition efficiencies of the different elements, a non-equimolar Ti₂₅V₂₀Nb₂₀Ta₃₅ powder chemical composition was selected on the basis of the literature review [10], [11]. The mixing process was performed manually in a controlled environment.

The powder mixture was then poured in a powder hopper connected to a commercial BeAM Mobile, an LPBF additive manufacturing machine, equipped with a 5-axis deposition system, a 500W Nd:YAG laser and a sealed chamber, used in this study to manufacture the samples under Ar atmosphere.

20mm-long single tracks and 4-layers wall samples were then printed in an inert atmosphere (<1000ppm O₂) on a Ti-6Al-4V substrate using variable laser powers (P), printhead speed (V) and powder mass-flow (ṁ). The upper and lower limits of these parameters are reported in Table 1.

<table><caption>Table 1: Upper and lower limits of the manufacturing parameters</caption><thead><tr><td>Laser Power [W]</td><td>Printhead speed [mm/min]</td><td>Powder mass-flow [g/min]</td></tr></thead><tbody><tr><td>400 - 500</td><td>400 - 1200</td><td>3.5 - 7.3</td></tr></tbody></table>

As suggested by Dobbelstein et al. [6], some samples were manufactured using a 2-step deposition process. First, a conventional deposition step, followed by a second re-melting pass, without adding more powder (Fig. 1). This process is meant to ensure the melting of the high melting-temperature elements.

<figure><img src="image_254.png" /></figure>

<figure><img src="image_255.png" /><figcaption>Fig. 1: Illustration of the two deposition strategies used (a): 1-step strategy and (b): 2-step strategy, with a re-melting step</figcaption></figure>

Each single-track sample was made 3 consecutive times to assess the repeatability of the process. In order to analyze the composition of the powder exiting the printhead, powder was deposited without turning on the laser on steel substrate covered with glue, which will thereafter be referred to as “Sticky Bands”. This was made to retrieve a single layer of powder particle and avoid errors in composition analysis linked with density-based stratification of the various elements.

The printed samples were then cut and polished using SiC abrasive papers and diamond suspensions. The cross-sections were observed using a Zeiss EVO-15 scanning electron microscope (SEM) equipped with an Energy Dispersive X-Ray (EDX) for composition analysis.
## Results

The results of the EDX analysis of the powders retrieved on three of the sticky bands is visible on Fig. 2.

The three bands, referred as S1, S2 and S3, where exposed to the powder flow under the same conditions at various moments of the sample manufacturing campaign. For each band, a minimum of 5 measurements was conducted. It appears that the composition of the retrieved powders strongly deviates from the expected one, and changes from one sticky band to the other. The mean deviation for one band remains low (>5%atm, except for the Ti on S1).

Regarding the SEM observations of the printed samples, the samples can be divided in two categories depending on whether or not they were subjected to a second re-melting laser-pass. A typical sectional view of a single-pass sample is visible on Fig. 3(a).

<figure><img src="image_252.png" /><figcaption>Fig. 2: Average composition of the powder retrieved on the sticky bands and associated mean deviation in the form of error bars, with corresponding examples of EDX mapping superimposed on SEM imageries.</figcaption></figure>

<figure><img src="image_253.png" /><figcaption>Fig. 3: SEM imaging of typical sectional views of (a) a single-pass sample printed with a P=500W, V=600mm/min and $\dot{m}$=7.3g/min and (b) a double-pass sample printed with the same parameters for the deposition steps and P=500W and V=400mm/min for the re-melting pass.</figcaption></figure>

The first thing visible is the large amount of non-melted Ta particles (white particles, highlighted by blue arrows), and of Nb particles (dark particles, highlighted by purple arrows) in lesser proportions, whose presence can be explained by their extremely high melting point ($T_{Ta}$=3020°C and $T_{Nb}$=2477°C). Moreover, a large fraction of the single-pass samples exhibits a Ti-rich outer edge (darker edge on Fig. 3(a)). This partitioning of the elements is proof that the
## Results

The results of the EDX analysis of the powders retrieved on three of the sticky bands is visible on Fig. 2.

The three bands, referred as S1, S2 and S3, where exposed to the powder flow under the same conditions at various moments of the sample manufacturing campaign. For each band, a minimum of 5 measurements was conducted. It appears that the composition of the retrieved powders strongly deviates from the expected one, and changes from one sticky band to the other. The mean deviation for one band remains low (>5%atm, except for the Ti on S1).

Regarding the SEM observations of the printed samples, the samples can be divided in two categories depending on whether or not they were subjected to a second re-melting laser-pass. A typical sectional view of a single-pass sample is visible on Fig. 3(a).

<figure><img src="image_256.png" /><figcaption>Fig. 2: Average composition of the powder retrieved on the sticky bands and associated mean deviation in the form of error bars, with corresponding examples of EDX mapping superimposed on SEM imageries.</figcaption></figure>

<figure><img src="image_257.png" /><figcaption>Fig. 3: SEM imaging of typical sectional views of (a) a single-pass sample printed with a P=500W, V=600mm/min and $\dot{m}$=7.3g/min and (b) a double-pass sample printed with the same parameters for the deposition steps and P=500W and V=400mm/min for the re-melting pass.</figcaption></figure>

The first thing visible is the large amount of non-melted Ta particles (white particles, highlighted by blue arrows), and of Nb particles (dark particles, highlighted by purple arrows) in lesser proportions, whose presence can be explained by their extremely high melting point ($T_{Ta}$=3020°C and $T_{Nb}$=2477°C). Moreover, a large fraction of the single-pass samples exhibits a Ti-rich outer edge (darker edge on Fig. 3(a)). This partitioning of the elements is proof that the
inner part of the sample cooled first, leading to a partial segregation of the low-melting temperature Ti in said edges.

A typical sectional view of a double-pass sample is visible on Fig. 3(b). This sample also exhibits a large concentration of non-melted Ta particles. However, the upper fraction of the sample's cross section appears to be homogeneous. These homogeneous areas have a depth between 200 and 250µm.

Typical high-magnification EDX composition maps can be seen on Fig. 4. For all samples, a dendritic microstructure can be observed. The dendrites are Ta-rich, whereas the inter-dendrites are mostly Ti, V and Nb-rich. This is coherent with the melting points of the various elements, and confirms that, even though pure Ta particle can be seen, a fraction of it was indeed melted.

<figure><img src="image_254.png" /><figcaption>Fig. 4: x1000 Magnification EDX mapping of the various element's concentrations. These maps where acquired on a single-pass sample manufactured with P=500W and V=700mm/min</figcaption></figure>

Nb concentration follows a peculiar behaviour compared to the other elements. Even though it appears to be concentrated in the inter-dendrites, the segregation does not appear as severe as it is for V and Ti. This hints towards a diffusion of Nb in the Ta-rich dendrites during the solidification process.

The composition of the single-pass samples was analysed, and the average concentrations of each elements can be seen in Table 2.

<table><caption>Table 2: Average concentration of each elements in the single-track samples and corresponding standard deviation</caption><thead><tr><th></th><th>Ti</th><th>V</th><th>Nb</th><th>Ta</th></tr></thead><tbody><tr><td>Average concentration (at.%)</td><td>28.0</td><td>11.9</td><td>14.9</td><td>45.2</td></tr><tr><td>Standard Deviation (at.%)</td><td>11.3</td><td>1.7</td><td>4.3</td><td>8.8</td></tr></tbody></table>

It appears that the composition of the samples strongly deviates from the expected equimolar one. V and Nb are underrepresented, whereas Ta is over-represented. The average Ti concentration is close to the expected one, but is highly variable from one sample to the next. This is most likely due to random levels of chemical contaminations from the substrates. This can be confirmed by analysing the concentration evolution of the various elements on the different layers of the wall-type samples.
inner part of the sample cooled first, leading to a partial segregation of the low-melting temperature Ti in said edges.

A typical sectional view of a double-pass sample is visible on Fig. 3(b). This sample also exhibits a large concentration of non-melted Ta particles. However, the upper fraction of the sample's cross section appears to be homogeneous. These homogeneous areas have a depth between 200 and 250µm.

Typical high-magnification EDX composition maps can be seen on Fig. 4. For all samples, a dendritic microstructure can be observed. The dendrites are Ta-rich, whereas the inter-dendrites are mostly Ti, V and Nb-rich. This is coherent with the melting points of the various elements, and confirms that, even though pure Ta particle can be seen, a fraction of it was indeed melted.

<figure><img src="image_258.png" /><figcaption>Fig. 4: x1000 Magnification EDX mapping of the various element's concentrations. These maps where acquired on a single-pass sample manufactured with P=500W and V=700mm/min</figcaption></figure>

Nb concentration follows a peculiar behaviour compared to the other elements. Even though it appears to be concentrated in the inter-dendrites, the segregation does not appear as severe as it is for V and Ti. This hints towards a diffusion of Nb in the Ta-rich dendrites during the solidification process.

The composition of the single-pass samples was analysed, and the average concentrations of each elements can be seen in Table 2.

<table><caption>Table 2: Average concentration of each elements in the single-track samples and corresponding standard deviation</caption><thead><tr><th></th><th>Ti</th><th>V</th><th>Nb</th><th>Ta</th></tr></thead><tbody><tr><td>Average concentration (at.%)</td><td>28.0</td><td>11.9</td><td>14.9</td><td>45.2</td></tr><tr><td>Standard Deviation (at.%)</td><td>11.3</td><td>1.7</td><td>4.3</td><td>8.8</td></tr></tbody></table>

It appears that the composition of the samples strongly deviates from the expected equimolar one. V and Nb are underrepresented, whereas Ta is over-represented. The average Ti concentration is close to the expected one, but is highly variable from one sample to the next. This is most likely due to random levels of chemical contaminations from the substrates. This can be confirmed by analysing the concentration evolution of the various elements on the different layers of the wall-type samples.
<figure><img src="image_255.png" /><figcaption>Fig. 5: Sectional view of a wall-type sample printed with the dual-pass method (P=500W, V=400mm/min and $\dot{m}$=3.5g/min for the deposition step, and P=500W and V=400mm/min for the re-melting step), and average concentration evolution of the various elements for the different layers for all wall type samples.</figcaption></figure>

As it can be seen on Fig. 5, the average Ti concentration drops significantly between the first and second layer. The composition then remains stable along the sample's height. Moreover, as shown on Fig. 5, Ta particles can also be found in the wall structures. For the dual-pass samples, the top part of each layer is homogeneous, leading to high concentration of Ta particles at the interface between layers. Moreover, these wall structures are free of cracks and porosities.

The dimensions of the single-track samples were measured through image analysis. Fig. 6 shows the evolution of the samples' heights as a function of the deposition speed, regardless of laser power.

<figure><img src="image_256.png" /><figcaption>Fig. 6: Sample height of the single-track samples as a function of deposition speed for various powder mass flow rate, regardless of the laser power</figcaption></figure>

It appears that a direct correlation exists between a sample's height and the deposition speed and powder mass flow with which it was deposited.
<figure><img src="image_259.png" /><figcaption>Fig. 5: Sectional view of a wall-type sample printed with the dual-pass method (P=500W, V=400mm/min and $\dot{m}$=3.5g/min for the deposition step, and P=500W and V=400mm/min for the re-melting step), and average concentration evolution of the various elements for the different layers for all wall type samples.</figcaption></figure>

As it can be seen on Fig. 5, the average Ti concentration drops significantly between the first and second layer. The composition then remains stable along the sample's height. Moreover, as shown on Fig. 5, Ta particles can also be found in the wall structures. For the dual-pass samples, the top part of each layer is homogeneous, leading to high concentration of Ta particles at the interface between layers. Moreover, these wall structures are free of cracks and porosities.

The dimensions of the single-track samples were measured through image analysis. Fig. 6 shows the evolution of the samples' heights as a function of the deposition speed, regardless of laser power.

<figure><img src="image_260.png" /><figcaption>Fig. 6: Sample height of the single-track samples as a function of deposition speed for various powder mass flow rate, regardless of the laser power</figcaption></figure>

It appears that a direct correlation exists between a sample's height and the deposition speed and powder mass flow with which it was deposited.
## Discussion

The current main problematic is the high concentration of Ta particles in the samples, due to their high melting temperature. As the machine available for this study is only equipped with a 500W laser, the only way to increase the energy input is to reduce the deposition speed. However, doing so with a single-pass strategy means that more powder will be deposited, negating the benefice of the increased energy output.

Upon analysing the depth of the area homogenized by the re-melting step, one can notice that said depth remains fairly consistent from one sample to the next.

This, combined with the sample's height trend identified previously (Fig. 6), tends to show that the optimal method to achieve a complete melting of the matter is to use a high deposition speed and/or a low powder mass flowrate to obtain slim seams. This, combined with a re-melting pass, could ensure that a large fraction of the seam is homogenised, as illustrated on Fig. 7.

<figure><img src="image_257.png" /></figure>

<table><caption>Fig. 7: Cross section of 3 single track double-pass samples manufactured with the same re-melting parameters (P=500W and V=400mm/min) and different deposition parameters. The homogenised part of the cross-section is highlighted in yellow.</caption><thead><tr><th colspan="3">Deposition Parameters</th></tr></thead><tbody><tr><td>P=500W V=400mm/min<br>ṁ=7.3g/min</td><td>P=500W V=400mm/min<br>ṁ=3.5g/min</td><td>P=500W V=800mm/min<br>ṁ=3.5g/min</td></tr></tbody></table>

Another current issue lies in the control of the composition of the powder feed. The analysis of the powder retrieved on the sticky bands shows that powder composition does not evolve much when measured on a single band, meaning that instantaneous variation in the composition are low. However, the composition differences between different bands are very large. This goes to show that the powder mixture in the printer's hopper is imperfect. Multiple root cause can be identified. First, the strong density difference between the various powders can lead to a de-mixing of the powders over time. Moreover, the manual mixing might not be enough to fully homogenise the mixture, even though only a small quantity of powders was mixed at a given time. Better mixing procedures, compatible with the high oxidation rate of pure elemental powders, needs to be put in place.

Even if the composition of the powder feed is variable, a clear tendency appears in the composition of the samples. The chemical analysis of the wall-type samples proves that the Ta has a much higher deposition efficiency compared to the other elements (Fig. 5). This goes against what can be seen in the literature regarding the manufacturing of RHEA through LMD, where the elements with the highest melting point are usually the ones with the lowest deposition efficiency due to them bouncing off the melt pool [9], [11], [12]. This phenomenon can be explained by the large fraction of Ta particles of smaller diameter, the later having a higher chance of sticking to the melt pool.

## Conclusions

From the present study, the following conclusions can be drawn:

1/ Manufacturing the NbTaTiV alloy by LMD with a 500W laser leads to the formation of a beads containing a large concentration of non-melted Ta and Nb powders, due to their high melting temperature.
## Discussion

The current main problematic is the high concentration of Ta particles in the samples, due to their high melting temperature. As the machine available for this study is only equipped with a 500W laser, the only way to increase the energy input is to reduce the deposition speed. However, doing so with a single-pass strategy means that more powder will be deposited, negating the benefice of the increased energy output.

Upon analysing the depth of the area homogenized by the re-melting step, one can notice that said depth remains fairly consistent from one sample to the next.

This, combined with the sample's height trend identified previously (Fig. 6), tends to show that the optimal method to achieve a complete melting of the matter is to use a high deposition speed and/or a low powder mass flowrate to obtain slim seams. This, combined with a re-melting pass, could ensure that a large fraction of the seam is homogenised, as illustrated on Fig. 7.

<figure><img src="image_261.png" /></figure>

<table><caption>Fig. 7: Cross section of 3 single track double-pass samples manufactured with the same re-melting parameters (P=500W and V=400mm/min) and different deposition parameters. The homogenised part of the cross-section is highlighted in yellow.</caption><thead><tr><th colspan="3">Deposition Parameters</th></tr></thead><tbody><tr><td>P=500W V=400mm/min<br>ṁ=7.3g/min</td><td>P=500W V=400mm/min<br>ṁ=3.5g/min</td><td>P=500W V=800mm/min<br>ṁ=3.5g/min</td></tr></tbody></table>

Another current issue lies in the control of the composition of the powder feed. The analysis of the powder retrieved on the sticky bands shows that powder composition does not evolve much when measured on a single band, meaning that instantaneous variation in the composition are low. However, the composition differences between different bands are very large. This goes to show that the powder mixture in the printer's hopper is imperfect. Multiple root cause can be identified. First, the strong density difference between the various powders can lead to a de-mixing of the powders over time. Moreover, the manual mixing might not be enough to fully homogenise the mixture, even though only a small quantity of powders was mixed at a given time. Better mixing procedures, compatible with the high oxidation rate of pure elemental powders, needs to be put in place.

Even if the composition of the powder feed is variable, a clear tendency appears in the composition of the samples. The chemical analysis of the wall-type samples proves that the Ta has a much higher deposition efficiency compared to the other elements (Fig. 5). This goes against what can be seen in the literature regarding the manufacturing of RHEA through LMD, where the elements with the highest melting point are usually the ones with the lowest deposition efficiency due to them bouncing off the melt pool [9], [11], [12]. This phenomenon can be explained by the large fraction of Ta particles of smaller diameter, the later having a higher chance of sticking to the melt pool.

## Conclusions

From the present study, the following conclusions can be drawn:

1/ Manufacturing the NbTaTiV alloy by LMD with a 500W laser leads to the formation of a beads containing a large concentration of non-melted Ta and Nb powders, due to their high melting temperature.
2/Using a dual-pass deposition strategy (deposition followed by re-melting) creates a homogenized zone in the upper part of the bead's cross-section. Reducing the bead's height by increasing the deposition speed and reducing the powder mass flowrate can allow to fully homogenize the cross-section.

3/ The samples all present a dendritic microstructure with Ta-rich dendrites. The samples all present a high Ta concentration, due to a higher deposition efficiency of Ta particles.

4/ The powder composition appears to be variable over time. This variation can be attributed to an imperfect mixing of the powder in the hoppers.

References

[1] J.-W. Yeh et al., "Nanostructured High-Entropy Alloys with Multiple Principal Elements: Novel Alloy Design Concepts and Outcomes," *Advanced Engineering Materials*, vol. 6, no. 5, pp. 299-303, 2004. https://doi.org/10.1002/adem.200300567

[2] B. Cantor, I. T. H. Chang, P. Knight, and A. J. B. Vincent, "Microstructural development in equiatomic multicomponent alloys," *Materials Science and Engineering: A*, vol. 375-377, pp. 213-218, Jul. 2004. https://doi.org/10.1016/j.msea.2003.10.257

[3] O. N. Senkov, G. B. Wilks, D. B. Miracle, C. P. Chuang, and P. K. Liaw, "Refractory high-entropy alloys," *Intermetallics*, vol. 18, no. 9, pp. 1758-1765, Sep. 2010. https://doi.org/10.1016/j.intermet.2010.05.014

[4] O. N. Senkov, G. B. Wilks, J. M. Scott, and D. B. Miracle, "Mechanical properties of Nb25Mo25Ta25W25 and V20Nb20Mo20Ta20W20 refractory high entropy alloys," *Intermetallics*, vol. 19, no. 5, pp. 698-706, May 2011. https://doi.org/10.1016/j.intermet.2011.01.004

[5] R. Feng et al., "Superior High-Temperature Strength in a Supersaturated Refractory High-Entropy Alloy," *Advanced Materials*, vol. 33, no. 48, p. 2102401, 2021. https://doi.org/10.1002/adma.202102401

[6] H. Dobbelstein, M. Thiele, E. L. Gurevich, E. P. George, and A. Ostendorf, "Direct Metal Deposition of Refractory High Entropy Alloy MoNbTaW," *Physics Procedia*, vol. 83, pp. 624-633, Jan. 2016. https://doi.org/10.1016/j.phpro.2016.08.065

[7] J. P. Couzinié et al., "Microstructure of a near-equimolar refractory high-entropy alloy," *Materials Letters*, vol. 126, pp. 285-287, Jul. 2014. https://doi.org/10.1016/j.matlet.2014.04.062

[8] C. Lee et al., "Lattice distortion in a strong and ductile refractory high-entropy alloy," *Acta Materialia*, vol. 160, pp. 158-172, Nov. 2018. https://doi.org/10.1016/j.actamat.2018.08.053

[9] M. A. Melia et al., "High-throughput additive manufacturing and characterization of refractory high entropy alloys," *Applied Materials Today*, vol. 19, p. 100560, Jun. 2020. https://doi.org/10.1016/j.apmt.2020.100560

[10] H. Dobbelstein, E. L. Gurevich, E. P. George, A. Ostendorf, and G. Laplanche, "Laser metal deposition of compositionally graded TiZrNbTa refractory high-entropy alloys using elemental powder blends," *Additive Manufacturing*, vol. 25, pp. 252-262, Jan. 2019. https://doi.org/10.1016/j.addma.2018.10.042

[11] H. Dobbelstein, E. L. Gurevich, E. P. George, A. Ostendorf, and G. Laplanche, "Laser metal deposition of a refractory TiZrNbHfTa high-entropy alloy," *Additive Manufacturing*, vol. 24, pp. 386-390, Dec. 2018. https://doi.org/10.1016/j.addma.2018.10.008

[12] H. Dobbelstein, E. P. George, E. L. Gurevich, A. Kostka, A. Ostendorf, and G. Laplanche, "Laser metal deposition of refractory high-entropy alloys for high-throughput synthesis and structure-property characterization," *Int. J. Extrem. Manuf.*, vol. 3, no. 1, p. 015201, Dec. 2020. https://doi.org/10.1088/2631-7990/abcca8
2/Using a dual-pass deposition strategy (deposition followed by re-melting) creates a homogenized zone in the upper part of the bead's cross-section. Reducing the bead's height by increasing the deposition speed and reducing the powder mass flowrate can allow to fully homogenize the cross-section.

3/ The samples all present a dendritic microstructure with Ta-rich dendrites. The samples all present a high Ta concentration, due to a higher deposition efficiency of Ta particles.

4/ The powder composition appears to be variable over time. This variation can be attributed to an imperfect mixing of the powder in the hoppers.

## References

[1] J.-W. Yeh et al., "Nanostructured High-Entropy Alloys with Multiple Principal Elements: Novel Alloy Design Concepts and Outcomes," *Advanced Engineering Materials*, vol. 6, no. 5, pp. 299-303, 2004. https://doi.org/10.1002/adem.200300567

[2] B. Cantor, I. T. H. Chang, P. Knight, and A. J. B. Vincent, "Microstructural development in equiatomic multicomponent alloys," *Materials Science and Engineering: A*, vol. 375-377, pp. 213-218, Jul. 2004. https://doi.org/10.1016/j.msea.2003.10.257

[3] O. N. Senkov, G. B. Wilks, D. B. Miracle, C. P. Chuang, and P. K. Liaw, "Refractory high-entropy alloys," *Intermetallics*, vol. 18, no. 9, pp. 1758-1765, Sep. 2010. https://doi.org/10.1016/j.intermet.2010.05.014

[4] O. N. Senkov, G. B. Wilks, J. M. Scott, and D. B. Miracle, "Mechanical properties of Nb25Mo25Ta25W25 and V20Nb20Mo20Ta20W20 refractory high entropy alloys," *Intermetallics*, vol. 19, no. 5, pp. 698-706, May 2011. https://doi.org/10.1016/j.intermet.2011.01.004

[5] R. Feng et al., "Superior High-Temperature Strength in a Supersaturated Refractory High-Entropy Alloy," *Advanced Materials*, vol. 33, no. 48, p. 2102401, 2021. https://doi.org/10.1002/adma.202102401

[6] H. Dobbelstein, M. Thiele, E. L. Gurevich, E. P. George, and A. Ostendorf, "Direct Metal Deposition of Refractory High Entropy Alloy MoNbTaW," *Physics Procedia*, vol. 83, pp. 624-633, Jan. 2016. https://doi.org/10.1016/j.phpro.2016.08.065

[7] J. P. Couzinié et al., "Microstructure of a near-equimolar refractory high-entropy alloy," *Materials Letters*, vol. 126, pp. 285-287, Jul. 2014. https://doi.org/10.1016/j.matlet.2014.04.062

[8] C. Lee et al., "Lattice distortion in a strong and ductile refractory high-entropy alloy," *Acta Materialia*, vol. 160, pp. 158-172, Nov. 2018. https://doi.org/10.1016/j.actamat.2018.08.053

[9] M. A. Melia et al., "High-throughput additive manufacturing and characterization of refractory high entropy alloys," *Applied Materials Today*, vol. 19, p. 100560, Jun. 2020. https://doi.org/10.1016/j.apmt.2020.100560

[10] H. Dobbelstein, E. L. Gurevich, E. P. George, A. Ostendorf, and G. Laplanche, "Laser metal deposition of compositionally graded TiZrNbTa refractory high-entropy alloys using elemental powder blends," *Additive Manufacturing*, vol. 25, pp. 252-262, Jan. 2019. https://doi.org/10.1016/j.addma.2018.10.042

[11] H. Dobbelstein, E. L. Gurevich, E. P. George, A. Ostendorf, and G. Laplanche, "Laser metal deposition of a refractory TiZrNbHfTa high-entropy alloy," *Additive Manufacturing*, vol. 24, pp. 386-390, Dec. 2018. https://doi.org/10.1016/j.addma.2018.10.008

[12] H. Dobbelstein, E. P. George, E. L. Gurevich, A. Kostka, A. Ostendorf, and G. Laplanche, "Laser metal deposition of refractory high-entropy alloys for high-throughput synthesis and structure-property characterization," *Int. J. Extrem. Manuf.*, vol. 3, no. 1, p. 015201, Dec. 2020. https://doi.org/10.1088/2631-7990/abcca8
# Finite element modeling (FEM) as a design tool to produce thin wall structures in laser powder bed fusion (L-PBF)

POLLARA Gaetano¹,ª*, PALMERI Dina¹,ᵇ, BUFFA Gianluca¹,ᶜ and
FRATINI Livan¹,ᵈ

¹Dipartimento di Ingegneria, Università Degli Studi di Palermo, Viale delle Scienze, Palermo,
90128, Italy

ªgaetano.pollara@unipa.it, ᵇdina.palmeri@unipa.it, ᵈgianluca.buffa@unipa.it,
lioivan.fratini@unipa.it

**Keywords:** Numerical Simulation, Laser Powder Bed Fusion, Ti-6Al-4V

**Abstract.** Laser Powder Bed Fusion (LPBF) has been widely adopted in many industrial sectors such as biomedical, automotive, and aerospace thanks to the possibility to produce objects with complex shapes and meet customers' needs. Despite all the advantages that LPBF can offer, the rise of residual stress due to the high thermal gradients generated during the process can limit its application. This is the case with thin-wall structures where the build-up of residual stress can compromise the success of the printing process. Being able to print this structure can be useful in fuel cell applications where the implementation of cooling channels in bipolar plates can improve their performance. This paper aims to provide guidelines for designing thin-wall structures produced by LPBF processes through numerical simulations by understanding the effect of residual stress on part distortion.

**Introduction**

Nowadays, where the energy demand is constantly increasing and fossil-based fuels are the principal contributors to meeting the global energy need, hydrogen-based technologies can offer a good alternative to non-renewable energy sources that are doomed to be exhausted one day. Moreover, thanks to its inherently immense energy content, availability, and zero-emission, hydrogen is considered the future fuel in the transport sector. Particularly, hydrogen can be used in electrical power generation systems such as fuel cells for vehicles [1]. These kinds of fuel cells which use chemical hydrogen bonds to store electricity are known as proton exchange membrane (PEM) fuel cells. Bipolar plates (BSs) are the main component of PEM fuel cells in terms of structural strength, weight, and cost [2]. Usually, BPs are made of graphite because of its high conductivity, high corrosion resistance, low density, and ease of fabrication. However, manufacturing of graphite plates under a certain thickness can be difficult due to their brittle nature. Furthermore, it is very challenging to create flow channels in order to apply liquid cooling and improve the PEM fuel cells' performance [3]. As a consequence, metallic plates are taking increasing interest due to their high thermal and electrical conductivity, good gas impermeability, high strength, and good formability. In addition, by using metals as the material to produce metallic plates, additive manufacturing (AM) methods can be adopted [4]. AM consists of a layer-by-layer production fashion in which the part is obtained by adding layers of material. Thus, allowing the fabrication of parts with a very complex shape or structure. In the case of metallic parts, laser powder bed fusion (L-PBF) is among the more suitable AM technologies. In L-PBF a powder bed is deposited on a build plate within a chamber filled with an inert gas and the metallic powder is selectively melted with a laser heat source according to the slicing file [5].

Titanium alloys can be used to produce metallic plates thanks to their high mechanical strength, high corrosion resistance, and lightweight [6,7]. In this study, Ti-6Al-4V was used in the L-PBF process to print bipolar plates. These BPs are characterized by thin-wall structures which can be

<figure><img src="image_258.png" /></figure>
# Finite element modeling (FEM) as a design tool to produce thin wall structures in laser powder bed fusion (L-PBF)

POLLARA Gaetano¹,ᵃ*, PALMERI Dina¹,ᵇ, BUFFA Gianluca¹,ᶜ and
FRATINI Livan¹,ᵈ

¹Dipartimento di Ingegneria, Università Degli Studi di Palermo, Viale delle Scienze, Palermo,
90128, Italy

ªgaetano.pollara@unipa.it, ᵇdina.palmeri@unipa.it, ᶜgianluca.buffa@unipa.it,
[livan.fratini@unipa.it]

**Keywords:** Numerical Simulation, Laser Powder Bed Fusion, Ti-6Al-4V

**Abstract.** Laser Powder Bed Fusion (LPBF) has been widely adopted in many industrial sectors such as biomedical, automotive, and aerospace thanks to the possibility to produce objects with complex shapes and meet customers' needs. Despite all the advantages that LPBF can offer, the rise of residual stress due to the high thermal gradients generated during the process can limit its application. This is the case with thin-wall structures where the build-up of residual stress can compromise the success of the printing process. Being able to print this structure can be useful in fuel cell applications where the implementation of cooling channels in bipolar plates can improve their performance. This paper aims to provide guidelines for designing thin-wall structures produced by LPBF processes through numerical simulations by understanding the effect of residual stress on part distortion.

**Introduction**

Nowadays, where the energy demand is constantly increasing and fossil-based fuels are the principal contributors to meeting the global energy need, hydrogen-based technologies can offer a good alternative to non-renewable energy sources that are doomed to be exhausted one day. Moreover, thanks to its inherently immense energy content, availability, and zero-emission, hydrogen is considered the future fuel in the transport sector. Particularly, hydrogen can be used in electrical power generation systems such as fuel cells for vehicles [1]. These kinds of fuel cells which use chemical hydrogen bonds to store electricity are known as proton exchange membrane (PEM) fuel cells. Bipolar plates (BSs) are the main component of PEM fuel cells in terms of structural strength, weight, and cost [2]. Usually, BPs are made of graphite because of its high conductivity, high corrosion resistance, low density, and ease of fabrication. However, manufacturing of graphite plates under a certain thickness can be difficult due to their brittle nature. Furthermore, it is very challenging to create flow channels in order to apply liquid cooling and improve the PEM fuel cells' performance [3]. As a consequence, metallic plates are taking increasing interest due to their high thermal and electrical conductivity, good gas impermeability, high strength, and good formability. In addition, by using metals as the material to produce metallic plates, additive manufacturing (AM) methods can be adopted [4]. AM consists of a layer-by-layer production fashion in which the part is obtained by adding layers of material. Thus, allowing the fabrication of parts with a very complex shape or structure. In the case of metallic parts, laser powder bed fusion (L-PBF) is among the more suitable AM technologies. In L-PBF a powder bed is deposited on a build plate within a chamber filled with an inert gas and the metallic powder is selectively melted with a laser heat source according to the slicing file [5].

Titanium alloys can be used to produce metallic plates thanks to their high mechanical strength, high corrosion resistance, and lightweight [6,7]. In this study, Ti-6Al-4V was used in the L-PBF process to print bipolar plates. These BPs are characterized by thin-wall structures which can be

<figure><img src="image_262.png" /></figure>

Content from this work may be used under the terms of the Creative Commons Attribution 3.0 license. Any further distribution of this work must maintain attribution to the author(s) and the title of the work, journal citation and DOI. Published under license by Materials Research Forum LLC.
difficult to print due to the rise of residual stress during the L-PBF process that can result in part distortion [8,9]. Numerical simulations can be a useful tool to predict part distortion and avoid job failure. Furthermore, compensation strategies can be adopted based on the results from the numerical simulation. In order to simulate the L-PBF process of large parts, a layer-by-layer approach is necessary to reduce the computational cost [10]. In this paper, a finite element method (FEM) is used to simulate the fabrication of a thin-walled structure through the L-PBF process of Ti-6Al-4V powder. The Bipolar Plate geometry was printed and acquired in order to validate the numerical model. An acquisition method based on the fringe projection approach was employed to obtain the actual geometries of the produced part. The aim is to understand the development of residual stress during the process and its effect on part distortion. Moreover, the influence of the number of real layers in the computational one was evaluated in terms of model accuracy and time efficiency.

## FEM model and printing process

Simulation Set-Up. Numerical simulations were carried out to understand the influence of residual stress on the deformation mechanism of thin-wall structures for fuel cell application fabricated with L-PBF. To this aim, the FEM commercial software DEFORM-3D™ v12.0 (V12.0, SFC, Columbus, OH, USA) was used. Because of the large dimensions of the bipolar plates being analyzed (144 x 2 x 114 mm³), a layer-by-layer approach was adopted. With this method, the whole layer is heated simultaneously during the time step. The element activation is strictly correlated on the used voxel mesh according to the born-dead-elements technique. In this way, by the means of a searching algorithm, the elements within the voxel mesh that are influenced by the heat source are activated. The elements not affected by the power input will not be activated and will not have any role in the computational analysis. For more details about the element activation mode and the use of the voxel mesh refer to [11]. The thin-wall structure designed with Autodesk Fusion 360 was simulated on a base plate of 160 x 160 x 25 mm³ (Fig. 1) with which sticking conditions were applied to avoid detachment during the simulation of the L-PBF process.

<figure><img src="image_259.png" /><figcaption>Fig. 1 Thin-wall structure and base plate used for the simulation of the L-PBF process with DEFORM-3D™ v12.0.</figcaption></figure>

In this study, two different voxel meshes were compared named V1 (0.5 x 0.5 x 0.3 mm³) and V2 (0.5 x 0.5 x 0.6 mm³) to find the best compromise between simulation accuracy and simulation time. In this way, in the V1 mesh 10 real layers were grouped in one computational layer, while for the V2 mesh, 20 real layers were included in one computational layer. The bipolar plate mesh
difficult to print due to the rise of residual stress during the L-PBF process that can result in part distortion [8,9]. Numerical simulations can be a useful tool to predict part distortion and avoid job failure. Furthermore, compensation strategies can be adopted based on the results from the numerical simulation. In order to simulate the L-PBF process of large parts, a layer-by-layer approach is necessary to reduce the computational cost [10]. In this paper, a finite element method (FEM) is used to simulate the fabrication of a thin-walled structure through the L-PBF process of Ti-6Al-4V powder. The Bipolar Plate geometry was printed and acquired in order to validate the numerical model. An acquisition method based on the fringe projection approach was employed to obtain the actual geometries of the produced part. The aim is to understand the development of residual stress during the process and its effect on part distortion. Moreover, the influence of the number of real layers in the computational one was evaluated in terms of model accuracy and time efficiency.

## FEM model and printing process

Simulation Set-Up. Numerical simulations were carried out to understand the influence of residual stress on the deformation mechanism of thin-wall structures for fuel cell application fabricated with L-PBF. To this aim, the FEM commercial software DEFORM-3D™ v12.0 (V12.0, SFC, Columbus, OH, USA) was used. Because of the large dimensions of the bipolar plates being analyzed (144 x 2 x 114 mm³), a layer-by-layer approach was adopted. With this method, the whole layer is heated simultaneously during the time step. The element activation is strictly correlated on the used voxel mesh according to the born-dead-elements technique. In this way, by the means of a searching algorithm, the elements within the voxel mesh that are influenced by the heat source are activated. The elements not affected by the power input will not be activated and will not have any role in the computational analysis. For more details about the element activation mode and the use of the voxel mesh refer to [11]. The thin-wall structure designed with Autodesk Fusion 360 was simulated on a base plate of 160 x 160 x 25 mm³ (Fig. 1) with which sticking conditions were applied to avoid detachment during the simulation of the L-PBF process.

<figure><img src="image_263.png" /><figcaption>Fig. 1 Thin-wall structure and base plate used for the simulation of the L-PBF process with DEFORM-3D™ v12.0.</figcaption></figure>

In this study, two different voxel meshes were compared named V1 (0.5 x 0.5 x 0.3 mm³) and V2 (0.5 x 0.5 x 0.6 mm³) to find the best compromise between simulation accuracy and simulation time. In this way, in the V1 mesh 10 real layers were grouped in one computational layer, while for the V2 mesh, 20 real layers were included in one computational layer. The bipolar plate mesh
consists of 709760 tetrahedral elements and was considered elastoplastic for thermomechanical analysis. For the base plate, instead, 32000 tetrahedral elements were used and the object was modeled as rigid. Ti-6Al-4V was assigned for both the bipolar plate and the build platform as the material. Heat exchange boundary conditions with the environment were defined for both objects by considering the build chamber at 40°C and filled with argon ($h_{conv} = 1$ W/m²·°C). To reduce the computational cost and simplify the model, the heat exchange with the powder bed was not taken into consideration. Moreover, the initial built platform temperature was set at 200°C. In order to solve the thermomechanical problem, a Newton-Rapson iteration method was used within DEFORM-3D™ v12.0 using a MUMPS (Multifrontal Massively Parallel Sparse) solver and a 12th Gen Intel (R) Core (TM) i9-12900 2.40 GHz processor.

Experimental Details. The bipolar plate was fabricated with a 3D system SLM 280 HL machine (SLM Solutions, Lubeka, Germany) to confirm the results from the numerical simulation. A laser power input $P$ equal to 350 W was used together with a scan speed (v) of 1400 mm/s, a hatch spacing (h) equal to 120 µm, and a layer thickness (t) of 30 µm. The platform was preheated to 200°C to reduce the amount of residual stress and the build chamber was filled with argon to lower the oxygen level below 0.1%. In this study, Ti-6Al-4V powder with a Gaussian size distribution of 20-63 µm and spherical shape were used in the L-PBF process.

The printed BP part was removed from the platform and a 3D COMET V system (Steinbichler, Germany) was used to acquire the geometry of the sample. The fringe projection approach is the basis for this 3D acquisition method, where a structured light pattern is projected onto the surface of the object using a projection unit. The image acquisition unit captures the image of the fringe pattern. To measure the component shape at different rotation angles, the printed specimen is placed on a rotating table during the 3D scanning procedure. A three-dimensional representation of the part is created by combining and meshing the point-cloud data sets. Quantifying the distortion of the part was achieved by using the difference in the y-axis (Dy) between the original CAD model and the geometry acquired after the printing process as represented in Fig. 2.

<figure><img src="image_260.png" /><figcaption>Fig. 2 Comparison between the CAD model of the bipolar plate and the 3D scanned AM part a) 3D view and b) side view with the distortion measurement along the y direction.</figcaption></figure>

## Results and discussions

Producing thin-wall structures with L-PBF is still a challenge, especially when large parts are considered as in the case of bipolar plates. The development of residual stresses during the printing process was highlighted as a key factor when creating the BP geometry using additive manufacturing. Significant in-process distortion (as shown in Fig. 3) caused the printing process to end prematurely for a z value of 84.7 mm, corresponding to layer 2851. In detail, a distortion
consists of 709760 tetrahedral elements and was considered elastoplastic for thermomechanical analysis. For the base plate, instead, 32000 tetrahedral elements were used and the object was modeled as rigid. Ti-6Al-4V was assigned for both the bipolar plate and the build platform as the material. Heat exchange boundary conditions with the environment were defined for both objects by considering the build chamber at 40°C and filled with argon ($h_{conv} = 1$ W/m²·°C). To reduce the computational cost and simplify the model, the heat exchange with the powder bed was not taken into consideration. Moreover, the initial built platform temperature was set at 200°C. In order to solve the thermomechanical problem, a Newton-Rapson iteration method was used within DEFORM-3D™ v12.0 using a MUMPS (Multifrontal Massively Parallel Sparse) solver and a 12th Gen Intel (R) Core (TM) i9-12900 2.40 GHz processor.

Experimental Details. The bipolar plate was fabricated with a 3D system SLM 280 HL machine (SLM Solutions, Lubeka, Germany) to confirm the results from the numerical simulation. A laser power input $P$ equal to 350 W was used together with a scan speed (v) of 1400 mm/s, a hatch spacing (h) equal to 120 µm, and a layer thickness (t) of 30 µm. The platform was preheated to 200°C to reduce the amount of residual stress and the build chamber was filled with argon to lower the oxygen level below 0.1%. In this study, Ti-6Al-4V powder with a Gaussian size distribution of 20-63 µm and spherical shape were used in the L-PBF process.

The printed BP part was removed from the platform and a 3D COMET V system (Steinbichler, Germany) was used to acquire the geometry of the sample. The fringe projection approach is the basis for this 3D acquisition method, where a structured light pattern is projected onto the surface of the object using a projection unit. The image acquisition unit captures the image of the fringe pattern. To measure the component shape at different rotation angles, the printed specimen is placed on a rotating table during the 3D scanning procedure. A three-dimensional representation of the part is created by combining and meshing the point-cloud data sets. Quantifying the distortion of the part was achieved by using the difference in the y-axis (Dy) between the original CAD model and the geometry acquired after the printing process as represented in Fig. 2.

<figure><img src="image_264.png" /><figcaption>Fig. 2 Comparison between the CAD model of the bipolar plate and the 3D scanned AM part a) 3D view and b) side view with the distortion measurement along the y direction.</figcaption></figure>

## Results and discussions

Producing thin-wall structures with L-PBF is still a challenge, especially when large parts are considered as in the case of bipolar plates. The development of residual stresses during the printing process was highlighted as a key factor when creating the BP geometry using additive manufacturing. Significant in-process distortion (as shown in Fig. 3) caused the printing process to end prematurely for a z value of 84.7 mm, corresponding to layer 2851. In detail, a distortion
value of 6.53 mm was observed. In the simulated process, although the level of distortion of the sample is evaluated with a reduced percentage of error, the process does not stop during the creation of the 2851 layer as the development of the distortions, connected to the residual stresses, occurs only following the cooling down phase. This is due to the grouping of several real layers in the computational one and the neglecting of the resting time between consecutive layers. In fact, without considering the interlayer resting time there is no time for the residual stress to develop during the simulation of the process. On the other hand, considering the interlayer resting time results in a significant increase in the simulation time.

<figure><img src="image_261.png" /><figcaption>Fig. 3 Post-exposure images acquired during the printing process for a) layer 2850 and b) layer 2851, red arrows indicate the distortion of the part. These images represent the cross-section of the bipolar plate (plane XY), perpendicular to the build direction.</figcaption></figure>

In Fig. 4 it is possible to observe the temperature profiles obtained from the numerical simulation at the end of the printing process.

<figure><img src="image_262.png" /><figcaption>Fig. 4 Temperature profiles for V1 and V2 from the numerical simulation at the end of the printing process.</figcaption></figure>

It can be noted how the temperature profile, in the case of a voxel mesh dimension of 0.3 mm along the z direction (V1), is similar to the one obtained for a voxel mesh dimension of 0.6 mm (V2). A significant difference between the V1 and V2 temperature profiles appears only in the upper part of the component where the temperature is above 700 °C.

Particularly, it can be noted how a finer voxel mesh (V1) better represents the temperature distribution along the z-direction for the last scanned layers with higher temperatures.

The thermal gradient along the z-direction is responsible for the development of the residual stress that can be observed in the numerical simulation after the cooling phase.
value of 6.53 mm was observed. In the simulated process, although the level of distortion of the sample is evaluated with a reduced percentage of error, the process does not stop during the creation of the 2851 layer as the development of the distortions, connected to the residual stresses, occurs only following the cooling down phase. This is due to the grouping of several real layers in the computational one and the neglecting of the resting time between consecutive layers. In fact, without considering the interlayer resting time there is no time for the residual stress to develop during the simulation of the process. On the other hand, considering the interlayer resting time results in a significant increase in the simulation time.

<figure><img src="image_265.png" /><figcaption>Fig. 3 Post-exposure images acquired during the printing process for a) layer 2850 and b) layer 2851, red arrows indicate the distortion of the part. These images represent the cross-section of the bipolar plate (plane XY), perpendicular to the build direction.</figcaption></figure>

In Fig. 4 it is possible to observe the temperature profiles obtained from the numerical simulation at the end of the printing process.

<figure><img src="image_266.png" /><figcaption>Fig. 4 Temperature profiles for V1 and V2 from the numerical simulation at the end of the printing process.</figcaption></figure>

It can be noted how the temperature profile, in the case of a voxel mesh dimension of 0.3 mm along the z direction (V1), is similar to the one obtained for a voxel mesh dimension of 0.6 mm (V2). A significant difference between the V1 and V2 temperature profiles appears only in the upper part of the component where the temperature is above 700 °C.

Particularly, it can be noted how a finer voxel mesh (V1) better represents the temperature distribution along the z-direction for the last scanned layers with higher temperatures.

The thermal gradient along the z-direction is responsible for the development of the residual stress that can be observed in the numerical simulation after the cooling phase.
In Fig. 5 the residual stresses, for the V1 and V2 cases, developed in the z-direction are shown.

<figure><img src="image_263.png" /><figcaption>Fig. 5 Residual stress profiles along the z-direction for V1 and V2 obtained from the numerical simulation before the part deformation. The dashed red line highlights the difference in the maximum compression region between V1 and V2.</figcaption></figure>

These stresses caused a large displacement along the y-direction (Fig. 6).

<figure><img src="image_264.png" /><figcaption>Fig. 6 Displacement profiles along the y-direction for V1 and V2 obtained from the numerical simulation in the ZX and ZY planes.</figcaption></figure>

It can be noted that the V1 case study shows a larger area of maximum compression stress with respect to the V2 case as highlighted by the dashed red line. Overall, the V1 voxel mesh allows a better representation of the temperature and residual stress profiles. As a consequence, the displacement from the numerical simulation for the V1 case is closer to the one obtained from the experimental measurement with respect to the V2 case. Also, as can be seen from the displacement maps in the ZX plane, the distribution of the displacement along the y-direction (Dy) is not symmetrical. Moreover, regarding the y displacements, the V1 voxel mesh offers a better match with the real acquired geometry. The difference between the simulated y-displacement and the actual one was evaluated by considering the Dy value at the z = 85.5 mm, which corresponds to the z coordinate of the last printed layer before the premature end of the printing process. As
In Fig. 5 the residual stresses, for the V1 and V2 cases, developed in the z-direction are shown.

<figure><img src="image_267.png" /><figcaption>Fig. 5 Residual stress profiles along the z-direction for V1 and V2 obtained from the numerical simulation before the part deformation. The dashed red line highlights the difference in the maximum compression region between V1 and V2.</figcaption></figure>

These stresses caused a large displacement along the y-direction (Fig. 6).

<figure><img src="image_268.png" /><figcaption>Fig. 6 Displacement profiles along the y-direction for V1 and V2 obtained from the numerical simulation in the ZX and ZY planes.</figcaption></figure>

It can be noted that the V1 case study shows a larger area of maximum compression stress with respect to the V2 case as highlighted by the dashed red line. Overall, the V1 voxel mesh allows a better representation of the temperature and residual stress profiles. As a consequence, the displacement from the numerical simulation for the V1 case is closer to the one obtained from the experimental measurement with respect to the V2 case. Also, as can be seen from the displacement maps in the ZX plane, the distribution of the displacement along the y-direction (Dy) is not symmetrical. Moreover, regarding the y displacements, the V1 voxel mesh offers a better match with the real acquired geometry. The difference between the simulated y-displacement and the actual one was evaluated by considering the Dy value at the z = 85.5 mm, which corresponds to the z coordinate of the last printed layer before the premature end of the printing process. As
reported in Fig. 1, this point will result in the z coordinate equal to 84.7 mm after the occurred distortion of 6.53 mm. The Dy values obtained for the V1 and V2 cases are reported in Table 1, equal to 6.03 and 5.72 respectively.

<table><caption>Table 2. Output of the numerical model in terms of displacement, accuracy, and simulation time.</caption><thead><tr><th></th><th>V1</th><th>V2</th></tr></thead><tbody><tr><td>Dy [mm]</td><td>6.03</td><td>5.72</td></tr><tr><td>Err(Dy)%</td><td>7.6</td><td>12.4</td></tr><tr><td>Simulation Time (gg:hh:mm:ss)</td><td>02:07:39:52 (200392 s)</td><td>00:20:44:25 (74665 s)</td></tr><tr><td>Time Reduction [%]</td><td>53.6</td><td>82.7</td></tr></tbody></table>

Moreover, the accuracy of the simulation model was expressed in terms of percentage error (Err(Dy)%) between the Dy measured and simulated at z = 85.5 mm. As shown in Table 1, an Err(Dy) % value of 7.6 was found for the V1 voxel mesh and 12.4 for the V2 one. A further analysis regarding the simulation times to simulate the entire process was carried out. In detail, the simulation time was compared to the real printing time introducing a time reduction factor. This factor was calculated as the percentage difference between the printed time and the simulation time of the L-PBF process, by considering a printing time of 432000 s. The V1 voxel mesh resulted in a simulation time of 200392 s and a time reduction of 53.6%, while the V2 resulted in a simulation time of 74665 s and a time reduction of 82.7%. It can be observed how the simulation time is significantly reduced by using a voxel mesh dimension of 0.6 mm (V2) including 20 real layers in one computational layer. The considerable time reduction obtained by using half of the V1 voxel mesh dimension results in one eight of V1 voxel mesh elements. The reduction of the computational time is at the expense of the accuracy of the model as can be seen in Table 1.

## Conclusions

In this study, the simulation of the L-PBF process to produce a large Ti-6Al-4V Bipolar Plate was carried out. The influence of the residual stress on the distortion of this thin-walled structure was evaluated together with the accuracy of the numerical model. The main findings of this study are the following:

- Due to the significant development of the residual stresses along the z-direction is not possible to print a Bipolar Plate of the dimension used in this study without any further strategy.

- Even if it is not possible to reproduce exactly when the printing process will end prematurely, the reported displacement from the numerical simulation after the cooling time is very close to the real ones.

- The number of real printed layers included in the computational one has a big impact on the numerical model results both in terms of accuracy and simulation time. In detail, the V1 voxel mesh with 10 real printed layers in one computational layer leads to a better fit with the experimental results. The V2 voxel mesh with 20 real layers in one computational layer, instead, leads to a greater simulation time reduction.

- The good compromise obtained from the numerical model in the case of V1 voxel mesh makes this model a suitable tool for helping engineers in the design of the L-PBF process.