New powder, additive manufacturing method for manufacturing a component from the new powder, and article made therefrom

Abstract

The invention provides an Al-based powder suitable for additive manufacturing. In one embodiment, the powder consists of 3-5.5 wt% Mn, 0.2-2 wt% Zr, 0.1-1.4 wt% Cr, 0-2 wt% Mg, up to 0.7 wt% Fe and Si in total, up to 0.7 wt% O as unavoidable impurities, and up to 0.5 wt% other unavoidable impurities, with the balance being Al. The invention also relates to an additive manufacturing method and to an article made by the additive manufacturing method.

Description

Technical Field

[0001] This invention relates to a novel aluminum alloy suitable for use in components made from the alloy by additive manufacturing (AM) methods. Background Technology

[0002] Metal-based additive manufacturing technology is used in many industries, including aerospace and automotive.

[0003] Aluminum alloys are used in a variety of applications. Aluminum alloy products are typically produced by forming casting or forging methods. Forming casting usually involves casting molten aluminum alloy into its final form, such as via high-pressure molds, permanent molds, wet and dry sand, encasing molds, or plaster casting. Forged products are typically produced by casting molten aluminum alloy into ingots or billets. Ingots or billets are usually further hot-worked, and sometimes cold-worked, to produce their final form.

[0004] Depending on the alloy composition and the established methods, aluminum alloys can be classified into several systems. Most Al-based casting alloys contain 9-12% Si, an element that forms a eutectic with excellent casting properties at approximately 12% Si. Other elements that can be used in casting alloys are Mg and Cu. The addition of these elements maintains good casting properties and combines with solid solution strengthening and precipitation strengthening from intermetallic compounds. Typical examples are Al-7Si-0.7Mg and Al9Si-3Cu.

[0005] To achieve higher strength, it is necessary to utilize other alloying elements and reduce the Si content. Forged aluminum alloys with higher strength are called the 2000 series, 6000 series, and 7000 series alloys. Tables 1 and 2 show the chemical composition and mechanical properties of common aluminum alloys.

[0006] Table 1.

[0007]

[0008] Al-Cu alloys (2000 series)

[0009] In this alloy, Mg is a minor additive. Other elements are Mn, Cr, and Zr. The main contribution to strength originates from intermetallic phases (e.g., Al). 20 Cu2Mn3, Al 18Solid solution hardening and precipitation hardening are induced by Mg3Cr2 or AlZr3. To achieve this hardening, a solid solution heat treatment followed by aging is performed. Iron and silicon are considered impurities because they form unfavorable intermetallic compounds. Titanium can be added to control the grain structure during ingot casting. This alloy has less favorable corrosion properties compared to other types of Al alloys. Parts made from this alloy have a high strength-to-weight ratio and can be used at temperatures up to 150°C. This alloy is generally not weldable but has good machinability.

[0010] Al-Mg-Si alloys (6000 series)

[0011] This alloy is an important component of high-strength Al alloys. Mg and Si are typically added in a specific ratio as Mg₂Si, the main precipitate. This alloy requires solution treatment and aging, but will not reach the strength levels of 2000 and 7000 series alloys. However, it is weldable and exhibits good formability, extrudability, machinability, and corrosion resistance.

[0012] Al-Zn alloy (7000 series)

[0013] This alloy composition features Zn as the primary alloying element, with the addition of Mg, Cu, and Cr. These elements form an intermetallic phase, which is used in solution treatment and aging to achieve extremely high strength. Resistance to stress corrosion cracking is reduced, and these materials are typically used under over-aging conditions to achieve an optimal combination of strength, fracture toughness, and corrosion resistance.

[0014] The mechanical properties of the alloys in different groups are summarized in Table 2.

[0015] Table 2.

[0016]

[0017] Corrosion is a systemic characteristic because it depends on the material and the environment in which it operates. Nevertheless, some general predictions can be made about the behavior of alloys. Aluminum typically forms a passivating oxide layer that minimizes the overall corrosion rate. If this layer breaks down due to localized changes in the composition, such as in the form of precipitates or interfaces between dissimilar metals, corrosion can occur in the form of crevice corrosion, pitting corrosion, or electrolytic corrosion. These are all extremely rapid processes that can lead to severe failure. When the alloying elements are in solid solution form within an fcc matrix, aluminum alloys generally exhibit strong resistance to all forms of localized corrosion, while alloys with Cu or Zn precipitates perform poorly in this regard.

[0018] In laser powder bed fusion (LPBF), a laser beam scans across a powder bed, melting powder particles in a predetermined pattern. New powder layers are then applied, and the process is repeated. To produce fully dense parts, it is necessary not only to melt the powder particles of the new layer but also to melt some of the previously formed layers. This ensures complete fusion between layers and removes any defects, such as porosity. The laser scans extremely quickly (200–7000 mm / s) and has a small spot size (40–100 µm); this means that the resulting melt pool is small, and the material remains in a liquid state for a very short time (approximately 0.2 ms). These manufacturing limits are designed to ensure that any segregation of alloying elements during solidification is minimized due to the short available interaction time.

[0019] However, the directionality of this method results in both the grains and subgrain structures oriented towards the heat source. The subgrain morphology is dendritic or cellular. This structure is typically large columnar grains with extremely fine cellular structures in the building direction. Most of the heat extracted from the built component is dissipated through this component to the building plate, which acts as a heat sink, thereby enhancing the directionality of the solidification process.

[0020] The conventional alloys described above are not always well-suited for LPBF (Limited Base Fluid) processes. Typically, only cast alloys can be used without problems in LPBF. A360.0 is an alloy well-suited for printing, with very few defects and a reasonable operating window suitable for industrial operations. However, this material offers only limited strength.

[0021] Alloys that achieve sufficient strength using conventional manufacturing methods are more difficult to process using the LPBF method without serious defects. In particular, the solidification behavior is unfavorable, and cracks can form during solidification.

[0022] Several inventors have partially solved these problems:

[0023] EP0105595A2 (Alcan) discloses an aluminum alloy composed of 1.5-7.0 wt% Cr, 0.5-2.5 wt% Zr, 0.25-4.0 wt% Mn, with the balance being aluminum and general impurities.

[0024] In another embodiment, the composition ranges as follows: 3.0-5.5 wt% Cr, 1.0-2.0 wt% Zr, 0.8-2.0 wt% Mn, with the balance being Al and general impurities.

[0025] The inventors also disclosed a method for producing powder by rapidly solidifying a molten alloy at a cooling rate of at least 1000°C / second, which produces soft particles in which most of the alloy additives are retained in a solid solution.

[0026] The resulting particles are solidified at temperatures of 300–500°C to achieve age hardening. In one example, powder of alloy A (Cr 5.25, Zr 1.75, Mn 1.75, with the remainder being Al) was produced by gas atomization, followed by sheet rolling at 350°C, resulting in a mixed microstructure with regions containing and without visible precipitates. This material exhibits a tensile strength of 588 MPa, a yield strength of 530 MPa, and an elongation at break of 6%.

[0027] WO2020 / 139427 (HRL Laboratories) discloses additively manufactured high-temperature aluminum alloys and raw materials for preparing the alloys. The disclosed raw materials consist of a powder mixture comprising 80-99 wt% aluminum-containing matrix powder and 1-20 wt% alloy powder with a particle size smaller than the aluminum-containing matrix powder. In some embodiments, the raw materials contain other alloying elements.

[0028] One embodiment discloses a gas-atomized matrix powder containing, by weight percent: Al 92.6, Cu 6.7, Mn 0.35, Ti 0.24. The matrix powder has a particle size distribution suitable for selective laser melting: D10 = 15 μm, D50 = 27 μm, and D90 = 44 μm. The matrix powder is mixed with zirconium powder having an average particle size of 0.5-1.5 μm and a content of 2% by weight.

[0029] WO2018 / 009359A1 (NanoAl) discloses a series of aluminum alloys with high strength and ductility, excellent corrosion resistance and weldability.

[0030] The aluminum alloy disclosed in this document contains approximately 1-10% by weight of Mg, 0.45-3% by weight of Zr, and the balance being Al. This alloy contains absolutely no intentionally added scandium, erbium, thulium, ytterbium, or lutetium.

[0031] In another embodiment, an aluminum alloy is disclosed comprising Mg and at least one element from Group 4 elements Ti, Zr, and Hf, Group 5B elements V, Nb, Ta, Group 6B elements Cr, Mo, and W, with the remainder being Al. The alloy is completely free of intentionally added scandium, erbium, thulium, ytterbium, or lutetium, wherein the alloy comprises nanoscale aluminum-transition metal precipitates in an aluminum matrix with an average diameter ranging from about 3 nm to about 50 nm, and wherein the transition metal is selected from the group consisting of Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, and W.

[0032] Scalmalloy / Al-Mg-Sc alloys for laser powder bed fusion have recently been developed.

[0033] EP3165620A1 and EP3181711A1 (Airbus Defense and Space GmbH) teach how to achieve high strength, elongation, and hardness through laser powder bed fusion processing. The produced parts are defect-free and exhibit high corrosion resistance. This invention is based on the use of scandium, an expensive and rare rare earth element.

[0034] Additive manufacturing (AM) methods for forming objects from powdered aluminum alloys are currently limited by the lack of suitable alloy powders for producing high-strength objects. Most available high-strength Al alloy compositions have been developed for casting and for AM in powder form, such as through laser powder bed fusion processes, but the resulting objects are prone to cracking. Some recently developed aluminum alloys for laser powder bed fusion processes are thought to be able to produce crack-free objects.

[0035] Therefore, there is a need for aluminum powder alloys that can be manufactured from a homogeneous melt, for example, through atomization, without forming excessive deposits that would interfere with the manufacturing process. Furthermore, such powders should be suitable for forming crack-free objects with high strength immediately after printing and after heat treatment via AM methods, such as LPBF.

[0036] One object of the present invention is to provide an aluminum powder suitable for additive manufacturing of crack-free articles, particularly suitable for aluminum powder used in laser powder bed fusion, i.e., LPBF. The alloy composition will be able to be produced in the form of a fully pre-alloyed powder, thereby avoiding segregation of particles and inhomogeneous powder compositions, and will be free of expensive alloying rare earth elements and copper. In particular, the alloy of the present invention should be completely free of intentionally added rare earth metals such as scandium, erbium, thulium, ytterbium, or lutetium, and free of intentionally added Cu. This object is satisfied in the first aspect of the invention.

[0037] Another objective of the present invention is to provide an article produced from the powder that is crack-free and has high strength both immediately after printing and after heat treatment. This objective is achieved in a second aspect of the invention.

[0038] Another object of the present invention is to provide a method for producing articles from the powder via additive manufacturing. The provided method is robust and has a wide operating window, meaning that variations in operating parameters do not affect the properties of the printed object. This object is satisfied in a third aspect of the invention.

[0039] Furthermore, because chromium is an environmental pollutant, high chromium content is generally undesirable, especially in consumer products. Therefore, one aspect of the present invention is to provide an aluminum-based alloy with low chromium content that has or provides the aforementioned desired properties when used.

[0040] The above-mentioned objectives are achieved by aspects and embodiments thereof as shown in the appended claims. Aspects of the invention are summarized under the heading of the following Invention Summary. Summary of the Invention

[0041] In the first aspect and embodiments, this document describes in detail a pre-alloyed Al-based powder suitable for additive manufacturing, which comprises, by weight of the total pre-alloyed Al-based powder, the following:

[0042] 3-5.5% by weight of Mn,

[0043] 0.2-2% by weight of Zr,

[0044] 0.2-1.4% by weight of Cr,

[0045] 0-2% by weight of Mg,

[0046] The total amount of Fe and Si is at most 0.7% by weight.

[0047] At most 0.7% by weight of O as an unavoidable impurity, and

[0048] Other unavoidable impurities, up to 0.5% by weight.

[0049] The balance is Al.

[0050] In one embodiment of this, a pre-alloyed Al-based powder is used, wherein the Zr content is 0.3-1.8 by weight.

[0051] In one embodiment of this, a pre-alloyed Al-based powder is used, wherein the Cr content is 0.3-1.4% by weight.

[0052] In one embodiment of this, a pre-alloyed Al-based powder is used, wherein the total content of Fe and Si is at most 0.2 wt%.

[0053] In one embodiment of this, a pre-alloyed Al-based powder is used, wherein the pre-alloyed Al-based powder does not contain intentionally added Mg.

[0054] In one embodiment of this, a pre-alloyed Al-based powder has a solidification cracking susceptibility index S below 20°C, calculated according to equation (1).

[0055] In one embodiment of this, a pre-alloyed Al-based powder has a particle size range of 10–53 µm as measured by laser diffraction according to ISO 13320-1 1999, wherein less than 10% of the particles fall outside this specified particle size range.

[0056] In the second aspect, this document describes in detail a method for producing articles by additive manufacturing, which includes the following steps:

[0057] - Provide a pre-alloyed Al-based powder according to any one of the embodiments of the first aspect;

[0058] - Deposit a layer of this pre-alloyed Al-based powder, thereby forming a powder bed;

[0059] - The pre-alloyed Al-based powder is heated according to a predetermined pattern by laser or electron beam to fuse or melt the particles to be combined;

[0060] - Cooling fused or molten particles;

[0061] - Repeat the deposition, heating, and cooling steps until an additively manufactured article is formed;

[0062] - Recycle the manufactured products.

[0063] In one embodiment of this, the manufactured articles are subsequently subjected to heat treatment.

[0064] In one embodiment of this, the heat treatment includes heating the produced article in an air atmosphere to a temperature of 150-450°C for a period of at least 0.5 hours.

[0065] In the third aspect, this document describes in detail an article formed by a melt fusion method of pre-alloyed Al powder according to any one of the embodiments of the first aspect.

[0066] In one embodiment of this aspect, an article is formed by a hot-melting method of pre-alloyed Al powder, wherein the melt fusion method is the method according to any one of the embodiments of the second aspect. Attached Figure Description

[0067] Figure 1 :

[0068] Solidification was simulated using the Scheil-Gulliver method. These lines show how the temperature at equilibrium between the residual liquid (solid phase line) and the solid varies with the mass fraction of the solid formed.

[0069] Figure 2 :

[0070] Operating window diagram for Alloy D. Relative density is displayed as a contour plot, where the vertical and horizontal axes represent laser power (W) and laser velocity (mm / s), respectively. The hatch distance is 100 µm. The operating window represents the region with a relative density higher than 98.5%.

[0071] Figure 3 :

[0072] Hardness changes with heat treatment. Aging treatment was performed in air at 410°C. The time corresponding to temperature is displayed on the horizontal axis.

[0073] Figure 4 :

[0074] A study on the effect of the combination of Fe and Si on the crack susceptibility of the alloy of this invention.

[0075] Figure 5 :

[0076] (a) Particle size (µm) as a percentage of volume, and

[0077] (b) Particle size (µm) relative to cumulative volume.

[0078] Figure 6 :

[0079] The DOE cube design shows all three planes and notches along the front plane for marking the airflow direction (front plane: XZ; top plane: XY; right plane: YZ; as well as the construction direction and airflow direction).

[0080] Figure 7 :

[0081] The design of experiments (DOE1) shows different processing parameters and the number of samples for each parameter. Layer thickness and laser power are kept constant.

[0082] Figure 8 :

[0083] The results of Experimental Design 6.1 show surface plots of the relative densities of all four alloys. a) Alloy A, b) Alloy B, c) Alloy C, d) Alloy D.

[0084] Figure 9 :

[0085] The surface plots of a) Experimental Design 6.2, b) Experimental Design 6.3, and c) / d) Experimental Design (1,2,3) are shown, respectively, showing the laser velocity of alloy D relative to the scanning distance and the laser velocity relative to the laser power.

[0086] Figure 10 :

[0087] a) X-ray images of all four alloys in their freshly printed condition.

[0088] b) A graph of alloy D in its freshly printed condition, with the inset showing the region between 2θ = 37° and 47°, indicating the formation of a minor phase in all four alloys. This is presumed to be an Al-Mn based precipitate. Δ = pure Al peak, and =Al-Mn precipitate peak.

[0089] Figure 11 :

[0090] a) Theoretical effects of Mn, Cr, and Zr on lattice parameters. Note that the dashed line represents intercept correction, assuming pure Al = 4.0478 Å instead of Al = 4.016 Å.

[0091] b) Lattice parameter values ​​based on X-ray diffraction patterns for all four alloys in their freshly printed condition. These values ​​are reasonable and within an accuracy range of 0.01 Å.

[0092] Figure 12 :

[0093] a) A 20X optical microscope image of powdered alloy C in a slightly atomized state, etched with Keller's reagent.

[0094] b) 50X optical microscope image of powder alloy C in a slightly atomized state, etched with Keller's reagent.

[0095] Figure 13 :

[0096] A top (XY), right (YZ), and front (XZ) view of an alloy D cube in a freshly printed state with a relative density of approximately 99.6%.

[0097] Figure 14 :

[0098] Two types of typical defects were observed after printing Al alloy (alloy D is shown in the figure); a) keyhole-shaped defects shown in low magnification (a) and close-up view (b); no fusion defects were observed in low magnification and close-up view (d).

[0099] Figure 15 :

[0100] SEM image of the powder of alloy C in a slightly atomized state, showing a dendritic structure (a); magnified image (b) showing nanoparticles within the dendritic region (marked with yellow arrows). Particles marked with red arrows are from residual silica polished with OP-U.

[0101] Figure 16 :

[0102] SEM images of alloy D (in printed form) along the following planes: a) front plane, where the construction direction is from bottom to top and the melt pool is marked in red; b) top plane, where the construction direction is from the image inward.

[0103] Figure 17 :

[0104] SEM images of the precipitates at the melt pool boundary are displayed at high resolution. Two categories of precipitates are identified: large precipitates (200–300 nm, blue arrows) and small precipitates (<100 nm, yellow arrows). Large precipitates are shown at a higher magnification in the inset in b).

[0105] Figure 18 :

[0106] SEM images of alloy D show a) a melt pool boundary with grain refinement in the melt pool boundary region, and b) magnified regions showing grain boundaries and subgrain boundary particles.

[0107] Figure 19 :

[0108] Chemical compositions of three categories of Mn-containing precipitates were characterized by EDS analysis; these included subgrain boundary precipitates and small and large precipitates at the melt pool boundary. Subgrain boundary precipitates, such as... Figure 18 As shown in the figure, and the precipitate at the boundary of the melt pool is as follows Figure 17 As shown in the image.

[0109] Figure 20 :

[0110] SEM images of alloy D show two categories of melt pools: a) a normal melt pool with long columnar grains epitaxially grown along the construction direction; b) a melt pool with refined grains (grain refinement shown in the inset image).

[0111] Figure 21 :

[0112] a) SEM image in SE2 mode, showing the top planar cross section of alloy D with Zr dendritic structure at 20kX.

[0113] b) SEM image in BSE mode, showing the top planar section of alloy C with Zr facets at 2.5 kX. Detailed Implementation

[0114] Integrated Computational Materials Engineering (ICME) is highly effective in optimizing computational methods for rapid solidification of alloys.

[0115] It can calculate the effects of different alloying elements on solidification behavior and precipitation processes in alloys, and can obtain good predictions of the properties of new alloy systems.

[0116] This method is based on a deep understanding of the relationship between phase relations and properties of specific target materials and procedures. It requires combining computation with selected experimental studies.

[0117] The advantage of the ICME development process is that it allows the study of previously unknown regions of the compositional space of a class of materials. This method is used in this invention to illustrate why alloy compositions differ significantly from available alloys, and how alloying elements can be used to obtain high-strength components.

[0118] As described by first-principle calculations, alloying elements are selected based on the potential solid solution strengthening effect.

[0119] This was then combined with previous knowledge about the expansion of solid solubility in Al-based alloys to estimate how much solute would be expected to dissolve and how strong the alloy would be when it was first printed.

[0120] The amount of supersaturation in the solid solution was confirmed by nonequilibrium calculations and by an assessment of the nucleation / coarsening resistance of second-phase particles based on the diffusion rate during LPBF processing and rapid cooling.

[0121] Provisional phase diagrams (from binary to quaternary) were generated using ThermoCalc software (version and database), and supercooling that may be caused by processing via LPBF was used as a factor to determine how high the percentage of soluble solute is compared to equilibrium.

[0122] The second important principle for alloy design is to select alloying elements based on Scheil solidification calculations and to make selections with negligible susceptibility to solidification cracking.

[0123] This method establishes a simplified path to avoid unnecessary problems during printing. The Scheil curve is generated using ThermoCalc software with the TCAL7 and MOBAL5 databases.

[0124] Unexpectedly, no low-melting-point phase was formed near the end of the solidification stage of the alloy of the present invention, thereby avoiding the risk of solidification cracking. This is indicated by the excellent cracking solidification index.

[0125] Furthermore, it was previously anticipated that by carefully selecting the elements and their contents, a Cr-containing phase would form during solidification at the time of printing, thereby reducing the solid solution hardening effect of Cr and unexpectedly potentially preventing the formation of such precipitates. Similarly, alloying elements Zr and Mn in the contents according to the invention do not form undesirable precipitates at the time of printing, thus enabling age hardening.

[0126] The pre-alloyed Al-based powder is designed using integrated computational materials engineering (ICME) for additive manufacturing methods, where solidification is rapid, exceeding 1000°C / s, and the formation of precipitates in the liquid is suppressed.

[0127] Alloying elements are responsible for strengthening the alloy and are effective only in rapidly solidifying microstructures. In all embodiments of the invention, the alloy has a stable solidification path until the solid mole fraction reaches 80%, and the drop in solidification temperature is less than 25°C, thereby allowing LPBF printing to proceed substantially without crack formation (see...). Figure 4 This is in contrast to alloys in existing technologies (see...). Figure 1 However, A360.0 is an exception. Furthermore, the most preferred alloy of this invention has a stable solidification path, i.e., it forms a single structure upon solidification and can be printed, for example, by LPBF without cracking (see...). Figure 4 ).

[0128] The printed product is characterized by a structure with a cellular / dendritic microstructure, wherein the phase composition of the microstructure is a single-phase fcc containing less than 5% by volume of Al6Mn, Al3Zr and other precipitates.

[0129] The pre-alloyed Al alloy powder according to the present invention can be used for additive manufacturing of Al alloy parts. These parts can be easily printed with a wide operating window. The parts exhibit high strength and can be heat-treated after the additive manufacturing step to further improve their properties.

[0130] Therefore, to obtain improved strength, the printed articles can be heat-treated (i.e., aged) in air at 150-450°C for at least 0.5 hours. This allows the production of articles with a hardness of at least 125 HV0.3 as measured according to ISO 6507-1-2018.

[0131] The microstructure of the product is characterized by a cellular / dendritic microstructure, wherein the phase composition of the microstructure is a single-phase fcc containing 2-15 vol% Al6Mn and Al3Zr, and up to 10 vol% other precipitates.

[0132] The alloy system is based on Al-Mn with added Cr and Zr. The results obtained from experiments and calculations are presented in detail in Table 3, where each component in the aluminum alloy is indicated by "X", which contributes to the alloy synergistic effect of the aluminum alloy of the present invention when combined with other elements.

[0133] Table 3 highlights the compositional regions of particular interest, where the elements in the alloys of the present invention exhibit the best synergistic effect. However, although working in the concentration regions surrounding these regions of particular interest is less preferred, alloys with such non-optimal elemental inclusions still exhibit better improvement properties than other alloys in the prior art and are therefore not excluded from the present invention.

[0134] According to the present invention, the following is described in detail herein:

[0135] A pre-alloyed Al-based powder suitable for additive manufacturing, comprising, by weight of the total pre-alloyed Al-based powder:

[0136] 3-5.5% by weight of Mn,

[0137] 0.2-2% by weight of Zr,

[0138] 0.2-1.4% by weight of Cr,

[0139] 0-2% by weight of Mg,

[0140] The total amount of Fe and Si is at most 0.7% by weight.

[0141] At most 0.7% by weight of O as an unavoidable impurity, and

[0142] Other unavoidable impurities, up to 0.5% by weight.

[0143] The balance is Al.

[0144] Oxygen is commonly present in aluminum alloys as an unavoidable impurity. The method of the present invention for producing the pre-alloyed powder provides oxygen as an unavoidable impurity at a concentration between 0.35 and 0.55% by weight (see Table 9), but this value will vary depending on the powder size. No particular effects or harmful effects of oxygen were observed in the experiments.

[0145] Generally, it is preferable to ensure that the content level of the other unavoidable impurities is as low as technically possible. Preferably, the concentration of the other unavoidable impurities is less than 0.4% by weight, less than 0.3% by weight, or even more preferably less than 0.2% by weight or less than 0.1% by weight, based on the weight of the pre-alloyed Al-based powder.

[0146] According to the present invention, alloying elements Cr and Mn are provided to induce a solid solution hardening effect, while alloying element Zr is provided to contribute to precipitation hardening during subsequent heat treatment after the printing process. Surprisingly, as observed in the experiments and indicated in Table 3, a synergistic effect of the added components within the claimed concentration range was observed.

[0147] The Zr content is 0.1-2.0% by weight of the pre-alloyed Al-based powder. In experiments (see Table 3), it was found that a Zr content higher than 2% would interfere with the atomization process, while a content lower than 0.1% by weight would not participate in the desired precipitation hardening effect after the product is heat-treated.

[0148] Preferably, the zirconium content is not less than 0.2% by weight, not less than 0.3% by weight, not less than 0.4% by weight, more preferably not less than 0.5% by weight, or even more preferably not less than 0.6% by weight, based on the weight of the pre-alloyed Al-based powder and in articles produced from Al-based powder.

[0149] Preferably, the zirconium content is less than 2% by weight, less than 1.9% by weight, less than 1.8% by weight, less than 1.7% by weight, more preferably less than 1.6% by weight, less than 1.5% by weight, or less than 1.25% by weight, based on the weight of the pre-alloyed Al-based powder and in articles produced from Al-based powder.

[0150] In a preferred embodiment, the zirconium content is 0.2-1.9 wt% based on the weight of the pre-alloyed Al-based powder and in the articles produced from the Al-based powder, preferably 0.3-1.8 wt%, 0.4-1.7 wt%, more preferably 0.5-1.6 wt%, 0.5-1.5 wt%, but most preferably 0.6-1.4 wt%, 0.7-1.3 wt%, or 0.8-1.1 wt%.

[0151] According to the present invention, the addition of Mn and Cr is believed to increase strength because Al atoms are replaced by Mn and Cr atoms. If the matrix is ​​supersaturated with Mn and Cr, precipitates will form in the fcc-Al matrix, which will lead to a reduction in the substitution strengthening effect.

[0152] Surprisingly, it was found (see Table 3) that, for certain concentrations of Cr and Mn, the presence of both Cr and Mn in the alloys of the present invention resulted in significantly less crack formation in articles formed by LPBF printing than expected. It is generally believed that the more Mn and Cr added, the stronger this effect; however, it was surprisingly found that this effect is maximized at approximately 0.7% wt% Cr and disappears when the chromium content in the alloys of the present invention exceeds 1.4% wt%.

[0153] Subsequent experiments revealed that adding more than 1.4% Cr to the composition of the present invention would result in the formation of unwanted intermetallic Al deposits in the melt at temperatures above 800°C. x Cr y Sediment.

[0154] Such precipitates formed in the melt will be too large to contribute to the strengthening of the alloy, and the removal of Cr from the solid solution reduces the substitution strengthening effect. Furthermore, the formation of intermetallic phases at high temperatures can cause problems during atomization, clogging the nozzle.

[0155] The Cr content is 0.1-1.4% by weight of the pre-alloyed Al-based powder. Preferably, the chromium content is not less than 0.2% by weight, not less than 0.3% by weight, not less than 0.4% by weight, more preferably not less than 0.5% by weight, or even more preferably not less than 0.6% by weight, based on the weight of the pre-alloyed Al-based powder and in articles produced from Al-based powder.

[0156] Preferably, the chromium content is less than 1.4 wt%, less than 1.35 wt%, less than 1.30 wt%, less than 1.25 wt%, more preferably less than 1.20 wt%, less than 1.15 wt%, less than 1.10 wt%, less than 1.05 wt%, or less than 1.0 wt%, based on the weight of the pre-alloyed Al-based powder and in articles produced from Al-based powder.

[0157] In a preferred embodiment, the chromium content, based on the weight of the pre-alloyed Al-based powder and in the articles produced from the Al-based powder, is 0.2-1.4 wt%, preferably 0.3-1.35 wt%, 0.4-1.30 wt%, more preferably 0.5-1.25 wt%, 0.5-1.15 wt%, or 0.6-1.4 wt%, or 0.7-1.3 wt%, or 0.8-1.1 wt%.

[0158] Therefore, the amount of Cr present in the pre-alloyed Al-based powder and in the articles produced from the powder is 0.1-1.4 wt%, preferably 0.2-1.4 wt%, preferably 0.3-1.4 wt%, more preferably 0.4-1.4 wt%, but even more preferably 0.5-1.4 wt%.

[0159] For the same reason, the amount of Mn present in the pre-alloyed Al-based powder and in the articles produced from the powder is 3-5.5% by weight.

[0160] Generally, the alloys of the present invention are observed to respond to the contents of Zr and Cr within this concentration range of Mn, but not outside this range. The observed specific dependence of Zr and Cr on Mn makes it possible that, in the preferred embodiment, the total concentration of Mn+Zr+Cr should be 3.4-8.5% by weight, based on the weight of the pre-alloyed Al-based powder and in the articles produced from the Al-based powder.

[0161] In its preferred embodiment, the total concentration of Mn+Zr+Cr, based on the weight of the pre-alloyed Al-based powder and in the articles produced from the Al-based powder, should be 3.7-8.0 wt%, 4-7.5 wt%, 4.25-7.25 wt%, or 4.5-7 wt%.

[0162] Optionally, the amount of Mg present in the pre-alloyed Al-based powder and in the articles produced from the powder may be up to 2% by weight. Mg will contribute to the strengthening of the solid solution. In one embodiment, the pre-alloyed Al-based powder does not contain intentionally added Mg, and therefore Mg is considered another unavoidable impurity in this embodiment.

[0163] The presence of excessive Fe and / or Si will lead to the formation of a low-melting-point liquid in the final stage of the solidification process. This liquid can be trapped at grain boundaries, which in turn will cause cracks to form during final shrinkage in solidification. Surprisingly, (see...) Figure 4 This phenomenon is not related to the individual contributions of Fe and Si, but depends solely on the total concentration of these two elements, and can be suppressed by limiting the total amount of Fe and Si.

[0164] An important feature has been shown is that the contents of Fe and Si are kept at low levels, at most 0.7% by weight, preferably at most 0.5% by weight, and more preferably at most 0.2% by weight.

[0165] The amount of other unavoidable impurities in the Al-based powder and the articles produced from the powder is at most 0.5% by weight, preferably at most 0.3% by weight.

[0166] The Al-based powder according to the invention has a melting range of up to 300°C, defined as the difference between the initial solidification temperature and the temperature at which all liquids have solidified.

[0167] However, to avoid cracking in the final stages of solidification, the decrease in melting temperature near complete solidification must be limited, see [link to relevant documentation]. Figure 1 A practical way to assess the tendency of an alloy to crack during solidification is to use a solidification crack susceptibility index.

[0168] The solidification crack susceptibility index is defined according to the following equation (1):

[0169]

[0170] Where S is the susceptibility index to solidification cracking, T is the temperature, and f s It is the mass fraction of the solid.

[0171] The Al-based powder according to the invention has a solidification crack susceptibility index below 20°C. This index is a measure of how well the final melt solidifies within a range of 80-100% solidification. The index measures the slope of temperature change within this range. A360.0 has a solidification crack susceptibility index of 44°C, and Al 6160 has an index of 257°C. The lower the index, the less susceptible the alloy is to solidification cracking.

[0172] Al alloy powder can be produced by any rapid solidification method, such as gas atomization, as well as jet deposition, melt spinning, melt extraction, etc., without forming excessive precipitates, which would otherwise be detrimental to production operations because precipitates tend to clog atomizing nozzles or otherwise interfere with the molten flow.

[0173] Powders with a predetermined composition are preferably produced by melting a material having the desired composition and atomizing the melt with an inert gas to obtain pre-alloyed atomized powder. Air or water can also be used as the atomizing medium for atomization, but for the purposes of this invention, the use of water would have a negative impact on powder production because the oxygen content would be too high, and the particle morphology may also be affected because the shape of powders atomized by air and water is often not as regular as that of powders formed by atomization with an inert gas.

[0174] The alloying elements in the powder of this invention are also selected to avoid clogging of the atomizing nozzle during atomization. This problem is avoided in some prior art powder compositions by using mixtures of different powders, which are themselves pure elemental powders or pre-alloyed to obtain the desired chemical composition. The disadvantage of mixtures is that the different powders can separate during processing and finishing, causing chemical inhomogeneity in the printed parts. This problem is avoided in this invention because the aluminum powder of this invention is easily pre-alloyed without loss of composition.

[0175] As is well known in the LPBF printing field, the size and morphology of powder particles used in printing are important for the ability to spread the powder into a uniform and homogeneous powder layer. Working sieves are known to separate particles in the range of 5–150 µm.

[0176] Smaller intervals are beneficial for the melting process and the homogeneity of the powder bed, resulting in a more stable melting process and fewer defects, such as porosity.

[0177] Common particle sizes range from 15-45 µm or 20-53 µm, but depending on the LPBF equipment used and the application requirements, any of the following particle sizes can be separated using sieves: 5-36 µm, 10-45 µm, 15-45 µm, 20-53 µm, 20-63 µm, 45-90 µm, 45-106 µm, 53-106 µm, 45-150 µm, 53-150 µm, 63-150 µm, 75-150 µm, 90-150 µm, 106-150 µm. A particular advantage of the Al-based powder of this invention is that such powders can be readily manufactured within the size range required for LPBF printing.

[0178] The particle size range mentioned means that up to 2% by weight of the powder has a particle size above the upper limit and up to 2% by weight of the powder has a particle size below the lower limit. Particle size is measured by laser diffraction according to ISO 13320-1 1999.

[0179] The shape (morphology) of the powder is also important in defining its spreading behavior. Spherical shapes produce more stable flow and spreading, resulting in fewer defects such as porosity and surface quality. Therefore, gas atomization is preferred when producing novel Al-based powders because the metal powders produced by this method exhibit a spherical shape.

[0180] The Al-based powder according to the invention can be used in a variety of additive manufacturing methods, provided the solidification rate is sufficient. AM by electron beam melting and direct energy deposition is a method that meets these requirements, and this method is illustrated below using a laser powder bed fusion method.

[0181] Example

[0182] Example 1

[0183] To design the composition of Al-based powders, an ICME (Integrated Computational Materials Engineering) strategy was employed. Phase diagram calculations (CALculation of Phase Diagrams; CALPHAD) were used to optimize the alloy composition of the new alloy.

[0184] CALPHAD calculations involve a complete thermodynamic description of all possible phases within the compositional space of Al alloys. Using this foundation, the expected phases and their compositions are predicted with good accuracy. The results are valid for equilibrium states.

[0185] In materials produced via LPBF, the microstructure and existing phases are a result of rapid solidification. Equilibrium calculations are irrelevant to the interpretation of these microstructures. However, thermodynamic and phase relationships in compositional space can be used to assess the tendency for phase formation by incorporating a kinetic description of phase formation. In this way, diffusion processes, nucleation, and growth are taken into account along with the thermodynamic description.

[0186] Using CALPHAD-type calculations, the solidification path and the tendency for alloy element segregation in the final melt region are predicted. For crack-free solidified structures, this segregation should be limited.

[0187] In the first step, calculations are used to determine a provisional composition to obtain an alloy with a stable solidification path, thereby avoiding thermal cracking.

[0188] Next, the compositional space of the alloy is calculated and the maximum limits for alloying elements are determined. Compositional limits are determined by calculating the phases formed during rapid solidification. Outside the limits marked in Table 3, undesirable phases will form, leading to problems with the atomization process or deterioration of mechanical properties.

[0189] By combining phase equilibrium calculations with a kinetic description of the alloy system, the nucleation, growth, and coarsening of the secondary hardening phase are simulated. This method allows for the adjustment of the alloy composition to optimize the precipitation of the secondary hardening phase during heat treatment.

[0190] The CALPHAD calculations are combined with models describing different strengthening mechanisms, including solid solution strengthening, precipitation strengthening, and grain refinement strengthening.

[0191] The primary strengthening mechanism of this alloy was found to be solid solution strengthening. Precipitation hardening and grain refinement strengthening were also identified as contributing to the alloy's strength. The combination of these strengthening mechanisms determines the properties of the new alloy.

[0192] Table 3 below shows the results of calculations performed to determine the range of alloying elements that would form the unwanted phase. To achieve solid solution strengthening, the highest possible amounts of alloying elements that would not form the unwanted phase were identified.

[0193] Compositions marked with "x" in Table 3 are acceptable in terms of the presence of undesirable phases and alloying effects, while compositions without any "x" are unacceptable.

[0194] The formation of Cr-containing precipitates in the melt or solid phase reduces the contribution of solid solution hardening because the precipitate removes Cr from the solid solution. Precipitates formed in the melt at high temperatures are coarse and do not contribute to precipitation hardening.

[0195] The formation of Mn-containing precipitates in the melt or solid phase will reduce the contribution of solid solution hardening. Large major precipitates do not contribute to precipitation hardening. An excessive proportion of major precipitates will lead to clogging of atomizing nozzles and production interruptions.

[0196] Table 3

[0197]

[0198] Example 2

[0199] Four different gas-atomized Al-based powders, A, B, C, and D, with chemical compositions as shown in Table 4, were generated using the calculated alloy compositions. Several parts were printed using different printing parameters. In particular, the laser power and laser speed were varied over a wide range.

[0200] A large area was found in which the relative density of the resulting product was higher than 98.5%. See [link / reference]. Figure 2 .

[0201] Further examination using optical microscopy revealed a single-phase structure without any cracks. According to ISO 6507-1-2018, the Vickers hardness HV0.3 was measured in both the freshly printed and heat-treated states, with the following results: Figure 3 As shown, alloys A, C, and D are age-hardenable; in particular, the hardness of alloys C and D is significantly increased to HV0.3, which is slightly below 130.

[0202] Table 4

[0203]

[0204] The role of chromium

[0205] Experiments suggest that while chromium primarily strengthens solid solutions, it also significantly and surprisingly influences the precipitation hardening sequence. This results in alloys with improved properties and temperature resistance.

[0206] The freshly printed sample showed no precipitation (as verified experimentally), meaning that all precipitation occurred during the hardening heat treatment. Therefore, the order of precipitation is influenced by composition.

[0207] Precipitation hardening heat treatment at 678 K involves the sequential precipitation of Al6(Mn,Cr) and Al. 12 The alloy is hardened using (Mn,Cr) and Al3Zr. The amount of chromium in the alloy will determine the precipitation sequence and the amount of precipitate and its stability to the coarsening. This means that the properties depend primarily on the specified amount of Cr as detailed in this paper.

[0208] Compared to similar materials without chromium, the chromium-containing alloys of this invention have the advantages of higher strength, and importantly and unexpectedly, increased temperature stability, which is a key parameter for industrial applications.

[0209] The role of Cr Figure 3 This has been confirmed. The hardening curve is shown below:

[0210] Alloy A (AlMnZr): Starts at around 90 HV and only reaches a peak of 105-110 HV.

[0211] Alloy B (AlMnCr): Starting from about 98 HV and reaching a peak of 115 HV at 648 K.

[0212] Alloy C (AlMnCrZr - this invention): starts at about 103 HV and reaches a peak of 140 HV at 648 K.

[0213] Alloy D (AlMnCrZr - this invention): starts at about 103 HV and reaches a peak of 140 HV at 648 K.

[0214] This experiment clearly demonstrates how chromium improves the hardening rate. Alloy A hardens by 20 HV (90→110 HV), but alloys C and D harden by 35 HV (103→140 HV). The hardening rate reflects the precipitation sequence, while the hardness at the point of initial printing reflects the solid solution hardening effect. The effect of chromium on alloys is a combination of both.

[0215] In this case, an unexpected synergistic effect was observed, as the coarsening rate was slowed down by the rearrangement of Mn and Cr among different precipitates, and the hardening effect was maintained for a longer period of time, resulting in better temperature stability.

[0216] These curves show that chromium plays a crucial role at the beginning of hardening. Using chromium alone will not yield alloys with the desired properties. When chromium is combined with zirconium, its hardening effect lasts for a longer period.

[0217] Alloys containing only Mn and Zr but no Cr exhibit significantly lower hardness. Higher Mn and Zr contents do not improve this situation.

[0218] Example 3

[0219] The solidification behavior of alloy D according to Example 2 and known forging alloys 6081, 7075 and A360.0 was simulated using the Hill-Gulliver method and the CALPHAD method, as also used in Example 1.

[0220] Figure 1 The simulation results are shown. It should be noted that this figure only includes the region nearing complete solidification, where the solids fraction ranges from 0.8 (80%) to 1 (100%). It can be seen that forged alloys 6061 and 7075 exhibit sloping curves, especially alloy 7075, which shows a sharply changing slope near the end of solidification. This behavior is characteristic of materials where hot cracking occurs at the end of solidification. Cast alloy A360.0 exhibits significantly better behavior, with only a slight decrease near complete solidification. Alloy D of the present invention does not exhibit this decrease at the end of solidification.

[0221] The solidification cracking susceptibility index S for each alloy was determined, and the results are shown in Table 5 below.

[0222] Table 5

[0223]

[0224] Example 4

[0225] In the experiments described in detail above (Example 1 and Table 3), it was found that Fe and Si can each be present in amounts of up to 0.7% by weight without adversely affecting the suitability of the Al-based powder of the present invention for LPBF printing. This is advantageous because it reduces the strain on the raw materials used to form the Al-based powder of the present invention, and allows the use of lower-cost recycled Al-based products to form the powder of the present invention.

[0226] In the numerical experiment (Example 1) detailed above, the effects of Fe and Si on the crack susceptibility of the alloy of the present invention were further investigated. The purpose was to determine whether there were any limitations on the content of Fe and Si if both of these impurities were present in the raw materials used to form the Al-based powder of the present invention.

[0227] like Figure 4 As shown, the solidification temperature is found to be dependent on the total content of Fe and Si. The higher the total content of these elements, the greater the risk of forming a melt with a low solidification temperature, which will lead to thermal cracking problems in LPBF-printed objects when using the powder of the present invention.

[0228] Therefore, although Fe and Si can each exist at a maximum of 0.7% by weight without causing a melt with an excessively low solidification temperature, the total combined concentration should not exceed 0.7% by weight. Figure 4 It is evident that the total amount of Fe and Si should preferably not exceed 0.5% by weight, and even more preferably not exceed 0.2% by weight.

[0229] Example 5

[0230] In other experiments using the LBPR printed objects reported herein, powder variants equivalent to those in the experiments according to Example 2 were produced by means of a nitrogen atomization process, and the resulting powders were classified in the nominal particle size range of 20-53 µm.

[0231] Four powder grades were designed with final chemical compositions as shown in Table 6. Particle size distribution was measured by laser diffraction using a Mastersizer 3000 from Malvern UK, according to ISO 13320-1 1999, and repeated five times (see Table 7). Figure 5 ).

[0232] The particle size range mentioned means that up to 2% by weight of the powder has a particle size above the upper limit and up to 2% by weight of the powder has a particle size below the lower limit.

[0233] These experiments clearly demonstrate the suitability of the Al-based alloy of the present invention for forming Al-based powders with high uniformity.

[0234] Table 6: Alloy Composition

[0235]

[0236] Table 7: Particle Size Distribution

[0237]

[0238] Example 6

[0239] Subsequent experiments tested the LB-PBF processing and experimental design used to establish full-density LB-PBF.

[0240] Different variants of the powder from Example 5 were processed in an EOS M100 machine with a 40 µm spot size and a 200 W (170 W nominal power) Yb fiber laser. Initially, 170 W power, 0.13 mm scan spacing, 1000 mm / s speed, and 0.03 mm layer thickness were maintained as the main processing inputs for the printed samples, along with a standard 67° scan rotation.

[0241] Before each printing, the powder sample is conditioned at 353 K for 4 hours using a drying process.

[0242] The sample was printed in a 10 mm × 10 mm × 10 mm cube format, designed to mark the direction of airflow within the chamber (e.g., ...). Figure 6 (As shown). After printing, these samples are cut with a saw.

[0243] The processability of the material is further derived based on a simple cubic factor design, thereby identifying the conditions of interest with high relative density (approximately 99.5%), such as... Figure 2 As shown in the image.

[0244] This design of experiment (DOE) uses the scanning spacing and laser speed as initial variables, compared with the main processing inputs, to identify the conditions for high-density processing.

[0245] A second DOE was then performed to confirm the processing parameter window, and the values ​​were verified to be experimentally acceptable by printing three additional samples for each point of interest.

[0246] Finally, a third DOE was performed on one of the quaternary alloys (alloy D) to complete the three-dimensional design, with the laser power serving as the third variable in the DOE.

[0247] Table 8 summarizes the range of processing parameters for each alloy used in DOE.

[0248] Table 8: Summary of Processing Parameters Used in Design of Experiments (DOE)

[0249]

[0250] Subsequently, the relative density of the sample was determined using cross-sectional optical microscopy. For example... Figure 6 As shown, the cross section under study is observed along three planes.

[0251] To identify and compensate for porosity, results were then obtained for all three planes, and the average of each sample was taken to minimize the orientation effect on the reported overall porosity values. This was achieved using ImageJ software and with each plane having a diameter of approximately 30-50 mm. 2 Image analysis was performed on the area of ​​each sample segment.

[0252] The first DOE was performed to assess the printability of all four alloys, with the goal of achieving a relative density greater than 95%.

[0253] The printing parameters were developed based on appropriate settings for the EOS M290 machine and modified to suit processing in the EOS M100, compensating for lower power and smaller laser spot size.

[0254] Based on this, the cubic DOE was set to have a fixed laser power of 170 W and a layer thickness of 0.03 mm, but the scanning speed was between 500 and 1500 mm / s and the scanning spacing was between 0.1 and 0.15 mm.

[0255] For each alloy, a total of 13 samples were printed using this DOE, and the results were measured as surface maps, such as... Figure 8 As shown in the image.

[0256] Next, a second DOE was performed to verify the existence of a target relative density with respect to the parameters described above.

[0257] A total of 12 samples were printed, and the resulting surface images are as follows. Figure 7 a) and Figure 7 As shown in b).

[0258] After the second DOE, values ​​much higher than the target relative density were established, and many samples (60-75 samples) of all alloys were printed using parameter settings for the maximum relative density (170 W power, 0.1 mm scan spacing and 1500 mm / s laser speed).

[0259] Finally, a similar conceptual design for the DOE was employed, involving 21 samples with varying laser power, scanning spacing, and speed. The range of processing parameters remained conservative only for samples requiring high density (>98% relative density). Figure 9 a) to Figure 9 b) shows the surface plot generated by the third DOE.

[0260] The final processing window is determined after overlaying all results from three DOEs, based on... Figure 9 c) to Figure 9 The surface plot shown in d) displays the processing parameter window that produces a relatively high density (>99%).

[0261] Example 7 – Microstructure Evaluation and Mechanical Testing

[0262] All samples were prepared as follows: in near all three planes (e.g.) Figure 6 The samples (shown) were center-cut and then mounted in an epoxy thermosetting resin called Polyfast (from Struers). These samples were then ground / polished on a Struers TegraPol 31 machine according to standard Struers aluminum alloy preparation methods. The samples were etched using standard Keller reagents to make the melt pool boundaries visible.

[0263] Example 8 – X-ray Diffraction

[0264] X-ray diffraction (XRD) of finely ground samples (up to 2000 abrasive grain size) was performed using a Bragg-Brentano HD X-ray machine equipped with a Cu source (Kα = 1.5406 Å), with 40 mA and 45 kV as generator settings, scanning between 2θ from 20° to 100°, with a step size of 0.007° and a scan step time of 68.59 seconds.

[0265] X-ray diffraction was performed on all four alloys (A, B, C, and D) in their freshly printed condition. This was performed on samples with high density (>99%).

[0266] exist Figure 10 In the XRD patterns shown, all four alloys exhibit similar information observed from the peaks. The Al peak was found to be shifted to a higher angle (2θ) compared to the expected position of pure Al, attributed to the solute dissolved in the solid solution. Secondary peaks observed in all four alloys were found at 40.4° and 43.0° (see [link to XRD pattern]). Figure 8 (close-up view in b), which is attributed to Al-Mn precipitates, and this was also confirmed by subsequent SEM analysis.

[0267] Figure 11 This paper presents an attempt to correlate the individual content of alloying elements in an alloy with the Al peak shift. This is done using the "rule of mixture" approach, based on T. Uesugi and K. Higashi, "First-principles studies on lattice constants and local lattice distortions in solid solution aluminum alloys," Computational Materials Science, Vol. 67, pp. 1-10, 2013, which calculates the effect of each element through first-principles calculations for binary Al alloys. The only modification to this calculation is changing the starting point of pure Al to 4.0478 Å (PDF nr.0040787), instead of the 4.016 Å assumed by Uesugi, because this modified value is significantly more reliable according to DIFFRAC.EVA software.

[0268] Therefore, the role of each element is to plot against 4.016 Å, and to perform linear fitting by extrapolating from 4.0478 Å rather than 4.016 Å (e.g.) Figure 10 (as shown in a). These values ​​were then verified against experimental values ​​derived from the strongest Al peak at approximately 45°, as shown in [the diagram]. Figure 10 As shown in b), the experimentally and calculated lattice parameters are observed to be at 0.01 Å.

[0269] Example 9 – Optical Microscopy

[0270] Optical microscopy was performed on a ZEISS Axioscope 7 instrument with auto-zoom, which can stitch together up to 50 mm at 10X optical zoom. 2 Cross-sectional image.

[0271] In the freshly atomized state, after etching with Keller's reagent, the dendritic structure of the powder variant is clearly visible. Figure 12 This image shows an optical microscope image of alloy C, which reveals the dendritic structure formed during the atomization process.

[0272] Optical microscopic images of all three cut sections of alloy D in its freshly printed state after polishing were also obtained at low magnification to show a nominally fully dense structure with an average relative density of 99.56%. Representative images of all three planes of the cut, as previously defined, are shown below. Figure 13 As shown.

[0273] When analyzing all samples, such as those observed for LPBF printed objects on aluminum alloys, two common types of defects were identified. Figure 14 (a) to Figure 14 The first one in (b) is considered to represent keyhole porosity or gas porosity, which occurs when the laser power is too high or the scanning speed is low; this means that the molten layer reaches the particle ejection or evaporation point in some cases. Figure 14 (c) to Figure 14 Another example in (d) shows a porosity without fusion, which occurs when there is too low laser power or too high scanning speed, resulting in insufficient conditions for melting the layer and bonding it with adjacent layers.

[0274] Therefore, the observed porosity may represent the maximum porosity under non-optimized LPBF printing conditions.

[0275] Example 10 – Electron Microscopy

[0276] Microstructure evaluation was performed on a Leo Gemini 1550 SEM equipped with a field emission gun. Imaging was performed using secondary electron (SE) and backscattered electron (BSE) receivers, depending on the type of image contrast required. X-ray energy-dispersive spectroscopy (EDS) for trace analysis was performed by combining the secondary electron receiver with INCA X-sight software for compositional analysis.

[0277] Selected samples were analyzed using a Zeiss Gemini SEM 450 scanning electron microscope (SEM) with a field emission gun source for micro-analysis. The microscope was equipped with a Bruker Quantax FlatQuad energy-dispersive X-ray spectroscopy (EDX) detector, which enables elemental mapping of microstructures at sub-micron resolution.

[0278] For all alloys, the powder in a just-atomized state exhibits a pure dendritic structure. Figure 15 Image a shows a cross-sectional view of the particles in alloy C. High-magnification imaging further reveals some fine nanoscale particles (<100 nm) within the dendrites, such as... Figure 15 As shown in b. The chemical composition of these particles cannot be identified because they are too small to be resolved using EDS in SEM.

[0279] After printing, the cubes are polished and lightly etched, and then characterized in SEM using backscattered electron microscopy (BSE) to easily identify grain contrast. Figure 16 The image shows images along the front and top planes, revealing a distinct columnar structure along the construction direction.

[0280] Further investigation of the anterior plane revealed the formation of two distinct types of precipitates within the alloy, characterized by their morphology and location within the sample:

[0281] 1. Precipitates located at the boundary of the melt pool.

[0282] 2. Precipitates located at grain boundaries / subgrain boundaries.

[0283] Precipitates located at the boundary of the melt pool are shown in Figure 17 Among them, it can be seen that the precipitate has two typical sizes: smaller precipitates with spherical morphology and average size <100 nm; and larger precipitates with size between 200 and 500 nm.

[0284] The second category of precipitates forms at grain boundaries / subgrain boundaries (cell boundaries). They are rod-shaped with a length of 200-400 nm and a small diameter (<50 nm). Figure 18 The precipitate is clearly visible in the image.

[0285] It is speculated that secondary phase particles aligned with themselves along grain boundaries / subgrain boundaries indicate cellular-type solidification. These particles can form along such boundaries during solidification, thereby producing cellular solidified structures.

[0286] Figure 19This section presents an overview of EDS point scans based on these three types of precipitates. Compared to the volume of the precipitate, EDS exhibits significantly greater diffusion for smaller particles due to the X-ray interaction volume generated by the electron beam of SEM. However, to compensate for this, a higher number of point scans (>25) are performed at high resolution to obtain consistent results.

[0287] The results showed that the Mn concentration in the precipitates at smaller melt pool boundaries and subgrain boundaries was higher than the matrix concentration, but this does not indicate that they are Al. 12 Mn (mesostable state) or Al6Mn (steady state). However, it has been shown that the precipitates at the larger melt pool boundary have stoichiometry close to that of Al6Mn.

[0288] After characterizing the size, morphology, and chemical properties of the secondary phase particles in the rigid atomized and rigidly printed states, the chemical composition of the matrix was assessed in SEM using EDS chemical analysis. The matrix was observed to be generally supersaturated with elements from the solid solution (excluding nanoparticles formed at melt pool boundaries or grain boundaries / subgrain boundaries). EDS analyses of all alloy matrices in the rigid atomized (AA) and rigidly printed (AP) states are summarized in Table 9. The balance in the alloy composition is aluminum.

[0289] Table 9: Chemical composition of all four Al alloys in the lightly atomized (AA) and lightly printed (AP) conditions, determined by EDS.

[0290]

[0291] Occasionally, some melt pools may exhibit a refined grain structure. Figure 20 b) This effect is shown in alloy D. Several Zr-containing particles are also observed in a freshly printed state, exhibiting faceted or dendritic structures; see [link to previous section]. Figure 21 Note that the size variation of this structure is an order of magnitude. The formation of larger faceted particles is not clear, and the average composition of the two precipitates is shown in Table 10, which shows the apparent composition of the faceted Al-Zr structure obtained by EDS, as follows: Figure 21 (a) and Figure 21 Seen in alloys C and D in (b).

[0292] Table 10: Apparent composition of precipitates

[0293]

[0294] Experimental Discussion

[0295] The solidification cracking problem commonly encountered in LPBF printing of certain Al alloys is addressed in this study. This problem is solved by including elements that are less prone to segregation during solidification in the alloy composition according to the invention.

[0296] This study introduces a new series of aluminum alloys comprising two ternary variants and two quaternary variants. The alloy compositions are designed to take advantage of the unique processing conditions provided by the rapid solidification and remelting that occur during laser-based powder bed fusion in additive manufacturing.

[0297] This alloy design principle follows the prediction of alloy solidification using ThermoCalc software. This provides a simple path to completely avoid solidification cracking. The basis of the alloy design is the development of aluminum alloys containing high amounts of solute in supersaturated solid solutions, which can then be used to achieve high strength through the integration of solid solution strengthening and secondary phase precipitate strengthening.

[0298] This alloy is expected to be suitable for high-temperature applications because it exhibits resistance to hardening up to 523 K in the freshly printed condition. Microhardness results show that, when aged at 678 K, the average hardness increases from 105 HV in the freshly printed condition to 130 HV in the potential peak aging condition.

[0299] Solidification cracking can be described as a shift in the concentration gradient of the liquid as solidification occurs, resulting in a higher solute content in the final liquid and thus causing greater volume shrinkage upon solidification. These cracks can form between solidified grains and span several grains. This implies the existence of a minimum temperature gradient, ensuring a sufficient amount of liquid metal is available for bonding along the two grain boundaries and preventing crack formation. Hill solidification curves show a significantly reduced susceptibility to solidification cracking in the novel Al alloy developed in this regard. Experimental design confirmed this crack resistance, and all fabricated samples were found to be crack-free in their freshly printed state.

[0300] Due to the rapid melting and solidification of thin layers (typically 20-40 µm) of metal powder, LPBF processing involves high cooling rates, resulting in cooling rates in the range of 10³-10⁵ K / s. This provides alloy designers with the opportunity to employ higher elemental supersaturation in the Al matrix during AM processing.

[0301] The alloying elements and composition are selected based on studies of rapidly solidifying materials, which advantageously utilize the benefits of incorporating transition metal elements into aluminum alloy solid solutions.

[0302] EDS results showed that, apart from the formation of a small amount of precipitates (subgrain boundary precipitates) during solidification and layer remelting observed near the melt pool boundary, almost all Mn and Cr dissolved in the matrix. This amount of precipitation is surprisingly small compared to other known alloys in the art.

[0303] Because of the ongoing effects of layer remelting and segregation along the solidification front during processing, the formation of secondary precipitates cannot be completely avoided in the freshly printed state, and nanophases containing Mn and Cr have been observed. In this field, other authors have experimentally demonstrated the segregation of Mn along grain boundaries in rapidly solidified nanocrystalline Al-Mn-based alloys, showing that higher concentrations of Mn lead to the formation of a "quasicrystalline" phase, a metastable phase with a composition close to Al6Mn, exhibiting fivefold symmetry with the Al matrix.

[0304] The strength of alloys proposed for rigid printing is a combination of solid solution strength and precipitate reinforcement. Nanoscale precipitates help restrict the movement of dislocations, and the smaller size of these precipitates helps to reinforce the matrix more efficiently.

[0305] As mentioned above, even higher strengths can be achieved after aging by balancing the coarsening of grain boundary / subgrain boundary precipitates with the nucleation / growth of new nanoscale L12Al3Zr precipitates.

[0306] Since some nanoprecipitates containing Mn and Cr can be quasicrystals or precursors of quasicrystals, it is necessary to develop thermal aging treatments to control their formation and prevent them from forming stable orthorhombic Al6Mn phases.

[0307] It can be inferred that by utilizing AI x The strengthening of this new alloy series by Mn and Al3Zr precipitates is a competitive mechanism between these two precipitates, which requires further research and development to properly tune the hardness and provisional strength values. Advantageously, the alloys in rigid-printed form can withstand aging for up to 24 hours at 523 K.

[0308] Ending Notes

[0309] Although the invention has been described in detail for illustrative purposes, it should be understood that such details are for illustrative purposes only, and that variations may be made therein by those skilled in the art in carrying out the claimed subject matter, based on the drawings, the disclosure and the appended claims.

[0310] It should be understood that the embodiments shown in the accompanying drawings are for illustrative purposes and should not be construed as limiting the invention. Unless otherwise indicated, the drawings (e.g., crosshairs, arrangement of components, scale, extent, etc.) are intended to be interpreted in conjunction with the specification and are considered an integral part of the entire written description of the invention.

[0311] The term "comprising" as used in the claims does not exclude other elements or steps. The indefinite article "a / an" as used in the claims does not exclude the plural form. A single processor or other unit may perform the functions of several means described in the claims. Reference numerals used in the claims should not be considered as limiting the scope.

Claims

1. A pre-alloyed Al-based powder suitable for additive manufacturing, consisting of, based on the total weight of the pre-alloyed Al-based powder: 3 - 5.5 wt% of Mn, 0.2 - 2 wt% of Zr, 0.6 - 1.10 wt% of Cr, 0 - 2 wt% of Mg, altogether at most 0.7 wt% of Fe and Si, at most 0.7 wt% of O as unavoidable impurities, and at most 0.5 wt% of other unavoidable impurities, the balance being Al.

2. The pre-alloyed Al-based powder according to claim 1, wherein the content of Zr is 0.3 - 1.9 wt%.

3. The pre-alloyed Al-based powder according to claim 1, wherein the total content of Fe and Si is at most 0.2 wt%.

4. The pre-alloyed Al-based powder according to claim 2, wherein the total content of Fe and Si is at most 0.2 wt%.

5. The pre-alloyed Al-based powder according to any one of claims 1 - 4, wherein the total content of Mn + Zr + Cr is 4 - 8.5 wt%.

6. The pre-alloyed Al-based powder according to any one of claims 1 - 4, wherein the pre-alloyed Al-based powder is produced by melting material having the desired composition and inert gas atomization of the melt.

7. The pre-alloyed Al-based powder according to claim 5, wherein the pre-alloyed Al-based powder is produced by melting material having the desired composition and inert gas atomization of the melt.

8. The pre-alloyed Al-based powder according to any one of claims 1 - 4, having a particle size range of 10 - 53 pm as measured by laser diffraction according to ISO 13320-1 1999, wherein less than 10% of the particles fall outside the specified particle size range.

9. The pre-alloyed Al-based powder according to claim 5, having a particle size range of 10 - 53 pm as measured by laser diffraction according to ISO 13320-1 1999, wherein less than 10% of the particles fall outside the specified particle size range.

10. The pre-alloyed Al-based powder according to claim 6, having a particle size range of 10 - 53 pm as measured by laser diffraction according to ISO 13320-1 1999, wherein less than 10% of the particles fall outside the specified particle size range.

11. The pre-alloyed Al-based powder according to claim 7, having a particle size range of 10 - 53 pm as measured by laser diffraction according to ISO 13320-1 1999, wherein less than 10% of the particles fall outside the specified particle size range.

12. A method for producing an article by additive manufacturing, comprising the following steps: - providing a pre-alloyed Al-based powder according to any one of claims 1 - 11; - depositing a layer of the pre-alloyed Al-based powder, thereby forming a powder bed; - heating the pre-alloyed Al-based powder by a laser or electron beam in a predetermined pattern to fuse or melt the particles to be bonded; - cooling the fused or melted particles; - repeating the depositing, heating and cooling steps until an additively manufactured article is formed. - recovering the produced article.

13. The method according to claim 12, followed by a heat treatment, wherein the heat treatment comprises heating the produced article in an air atmosphere to a temperature of 150-450 °C for a time period of at least 0.5 hours.

14. An article formed by the method of claim 12 or 13.

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