Technical field

The present invention belongs to the field of materials that can be processed using additive manufacturing (AM). In particular, the present invention relates to an aluminium-based alloy and a 3D-printed product obtained by processing said alloy by means of AM techniques.

Background art

In the last few years, additive manufacturing (AM) has attracted more and more interest in a wide range of technological fields as it allows the manufacture of products with complex shapes and small amount of waste. The most widespread AM technology, in particular to produce prototypes and series of 3D-printed products, is the Laser Powder Bed Fusion (LPBF), also known as Selective Laser Melting (SLM), Direct Metal Laser Sintering (DMLS), Laser Cusing, Additive Laser Manufacturing (ALM), etc, which is a powder bed fusion process used to produce objects from powdered materials using one or more lasers to selectively fuse or melt the particles at the surface, layer by layer, in an enclosed chamber.

Although there is a strong request for a wide range of alloys for AM, the number of materials available on the market is still limited. This is mainly due to the poor processability of many standard alloys normally used, for example, for traditional casting or plastic deformation. In particular, the lack of high-strength aluminum alloys that can also be processed using AM remains a big issue for the diffusion of such innovative techniques in many technological sectors such as the aerospace and automotive industries. Many high strength aluminum alloys currently available on the market, such as the alloys belonging to the 2xxx, 6xxx and 7xxx series, are difficult to process with additive manufacturing because they have high susceptibility to hot cracking. Hot cracking is a metallurgical phenomenon that can occur in casting and welding of many metallic materials including high strength Al alloys. During the terminal stage of solidification, thermal shrinkage promotes build-up of tensile stresses in the semi-solid metal, causing nucleation and propagation of cracks mainly along grain boundaries. Therefore, there is still a need to provide high-strength aluminum alloys that can be processed using additive manufacturing, in particular laser-based AM, without incurring the above-mentioned phenomena and that, at the same time, allow to obtain final products with the desired mechanical and functional properties. A few works on the modification of chemical composition of high-strength Al alloys have been recently published. For example, such research works are focused on the effect of the addition of nucleants on the solidification behavior of Al alloys processed by LPBF.

H. Zhang et al. ("Effect of Zirconium addition on crack, microstructure and mechanical behavior of selective laser melted Al-Cu-Mg alloy", Scr. Mater., vol. 134, pp. 6-10, 2017) studied the properties of an Al-Cu-Mg alloy modified with Zr and the corresponding specimens obtained by 3D printing (specifically by Selective Laser Melting), compared to the properties of the unmodified Al-Cu-Mg alloy and the corresponding specimens. The modification of the starting Al-Cu-Mg alloy was made by mechanically mixing the alloy in the form of an atomized powder with a micrometric Zr powder in a quantity equal to 2% by weight with respect to the total weight of the starting alloy.

This study showed a reduction in the hot cracking phenomenon for the specimens thus obtained compared to the specimens produced with the unmodified alloy, due to the grain refining effect triggered by the addition of the Zr powder.

P. Wang et al. (“A heat treatable TIB2/AI-3.5Cu-1.5Mg-1 Si composite fabricated by selective laser melting: microstructure, heat treatment and mechanical properties”, Compos. Part B Eng., vol. 147, pp. 162-168, 2018) investigated the properties of an Al-Cu-Mg-Si alloy modified by addition of TiB2 particles and the corresponding specimens obtained by AM, compared to the properties of the unmodified Al-Cu-Mg-Si alloy and the corresponding specimens. The modification of the starting Al-Cu-Mg-Si alloy was made by mechanically mixing the alloy in the form of atomized powder with a micrometric TiB2 powder. This study showed the positive effect of the addition of 5 vol.% TiB2 particles on the processability of the Al-Cu-Mg-Si alloy.

Likewise, X. Wen et al. (“Laser solid forming additive manufacturing T1B2 reinforced 2024AI composite: microstructure and mechanical properties, Materials Science and Engineering: A, Vol. 745, pp. 319-325, 2019) showed that the addition of TiB2 particles within the Al-Cu-Mg alloy (AA2024) resulted in a refinement of the grain structure and improved mechanical properties of the final product.

However, none of these works concerns the use of an Al-Ti-Cu-Mg-B-Ni- Fe-Si powder.

Furthermore, Al alloys such as Al-Cu or Al-Cu-Mg are also considered as materials with a particularly low AM processability (especially when L-PBF technique is used) and highly susceptible to yield solidification cracks in the manufactured product. However, the Applicant found that an Al-Ti-Cu- Mg-B-Ni-Fe-Si alloy powder shows excellent processability by AM. Furthermore, the proposed alloy of the invention yields a 3D printed product endowed with high mechanical characteristics and performance both at room and high temperature ranges.

Summary of the invention

The present invention relates to an Al-Ti-Cu-Mg-B-Ni-Fe-Si powder for additive manufacturing, preferably for laser-based additive manufacturing, more preferably for Laser Powder Bed Fusion (LPBF), comprising or consisting of:

- from 2 to 4 wt. % of Ti;

- from 0.5 to 1 .5 wt. % of B;

- from 2 to 3.5 wt. % of Cu;

- from 1 to 2 wt. % of Mg;

- from 0.5 to 1 .2 wt. % of Fe; - from 0.8 to 1 .5 wt. % of Ni;

- from 0.1 to 0.5 wt. % of Si;

- balance of Al.

The present invention also relates to a 3D-printed product obtained by:

(i) processing said Al-Ti-Cu-Mg-B-Ni-Fe-Si powder by means of an additive manufacturing technique, preferably by means of a laser-based additive manufacturing technique, more preferably by means of Laser Powder Bed Fusion (LPBF); said 3D-printed product comprising an Al-Ti- Cu-Mg-B-Ni-Fe-Si alloy comprising or consisting of:

- from 2 to 4 wt. % of Ti;

- from 0.5 to 1 .5 wt. % of B;

- from 2 to 3.5 wt. % of Cu;

- from 1 to 2 wt. % of Mg;

- from 0.5 to 1 .2 wt. % of Fe;

- from 0.8 to 1 .5 wt. % of Ni;

- from 0.1 to 0.5 wt. % of Si;

- balance of Al.

According to an embodiment of the invention, the 3D-printed product can optionally be further subjected to:

(ii) stress relief at a temperature comprised between 100 and 400 °C; and/or

(iii)a heat treatment comprising solution annealing at a temperature preferably comprised between 400 and 600 °C followed by water quenching; and/or

(iv) aging at a temperature preferably comprised between 120 and 220 °C for a time preferably ranging from 1 hour to 48 hours.

Preferably, when the 3D-printed product according to the present invention is not subjected to the heat treatment of steps (ii) and/or (iii) and/or (iv), it comprises second-phase compounds preferably selected from the group consisting of: TiB2, AI?Cu4Ni, AhTi, AhFeNi, O-AhCu and |3-Mg2Si and combination thereof; said second-phase compounds being positioned at cell boundaries and/or within cell cores of the solidification structure of the 3D-printed product.

According to another embodiment of the present invention, when the 3D- printed product is subjected to step (iii) or to steps (ii)+(iii), it comprises second-phase compounds preferably selected from the group consisting of: TiB2, AI?Cu4Ni, AhTi, AlgFeNi, and combination thereof; said second- phase compounds being preferably in the form of coarsen particles. Preferably, said coarsen particles have a mean diameter, measured by electron microscopy, comprised between 0.1 and 5 pm.

According to a preferred embodiment, the 3D-printed product of the present invention has at least one of the following characteristics: a) material relative density higher than 99.4%; b) absence of solidification cracks within the microstructure; c) yield strength or upper yield strength comprised between 300 and 500 MPa, measured at room temperature; d) ultimate tensile strength comprised between 400 and 600 MPa, measured at room temperature; e) hardness higher than 95 HV; f) elongation at fracture higher than 2%, measured at room temperature; g) yield strength or upper yield strength higher than 250 MPa up to 200°C and higher than 90MPa up to 300°C; h) ultimate tensile strength higher than 270 MPa up to 200°C and higher than 10OMPa up to 300°C.

According to an embodiment of the present invention, the 3D-printed product possesses a refined crack-free equiaxed grain structure.

The present invention also refers to a use of the 3D-printed product for high temperature applications (up to 300°C), such as aerospace and automotive/racing applications, preferably for the realization of engine heads, heat exchangers, pumps.

Brief description of drawings Figure 1 shows: a) Low magnification LOM image, b) high magnification LOM image and c) SEM micrographs of the section parallel to the building direction of the as-built sample 1 ; d) Low magnification LOM image and e) SEM micrographs of the section parallel to the building direction of the as- built sample 1 .1 as reported in Example 3.

Figure 2 shows: a) XRD diffractograms of the as-built sample 1 ; background signal in the low angle 20 region of the b) as-built and c) solution treated samples 1 , as reported in Example 4.

Figure 3 shows: a) XRD diffractograms of the as-built sample 1.1 ; background signal in the low angle 20 region of the b) as-built and c) solution treated samples 1.1 , as reported in Example 4. Figure 4 shows SEM images of aged a) as-built and b) solution treated samples 1 after Keller’s etching as reported in Example 5.

Figure 5 shows room-temperature tensile curves of the as-built, T5 and T6 a) Z- and b) XY-samples 1 ; c) XY- and d) Z-sample 1.1 ; as reported in Example 6.

Figure 6 shows: a) YS vs. temperature, b) UTS vs. temperature and c) elongation at fracture vs temperature plots of the sample produced with an Al-Cu-Mg-Ni-Fe-Si alloy (AI2618 wrought), the samples produced with the pre-alloyed Al-Ti-Cu-Mg-B-Ni-Fe-Si powder according to the present invention (sample 1 and 1.1 ), and of the sample produced with a commercial AISil OMg alloy, in T6 conditions, as reported in Example 6.

Detailed description of preferred embodiments of the invention

For the purposes of the present invention, “Laser Powder Bed Fusion (L- PBF)”, “Selective Laser Melting (SLM)”, “Direct Metal Laser Sintering (DMLS)”, “Laser Cusing”, and “Additive Laser Manufacturing (ALM)”, are used as interchangeable synonyms.

The present invention relates to an Al-Ti-Cu-Mg-B-Ni-Fe-Si powder suitable for being processed by additive manufacturing techniques, preferably by means of laser-based additive manufacturing, more preferably by means of Laser Powder Bed Fusion (L-PBF). According to one embodiment of the present invention, said Al-Ti-Cu-Mg- B-Ni-Fe-Si powder is a pre-alloyed Al-Ti-Cu-Mg-B-Ni-Fe-Si powder. Preferably, said pre-alloyed Al-Ti-Cu-Mg-B-Ni-Fe-Si powder is obtained by means of gas atomization.

Said Al-Ti-Cu-Mg-B-Ni-Fe-Si powder comprises or consists of:

- from 2 to 4 wt. %, preferably from 3 to 4 wt.%, more preferably from 2 to 3 wt.% of Ti;

- from 0.5 to 1 .5 wt. %, preferably from 1 to 1 .5 wt%, more preferably from 0.5 to 1 wt.% of B;

- from 2 to 3.5 wt. %, preferably from 2 to 2.5 wt.% of Cu;

- from 1 to 2 wt. % of Mg;

- from 0.5 to 1 .2 wt. % of Fe;

- from 0.8 to 1 .5 wt. % of Ni;

- from 0.1 to 0.5 wt. % of Si;

- balance of Al.

According to a preferred embodiment of the invention, said Al-Ti-Cu-Mg-B- Ni-Fe-Si powder comprises or consists of:

- from 3 to 4 wt. % of Ti;

- from 1 to 1 .5 wt. % of B;

- from 2 to 3.5 wt. % of Cu;

- from 1 . to 2 wt. % of Mg;

- from 0.5 to 1 .2 wt. % of Fe;

- from 0.8 to 1 .5 wt. % of Ni;

- from 0.1 to 0.5 wt. % of Si;

- balance of Al.

According to another preferred embodiment of the invention, said Al-Ti-Cu- Mg-B-Ni-Fe-Si powder comprises or consists of:

- from 2 to 3 wt. % of Ti;

- from 0.5 to 1 wt. % of B;

- from 2 to 2.5 wt. % of Cu;

- from 1 to 2 wt. % of Mg; - from 0.5 to 1 .2 wt. % of Fe;

- from 0.8 to 1 .5 wt. % of Ni;

- from 0.1 to 0.5 wt. % of Si;

- balance of Al.

The present invention also relates to a 3D-printed product obtained/obtainable by:

(i) processing the Al-Ti-Cu-Mg-B-Ni-Fe-Si powder as described above by means of an additive manufacturing technique, preferably by means of laser-based additive manufacturing, more preferably by means of Laser Powder Bed Fusion (L-PBF).

The 3D-printed product according to the present invention comprises the Al-Ti-Cu-Mg-B-Ni-Fe-Si alloy, which is formed from the Al-Ti-Cu-Mg-B-Ni- Fe-Si powder after step (i) and comprises or consists of:

- from 2 to 4 wt. %, preferably from 3 to 4 wt.%, more preferably from 2 to 3 wt.% of Ti;

- from 0.5 to 1 .5 wt. %, preferably from 1 to 1 .5 wt%, more preferably from 0.5 to 1 wt.% of B;

- from 2 to 3.5 wt. %, preferably from 2 to 2.5 wt.% of Cu;

- from 1 to 2 wt. % of Mg;

- from 0.5 to 1 .2 wt. % of Fe;

- from 0.8 to 1 .5 wt. % of Ni;

- from 0.1 to 0.5 wt. % of Si;

- balance of Al.

According to a preferred embodiment of the invention, said Al-Ti-Cu-Mg-B- Ni-Fe-Si alloy comprises or consists of:

- from 3 to 4 wt. % of Ti;

- from 1 to 1 .5 wt. % of B;

- from 2 to 3.5 wt. % of Cu;

- from 1 to 2 wt. % of Mg;

- from 0.5 to 1 .2 wt. % of Fe;

- from 0.8 to 1 .5 wt. % of Ni; - from 0.1 to 0.5 wt. % of Si;

- balance of Al.

According to a preferred embodiment of the invention, said Al-Ti-Cu-Mg-B- Ni-Fe-Si alloy comprises or consists of:

- from 2 to 3 wt. % of Ti;

- from 0.5 to 1 wt. % of B;

- from 2 to 2.5 wt. % of Cu;

- from 1 to 2 wt. % of Mg;

- from 0.5 to 1 .2 wt. % of Fe;

- from 0.8 to 1 .5 wt. % of Ni;

- from 0.1 to 0.5 wt. % of Si;

- balance of Al.

According to a particularly preferred embodiment of the invention, the 3D- printed product consists essentially of said Al-Ti-Cu-Mg-B-Ni-Fe-Si alloy.

According to an even more preferred embodiment of the invention, the 3D- printed product consists of said Al-Ti-Cu-Mg-B-Ni-Fe-Si alloy.

Preferably, step (i) is carried out in an apparatus having a chamber and comprises the steps of:

(ia) forming a layer of the Al-Ti-Cu-Mg-B-Ni-Fe-Si powder on a base plate, preferably while maintaining an oxygen-low environment;

(ib) melting the powder locally by exposing the powder to an energy beam, preferably a laser beam or an electron beam, during a period of time sufficient to form at least one melt pool;

(ic) letting the melted powder in the at least one melt pool solidify into the Al-Ti-Cu-Mg-B-Ni-Fe-Si alloy;

(id) applying further additional layers of the Al-Ti-Cu-Mg-B-Ni-Fe-Si powder on top of the previous layer by repeating the steps (ia)-(ic), wherein step (ib) comprises placing the additional layer on top of the previous layer once it is solidified after step (ic).

Preferably, the 3D-printed product being built in steps (ia)-(id) is kept heated at a temperature comprised between room temperature and 400°C, preferably between 40 and 200 °C, during the above-mentioned steps (ia)-(id).

Preferably, the environment in the chamber may comprise inert gases, preferably selected from the group consisting of argon, nitrogen (N2) or a combination thereof.

Preferably the energy beam is swept across the powder in a pattern.

According to an embodiment of the present invention, the 3D-printed product, obtained after step (i) can optionally be further subjected to:

(ii) stress relief at a temperature comprised between 100 and 400 °C; and/or

(iii) a heat treatment comprising solution annealing at a temperature preferably comprised between 400 and 600 °C followed by water quenching; and/or

(iv) aging at a temperature preferably comprised between 120 and 220 °C, more preferably between 150 and 200 °C, for a time preferably ranging from 1 hour to 24 hours, more preferably from 3 to 24 hours.

Preferably, said solution annealing is carried out at a temperature between 500 and 600 °C and preferably for a time comprised between 0.5 and 2 hours.

Without wishing to be bound to a specific theory, the Applicant has nevertheless found that the 3D-printed product according to the present invention, when not subjected to step (ii) and/or (iii) and/or (iv), may comprise second-phase compounds preferably selected from the group consisting of: TiB2, AI?Cu4Ni, AhTi, AlgFeNi, 0-AI2CU and p-Mg Si and combination thereof.

Said second-phase compounds are positioned at cell boundaries and/or within cell cores of the solidification structure of the 3D-printed product.

Instead, when the 3D-printed product is subjected to step (iii) or to steps (ii)+(iii), the Applicant has found that said step(s) modify(ies) the microstructure of the 3D-printed product as it preferably comprises second-phase compounds selected from the group consisting of: TiB2, AI?Cu4Ni, AhTi, AlgFeNi, and combination thereof (namely, it does not comprise Q-AhCu and p-MggSi phases anymore or it comprises said phases in a reduced amount since their dissolution is promoted by the solution annealing); said second-phase compounds preferably being in form of coarsen particles preferably having a mean diameter, measured by electron microscopy, comprised between 0.1 and 5 pm, more preferably between 0.2 and 2 pm, even more preferably between 0.5 and 1 .2 pm.

According to an embodiment of the present invention, the 3D-printed product has at least one of the following characteristics: a) material relative density higher than 99.4%, preferably comprised between 99.80 and 100 %, more preferably between 99.85 and 100%; b) absence of solidification cracks within the microstructure; c) yield strength or upper yield strength comprised between 300 and 500 MPa, preferably between 370 and 500 MPa, more preferably between 460 and 500 MPa, measured at room temperature; d) ultimate tensile strength comprised between 400 and 600 MPa, preferably between 440 and 490 MPa, more preferably between 450 and 490 MPa, measured at room temperature; e) hardness higher than 95 HV, preferably comprised between 100 and 150 HV; f) elongation at fracture higher than 2%, preferably comprised between 2 and 25%, more preferably between 7 and 20%, measured at room temperature; g) yield strength or upper yield strength higher than 250 MPa up to 200°C and higher than 90MPa up to 300°C; h) ultimate tensile strength higher than 270 MPa up to 200°C and higher than 10OMPa up to 300°C.

Preferably, the 3D-printed product has at least characteristic c), namely yield strength or upper yield strength comprised between 300 and 500 MPa, preferably between 370 and 500 MPa, more preferably between 460 and 500 MPa, measured at room temperature.

Preferably, the 3D-printed product has at least two of the above- mentioned characteristics, preferably at least c) and d).

More preferably, the 3D-printed product has at least three of the above- mentioned characteristics, preferably at least c), d) and f).

More preferably, the 3D-printed product has at least four of the above- mentioned characteristics, preferably at least c), d), f) and g).

More preferably, the 3D-printed product has at least five of the above- mentioned characteristics, preferably at least c), d), f), g) and h).

More preferably, the 3D-printed product has at least six of the above- mentioned characteristics, preferably at least c), d), f), g), h) and e).

More preferably, the 3D-printed product has at least seven of the above- mentioned characteristics, preferably at least c), d), f), g), h), e) and b).

More preferably, the 3D-printed product has all the above-mentioned characteristics.

As far as characteristic c) is concerned, the yield strength or upper yield strength is:

- preferably comprised between 440 and 470 MPa, when the 3D-printed product, obtained after step (i), is not further subjected to steps (iii) and/or (iv), and optionally to step (ii);

- preferably comprised between 470 and 500 MPa, when the 3D-printed product, obtained after step (i), is further subjected to step (iv) and optionally to step (ii) but not to step (iii);

- preferably comprised between 370 and 400 MPa, when the 3D-printed product, obtained after step (i), is further subjected to steps (iii) and (iv) and optionally to step (ii).

As far as characteristic d) is concerned, the ultimate tensile strength is:

- preferably comprised between 440 and 450 MPa, when the 3D-printed product, obtained after step (i), is not further subjected to steps (iii) and/or (iv) and optionally to step (ii);

- preferably comprised between 450 and 460 MPa, when the 3D-printed product, obtained after step (i), is further subjected to step (iv) and optionally to step (ii) but not to step (iii);

- preferably comprised between 450 and 490 MPa, when the 3D-printed product, obtained after step (i), is further subjected to steps (iii) and (iv) and optionally to step (ii).

As far as characteristic f) is concerned, the elongation at fracture is:

- preferably comprised between 6 and 8%, when the 3D-printed product, obtained after step (i), is not further subjected to steps (iii) and/or (iv) and optionally to step (ii);

- preferably comprised between 2 and 6%, when the 3D-printed product, obtained after step (i), is further subjected to step (iv) and optionally to step (ii) but not to step (iii);

- preferably comprised between 8 and 12% or comprised between 12 and 25% (measured at a temperature ranging from 150 °C to 300 °C), when the 3D-printed product, obtained after step (i), is further subjected to steps (iii) and (iv) and optionally to step (ii).

According to an embodiment of the present invention, the 3D-prited product possess a refined crack-free equiaxed grain structure.

According to an embodiment of the present invention, the 3D-printed product is a porous product with a density preferably comprised between 20 and 99 %, preferably between 30% and 80%, more preferably between 40 and 70%.

According to an embodiment of the present invention, the 3D-printed product possesses a functionally graded density comprised between 20 and 99 %, preferably between 30% and 80%, more preferably between 40 and 70%.

Preferably said 3D-printed product is selected from the group consisting of: a filter, a catalyst, a lattice structure, or a combination thereof.

Without wishing to be bound to a specific theory, the Applicant has found that the presence of a relative high amount of Fe and Ni (from 0.5 to 1 .2 wt% and from 0.8 to 1.5 wt.%, respectively) in the Al-Ti-Cu-Mg-B-Ni-Fe-Si powder used to produce the 3D-printed product according to the present invention, advantageously allows the formation of different intermetallic compounds (i.e. second-phase compounds such as Fe-, Ni-, Cu-based phases mentioned above, in particular AI?Cu4Ni and AlgFeNi) that, along with the other compounds (i.e. TiB2 + AhTi phases) contribute to stabilize the microstructure and improve the mechanical properties at high temperatures.

Furthermore, it should be noted that the heterogeneous nucleation of a-AI grains is also enhanced by the Al-Ti-Cu-Mg-B-Ni-Fe-Si powder of the invention because the presence of the Ti- B-based compounds contributes to generate a particularly refined microstructure and to the effective suppression of the solidification cracking phenomenon thereby ensuring processability by AM. Furthermore, as also demonstrated in the Examples, at high temperature, the 3D-printed product comprising the Al-Ti-Cu-Mg-B- Ni-Fe-Si alloy of the present invention, shows higher mechanical properties with respect to the products obtainable (with traditional techniques such as molding or forging) from similar alloys which, however, do not comprise Ti and B and/or comprise lower amounts of Ni and/and Fe.

Without wishing to be bound to any specific theory, the Applicant has found that the presence of the above-mentioned relative high amount of Fe in combination with the above-mentioned relative high amount of Ni, surprisingly allows to enhance the mechanical properties of the 3D-printed product obtained from the Al-Ti-Cu-Mg-B-Ni-Fe-Si powder of the invention, especially at high temperatures (preferably up to 300 °C)

The present invention also refers to a use of the 3D-printed product in high temperature applications (up to 300°C), such as aerospace and automotive/racing applications, preferably for the manufacturing of engine heads, heat exchangers, pumps.

Examples

Example 1 An Al-Ti-Cu-Mg-B-Ni-Fe-Si powder according to the present invention was produced by gas atomization.

The chemical composition of the powder is shown in Table 1 .

Table 1

Example 1.1

Another Al-Ti-Cu-Mg-B-Ni-Fe-Si powder according to the present invention was produced by gas atomization.

The chemical composition of the powder is shown in Table 1.1. With respect to the Al-Ti-Cu-Mg-B-Ni-Fe-Si powder of Example 1 , this powder comprises Ti and B in a lower content.

Table 1.1

Example 1 .2 (comparative)

In order to carry out comparative tests, a powder of the commercial aluminum AISiWMg alloy was employed.

The chemical composition of the AISiWMg alloy powder, is shown in Table

1.2.

Table 1.2

Example 2

A LPBF system was used to manufacture cubic samples (10 mm x 10 mm x 10 mm) using the powders of Examples 1 (hereinafter “sample 1 ”), 1.1 (hereinafter “sample 1.1 ”) and 1.2 (hereinafter “AISiWMg alloy”). Optimization of LPBF parameters was performed by investigating the relative density of 18 specimens, in accordance with full-factorial DOE statistical analysis. The relative density of samples was evaluated using image analysis. The highest density obtained for each sample is shown in Table 2.

Table 2

Furthermore, horizontal and vertical cylindrical samples with diameter of 11 mm and length of 80 mm were produced by LPBF using the best combination of parameters and machined to produce dogbone specimens. Geometry of specimens (gauge length of 30 mm and cross-section diameter of 6 mm) is in accordance with the ASTM E8 Standard Test Methods for Tension Testing of Metallic Material (ASTM Committee on Mechanical Testing, ASTM Int., vol. ASTM Stds., no. Designation: E8/E8M-13a, pp. 1 -28, 2013). Tensile tests were performed using a Zwick Roell Z100 universal testing machine equipped with extensometer. Tensile tests were performed at room temperature and at 150, 200, 250 and 300 °C.

Microstructure analysis was carried out by field emission scanning electron microscope (FE-SEM) mod. Zeiss Sigma 500 equipped with energy dispersive X-ray analysis (EDX) mod. Oxford Instruments Ultim Max and by light optical microscope (LOM). Chemical etching was performed using Keller’s reagent. A heat treatment comprising solution annealing followed by water quenching (called hereinafter “solution treatment-water quenching” - ST-WQ) was carried out at 530 °C; the treatment was carried out for 1 hour. The hardness tests were performed with a load of 300 g and a dwell time of 15 seconds. XRD investigation was carried out using a 0-0 diffractometer. The data were collected with a scan rate of 1 °/min and a step size of 0.02°.

Example 3

Prismatic samples were produced by LPBF using the powder of the present invention having chemical composition as reported in Table 1 , Example 1 and in Table 1 .1 , Example 1.1.

The samples thus produced namely, sample 1 and sample 1.1 , comprise the Al-Ti-Cu-Mg-B-Ni-Fe-Si alloy according to the present invention having the same chemical composition (measured by means of inductively coupled plasma optical emission spectrometry (ICP-OES)) of the Al-Ti-Cu- Mg-B-Ni-Fe-Si powder of Example 1 and Example 1.1 , respectively. Figures 1 a and 1d show a low magnification optical micrograph of the section parallel to the building direction of the “as-built” sample 1 and sample 1.1 , respectively after chemical etching with Keller’s solution. The fish-scale pattern generated by aligned solidified melt pools is noticeable.

A Light Optical Microscopy (LOM) image of the as-built sample taken at higher magnification (reported in this case only for sample 1 ) is shown in Figure 1 b and reveals a fine microstructure, characterized by micro-sized equiaxed grains. The Scanning Electron Microscopy (SEM) micrograph, depicted in Figures 1 c and 1 e, shows the solidification structure of the as- built sample 1 and sample 1.1 respectively with fine a-AI cells, coarse second-phase particles (size in the order of a few microns) and finer particles. The finer particles are located both at the boundaries and at the core regions of the solidification cells.

In Table 3 and Table 3.1 , the results of Energy-dispersive X-ray spectroscopy (EDS) chemical analysis, which was performed on spots labelled as A and B in the micrograph of Figure 1 c and 1 e respectively, are shown. These results reveal that the coarse particles have higher content of Ti and B than the Al cells.

Table 3

Table 3.1

The SEM analysis showed that the as-built samples (both in the case of sample 1 and sample 1.1 ) are characterized by refined cells, typical of rapidly solidified materials, that are surrounded by zones with pronounced solute segregation (Figure 1 c and 1 e, respectively).

Example 4

The XRD diffractograms of the as-built and of the solution treated (ST- WQ) samples of the present invention are reported in Figure 2 (for sample 1 ) and in Figure 3 (for sample 1.1 ). In Figures 2a and 3a, the full diffractogram of the as built samples 1 and 1.1 are respectively shown together with the Miller indexes of the main reflections of Al.

A magnification of the background signal in the low angle 20 region of the as-built and solution treated sample 1 is shown in Figure 2b and 2c, respectively.

A magnification of the background signal in the low angle 20 region of the as-built and solution treated sample 1.1 is shown in Figure 3b and 3c, respectively.

The results show the high intensity of Al peaks with respect to those of the second phases. A high number of peaks are noticeable at low angles, both for the as-built and solution-treated samples. The characteristic peaks of TiB2, Al CiuNi, AhTi and AlgFeNi compounds were detected in both conditions (the XRD diffractograms therefore confirm the existence of TiB2 and AhTi phases in the as-built samples as well). Reflections of 0-AhCu and p-MggSi phases were only identified in the diffractogram of the as-built samples. As expected, the solution treatment performed at 530 °C followed by water quenching was able to dissolve said 0 and [3 phases. Example 5

SEM analysis was performed on samples aged at 180 °C for 3 hours to achieve the maximum hardness, both starting from the as-built and solution treated (solution annealing carried out at 530 °C for 1 hour followed by water quenching) conditions, namely T5 and T6 conditions, respectively.

The analyses are shown only for sample 1 in Figures 4a and 4b. In the T5- treated sample (as-built + aging, Figure 4a), the solidification structure with second phases at cell boundaries and within the cell cores is still visible. On the contrary, solution annealing led to a radical change in microstructure. Indeed, solidification cells are no longer noticeable in the T6-treated sample (solution treatment + aging, Figure 4b) and second particles evidently coarsen.

The direct aging performed from the as-built condition does not modify the morphology of second phases formed on solidification (see Figure 4a). On the contrary, as mentioned above, solution treatment leads to a drastic change in microstructure as shown in Figure 4b. The solute-rich intercellular network made by fine second phases is indeed replaced by coarser particles, which are identified as TiB2, AhTi, AI?Cu4Ni and AlgFeNi by XRD analysis (as reported in Figure 3c).

Example 6

Tensile tests were carried out at room temperature on as-built, T5-treated (as-built + aging) and T6-tretated (solution annealing carried out at 530 °C for 1 hour followed by water quenching + aging) dogbone samples with longitudinal axis parallel and orthogonal to the building direction, which were named as Z- and XY-samples, respectively.

In order to achieve the maximum hardness, the aging was performed at 180 °C for 3 hours for sample 1 and at 160 °C for 3 hours for sample 1.1. Representative tensile curves are shown in Figures 5a (Z) and 5b (XY) for sample 1 , in Figures 5c (XY) and 5d (Z) for sample 1.1 and the upper yield strength (UYS), ultimate tensile strength (UTS) and elongation at fracture values (E%R) at room are reported in Table 4 for sample 1 and in Table 4.1 for sample 1.1. Table 4

Table 4.1 As shown in Figure 6, the curves of sample 1 are characterized by a lower and an upper yield strength, especially pronounced in the as-built and T5 samples records. Significative strain-hardening was noticeable only in the tensile curve of the T6 samples. The alloy in the T5 condition revealed the highest UYS, equal to 495.6 MPa and 478.5 MPa for the XY- and Z- samples, respectively.

The high solidification and cooling rates generated by L-PBF are responsible for the formation of extended solid solution and refined microstructure in the as-built material, leading to high material strength.

The T6-treated alloy showed higher elongation at fracture with respect to the other conditions.

As shown by the results of tensile tests performed along the Z- and XY- direction, the sample exhibits an almost isotropic behavior. The epitaxial and competitive growth of coarse columnar grains is indeed suppressed by heterogeneous nucleation of grains stimulated by the addition of nucleants.

Similarly, as shown in Figure 6, the curves of sample 1.1 are characterized by lower and upper yield points as well as serrations and significant strainhardening. The T6 condition revealed the highest tensile strength values, reaching 466.8 MPa and 468.6 MPa in the XY- and Z-specimens, respectively. The T6 alloy showed higher elongation at fracture (about 14%) with respect to the other conditions.

The Z-samples were also tested in the T6 condition at 150°C, 200°C, 250°C and 300°C for sample 1 and at 150 °C and 250 °C for sample 1.1. Tensile properties of the materials at high temperature are also shown in Figure 6 and summarized in Table 5 and 5.1 , respectively.

A strong drop in yield and tensile strength was noticed above 250 °C for sample 1 .

Table 5

Table 5.1

The Al-Ti-Cu-Mg-B-Ni-Fe-Si alloys printed by LPBF according to the present invention (sample 1 and 1.1 ) shows higher mechanical strength at elevated temperatures both with respect to the wrought AA2618 Al alloy (Al-Cu-Mg-Ni-Fe-Si), typically used for high temperature applications, and with respect to the AISiWMg alloy, typically used in the sector of LPBF. Figure 6 shows the YS vs. temperature (a) and UTS vs. temperature (b) plots of wrought 2618, AISiWMg and of Al-Ti-Cu-Mg-B-Ni-Fe-Si (i.e. Al-Ti- Cu-Mg-B-Ni-Fe-Si alloy according to the present invention, namely sample 1 and sample 1.1 ) produced by LPBF, both in T6 conditions.

Without wishing to be bound to a specific theory, the applicant has found that the Ti, B, Fe, Ni and Al- based compounds that forma at high temperature (such as, for example, TiB2 and AhTi particles) advantageously stabilize the microstructure of the alloy and enhance the mechanical behavior of the Al-Ti-Cu-Mg-B-Ni-Fe-Si alloy at higher temperatures.