Preparation method of aluminum alloy with high-density twin crystal and 9R phase

Abstract

The invention discloses a preparation method of an aluminum alloy with high-density twin crystals and a 9R phase, and belongs to the technical field of metal 3D printing. The alloy comprises the following components in percentage by weight: 3.5 to 9 percent of Mg, 0.4 to 0.8 percent of Mn, 0.5 to 1.5 percent of Er, 0.5 to 2.5 percent of Zr and the balance of Al. The Er and Zr are prepared into powder and then are subjected to 3D printing, under the influence of the rapid cooling rate of 3D printing, the solid solubility of Er and Zr is greatly improved, the stacking fault energy of Al is reduced, and the possibility of introducing nano-scale plane defects such as twin crystals and 9R phases is increased. The L12-Al3 (Er, Zr) primary phase serves as a transition structure and becomes a starting core of twin crystal growth, the mechanical property of the aluminum alloy is greatly improved, the tensile strength of the aluminum alloy reaches 531 MPa, the yield strength reaches 455MPa, and the ductility reaches 27%.

Description

Technical Field

[0001] The present invention belongs to the technical field of Al alloy preparation (additive manufacturing field), and specifically relates to an additive manufacturing aluminum alloy material with high-density twins and 9R phase. The Al alloy with the above-mentioned nano-scale planar defects such as twins has excellent comprehensive properties and will have important applications in industrial preparation and aerospace fields. Background Art

[0002] Due to excellent physical and chemical properties such as excellent specific strength, low density, electrical conductivity, thermal conductivity, etc., aluminum alloys have become the most widely used light alloys in fields such as aerospace, engineering manufacturing, and transportation. Laser powder bed fusion (L-PBF) is an additive manufacturing technology that can quickly realize the manufacturing of complex parts. Therefore, improving the mechanical properties of LPBFed Al alloys is of great significance for their industrial applications. However, like most metal materials, aluminum alloys also face the trade-off between strength and ductility. An interesting solution is to introduce twins and 9R phase into the alloy.

[0003] Twin boundaries are strong barriers to dislocation movement and can disperse the strain flow, which is crucial for increasing strength and ductility. Twins tend to form in face-centered cubic (FCC) metals with low stacking fault energy (SFE) (less than 50 mJ / m 2 )). The formation energy of the 9R phase is even higher than that of twins. Therefore, it is more difficult to form the 9R phase-containing material than twins in FCC materials. The 9R phase with a higher twin formation energy consists of periodic stacking faults (SFs), usually one SF per three atomic layers. As is well known, FCC Al has an extremely high SFE, up to 166 mJ / m2. Therefore, it is extremely difficult to obtain bulk Al alloys with twins and 9R phase. By using some SFE elements (e.g., Mg and Zr), the stacking fault energy can be reduced. However, elements with high solubility (e.g., Mg) have limited effect on reducing the stacking fault energy, or the element solubility is limited (e.g., the maximum solid solubility of Zr: 0.046 at.%). Therefore, twins and 9R phase are usually obtained in nanocrystalline Al-based materials (e.g., by magnetron sputtering). It is very difficult to occur twinning in coarse-grained bulk Al alloys unless extremely high strain rates or low temperatures are applied. The rapid cooling characteristics of the LPBF process can generate high internal stress and increase the solid solubility of elements, providing the possibility of introducing twins and 9R phase into coarse-grained (>1 μm) or ultrafine-grained (100 nm - 1 μm) bulk Al materials without applying deformation or low-temperature conditions. Therefore, it is necessary to explore reasonable Al alloy composition design and reproducible and scalable schemes to obtain Al alloys with high-density twins and 9R phase and improve the mechanical properties of the alloys.

[0004] A typical way to create high-density growth twins in Al is to introduce a buffer layer, but this method is currently only achieved in the production of Al thin films by magnetron sputtering. The difficulties in introducing growth twins into bulk Al alloys through buffer layers mainly focus on the following aspects: (1) Finding a suitable buffer structure and maintaining a coherent relationship with Al as much as possible. (2) Being able to play an epitaxial growth role on Al grains. (3) Being able to be used in the production of bulk Al alloy materials, such as not causing changes in the melting point of Al alloys. (4) Preferably being elements commonly used in the industrial production of bulk Al alloys, or even low-cost elements. (5) Selecting a preparation method and preparation process that match the selected elements.

[0005] The most commercially available Al-Mg alloy is that of European aircraft manufacturer Airbus In the as-cast state, the tensile strength is 418 MPa, the yield strength is 231 MPa, and the elongation is 10%. In the heat-treated state, the tensile strength is 539 MPa, the yield strength is 509 MPa, and the elongation is 8%. This alloy has a unique microstructure and can remain stable at high temperatures. It is better than ordinary aluminum alloys in terms of fatigue resistance, weldability, specific strength, and ductility, and is very suitable for the aerospace, defense, and transportation fields. Summary of the Invention

[0006] The present invention aims to provide a 3D printing process for preparing an aluminum alloy with twins and 9R equivalent nanoscale planar defects. The aluminum alloy prepared by the present invention breaks through the microstructure mode of traditional aluminum alloys, proposes a new method for preparing aluminum alloys with nanoscale planar defects, enables aluminum alloys to break the strength-ductility trade-off, further improves the performance compared to commercial alloys, and almost has the best strength-ductility matching as-cast aluminum alloy to date, significantly enhancing its service performance. Even without heat treatment, it exhibits highly competitive mechanical properties.

[0007] The aluminum alloy material of the present invention mainly consists of elements Mg, Mn, Er, and Zr, and their mass percentages are as follows:

[0008] The content of Mg is 3.5 - 9%, the content of Mn is 0.4 - 0.8%, the content of Er is 0.5 - 1.5%, the content of Zr is 0.5 - 2.5%, and the rest is Al and inevitable impurities.

[0009] The above components are preferably: the content of Mg is 5.5 - 8.0%, the content of Mn is 0.6 - 0.8%, the content of Er is 0.8 - 1.2%, the content of Zr is 1.8 - 2.2%, and the rest is Al and inevitable impurities.

[0010] The method for preparing an aluminum alloy with nanoscale planar defects provided by the present invention is characterized in that it uses a 3D printing and heat treatment process, including the following steps:

[0011] Step 1: Alloy powder preparation

[0012] Select high-purity Al, high-purity Mg, Al-Mn master alloy, Al-Er master alloy and Al-Zr master alloy, and put them into the crucible according to the ratio; after the melting equipment is pre-evacuated, introduce argon to melt the alloy at 750 °C. After the alloy is fully melted, raise the temperature to 800 - 840 °C; then perform gas atomization (or vacuum gas atomization) to produce powder in an inert atmosphere (such as argon), and set the gas pressure to 2 MPa; after cooling, screen the powder to obtain alloy powder with a particle size range of 20 - 50 μm;

[0013] Step 2: Powder drying

[0014] Place the alloy powder obtained in Step 1 in a vacuum drying oven and dry it at 80 °C to remove the moisture and adsorbed gas on the powder surface;

[0015] Step 3: 3D printing

[0016] Use a selective laser melting (SLM) metal printer to perform 3D printing on the powder dried in Step 2. The specific process parameters are as follows:

[0017] Laser power (P): 280 - 420 W, scanning speed (V): 800 - 1800 mm / s, hatch distance (h): 100 - 200 μm, powder layer thickness (t): 20 - 30 μm, (substrate for powder spreading) preheating temperature: 80 - 200 °C, volumetric energy density (E): 80 - 500 J / mm 3 , E = P / (V * h * t), defocus distance (Δf): 0 - 2.5 mm.

[0018] Different compositions have different optimal process parameters, which need to be explored according to the specific composition. The criterion for selecting the optimal process is the relative density, and the process parameters with the highest relative density are the optimal process.

[0019] The preferred printing process parameters for the preferred composition are: laser power: 280 - 300 W, scanning speed: 800 - 1200 mm / s, hatch distance 120 - 140 μm, layer thickness: 20 - 30 μm, preheating temperature: 100 - 150 °C, defocus distance 2.5 mm, volumetric energy density: 55 - 155 J / mm 3 .

[0020] The obtained product has equiaxed ultrafine grains (grain size: 100 nm - 1 μm) and micron-sized coarse grains (grain size ≥ 1 μm), and the center of the melt pool is coarse grains, while the boundary of the melt pool is ultrafine grains, and the proportion of ultrafine grains is extremely high.

[0021] The obtained product has twins and the 9R phase, and both ultrafine grains and coarse grains contain twins and the 9R phase. The growth of Al twins is carried out with the L12-Al3M phase as the core (buffer structure).

[0022] The aluminum alloy with high-density twins and the 9R phase obtained in this invention has excellent mechanical properties. Under the preparation process, the tensile strength of the as-formed alloy at room temperature can reach 531 MPa, the yield strength can reach 455 MPa, and the elongation rate is as high as 27%, significantly improving the comprehensive service performance of the aluminum alloy. It breaks the strength-ductility trade-off and has the best strength-elongation rate matching among the currently additively manufactured as-formed aluminum alloys.

[0023] The additively manufactured aluminum alloy of this invention shows a more excellent strength-elongation rate matching compared with the currently commercial aluminum alloys. More importantly, the cost is much lower than that of the Sc-containing aluminum alloy, which is due to the fact that the cost of the Er-containing master alloy is much lower than that of the Sc-containing master alloy.

[0024] This invention has the following beneficial effects:

[0025] This invention proposes a new manufacturing method for additively manufactured bulk aluminum alloys with twins and the 9R phase, as well as a new mechanism for the generation of Al growth twins. By adding Er and Zr elements that can reduce the stacking fault energy, the extremely high cooling rate during the LPBF process increases the solid solubility of Er and Zr, thus greatly reducing the stacking fault energy of the aluminum alloy and further increasing the possibility of twin generation. Most importantly, during the additive manufacturing process, primary L12-Al3(Er,Zr) phases with heterogeneous nucleation effects will be generated. The Al3(Er,Zr) phase is an enrichment area of Er and Zr elements and has a lower stacking fault energy. In particular, the Al3(Er,Zr) phase has a completely coherent structure with the Al matrix. Therefore, through the low stacking fault energy Al3(Er,Zr) phase as a buffer structure (attached Figure 2 ), the coherent Al3(Er,Zr) / Al interface allows the replication of twins and stacking defects from Al3(Er,Zr) to Al (attached Figure 2 ), thereby generating additively manufactured bulk aluminum alloys with high-density growth twins (attached Figure 1 , Figures 2 and 3 specifically show). That is, the formation mechanism of Al growth twins with the twin L12 structure as the core. During the LPBF process, due to the generation of extremely high internal stresses, the incoherent twin boundaries of coarse grains or ultrafine grains dissociate to form the 9R phase. Finally, additively manufactured aluminum alloys with high-density twins and the 9R phase are obtained. Twin boundaries are powerful barriers to dislocation movement and can disperse the strain flow at the same time, which is crucial for increasing strength and ductility. Through the above beneficial effects, the high-strength and high-toughness additively manufactured aluminum alloy as described above is obtained. Description of the Drawings

[0026] Figure 1: Electron backscatter diffraction (EBSD) pattern of the as-cast alloy in Example 1 (thin lines are grain boundaries and thick lines are twin boundaries).

[0027] Figure 2 : Transmission electron microscopy (TEM) image of the growth of Al twins in the as-cast alloy in Example 1 with twin Al3(Er,Zr) phase as the nucleus (buffer structure).

[0028] Figure 3 : TEM images of the as-cast alloy in Example 1, (a) TEM image of the twin boundary; (b) high-resolution TEM image of the twin boundary; (c) diffraction spots of the twin boundary; (d) high-resolution TEM image of the 9R phase; (e) enlarged HRTEM image of the white box; (f) Fourier transform image derived from the 9R phase.

[0029] Figure 4 : Stress-strain curves of Example 1 and the comparative example. Specific implementation manners

[0030] The present invention will be further described below in conjunction with the embodiments, but the present invention is not limited to the following embodiments.

[0031] Example 1

[0032] This example discloses a high-strength and high-toughness aluminum alloy with high-density twins and 9R phase and its preparation method. The specific implementation steps are as follows:

[0033] I. Alloy composition design

[0034] The chemical composition of the aluminum alloy is as follows by mass percentage: Mg 7.8%, Mn 0.8%, Er 0.9%, Zr 1.9%, and the balance is Al and unavoidable impurity elements. The alloy element ratio design is based on the principle of multi-component synergistic strengthening, and the material properties are improved through the microalloying effect of rare earth Er and transition elements.

[0035] II. Powder preparation process

[0036] Raw material pretreatment: Select high-purity Al ingots (99.99%), high-purity Mg ingots (99.95%) and Al-10Mn, Al-5Er, Al-10Zr master alloys, weigh them according to the designed ratio and place them in a graphite crucible;

[0037] Vacuum melting: Heat up to 750 °C for melting under a vacuum of 10 -2 Pa, and after complete melting, introduce high-purity argon (99.999%) as the protective atmosphere;

[0038] Gas atomization powder making: After superheating the melt to 800 - 840 °C, atomize it with high-pressure argon (2 MPa), and obtain pre-alloyed powder with a sphericity > 95% by controlling the nozzle structure.

[0039] Powder classification: After rapid solidification in an air cooling system, qualified powder with a particle size range of 20-50 μm is screened out using a vibration screening machine and dehydrated in a vacuum drying oven at 120°C for 4 hours.

[0040] 3. Additive Manufacturing Process

[0041] The specimens were formed using the selective laser melting (SLM) technology, and the main process parameters were:

[0042] Laser system: SLM Solutions M280

[0043] Laser power: 290W

[0044] Scanning strategy: chessboard scanning, speed 1000mm / s

[0045] Forming parameters: layer thickness 30 μm, hatch distance 130 μm.

[0046] Environmental control: substrate preheating 150°C, defocusing amount +2.5mm.

[0047] Volume energy density: 74J / mm 3 .

[0048] 4. Post-processing process

[0049] After the molded part is cut and the substrate is peeled off, no heat treatment is required. The standard tensile specimen is made by warp cutting and surface polishing in accordance with GB / T228.1-2010 standard.

[0050] By optimizing alloy composition and synergistically controlling process parameters, this approach achieved a density of 99.8%. In particular, it produced an additively manufactured bulk aluminum alloy with a high density of twins and 9R phases. This overcomes the significant technical challenge of introducing nanoscale planar defects, such as twins, into bulk aluminum alloys, a key innovation of this paper. Tensile testing demonstrated that the resulting alloy achieved a room-temperature tensile strength of 530 MPa and an elongation exceeding 27%, making it suitable for direct forming of complex structural components in the aerospace industry.

[0051] Electron backscatter diffraction (EBSD) pattern of the alloy as formed in Example 1 (thin lines are grain boundaries, thick lines are twin boundaries) is shown in Figure 1 .

[0052] Example 1 Transmission electron microscope (TEM) image of Al twin growth with twin Al3 (Er, Zr) phase as core (buffer structure) is shown in Figure 2 .

[0053] TEM images of the alloy in Example 1, (a) TEM image of the twin boundary; (b) high-resolution TEM image of the twin boundary; (c) diffraction spots of the twin boundary; (d) high-resolution TEM image of the 9R phase; (e) enlarged HRTEM image of the white frame; (f) Fourier transform image derived from the 9R phase, see Figure 3 。

[0054] Stress-strain curves of Example 1 and the comparative example, see Figure 4 。

[0055] Example 2

[0056] This example discloses a high-strength aluminum alloy and its preparation method. The specific implementation steps are as follows:

[0057] I. Alloy composition design

[0058] The chemical composition of the aluminum alloy is as follows by mass percentage: Mg 4.5%, Mn 0.7%, Er 0.6%, Zr 1.3%, and the balance is Al and inevitable impurity elements. The alloy element ratio design is based on the principle of multi-component synergistic strengthening, and the material properties are improved through the microalloying effect of rare earth Er and transition elements.

[0059] II. Powder preparation process

[0060] Raw material pretreatment: Select high-purity Al ingots (99.99%), high-purity Mg ingots (99.95%) and Al-10Mn, Al-5Er, Al-10Zr master alloys. After weighing according to the designed ratio, place them in a graphite crucible;

[0061] Vacuum melting: Heat up to 750 °C for melting under a vacuum of 10 -2 Pa, and after complete melting, introduce high-purity argon (99.999%) as the protective atmosphere;

[0062] Gas atomization powder making: After superheating the melt to 800 - 840 °C, atomize it with high-pressure argon (2 MPa), and obtain pre-alloy powder with a sphericity > 95% by controlling the nozzle structure;

[0063] Powder classification: After rapid solidification by the gas cooling system, use a vibrating sieve to screen out qualified powder with a particle size range of 20 - 50 μm, and dehydrate it in a vacuum drying oven at 120 °C for 4 hours.

[0064] III. Additive manufacturing process

[0065] Use selective laser melting (SLM) technology to form specimens. The main process parameters are:

[0066] Laser system: SLM Solutions M280 type equipment

[0067] Laser power: 350 W

[0068] Scanning strategy: chessboard scanning, speed 1200 mm / s

[0069] Forming parameters: layer thickness 30 μm, hatch distance 130 μm.

[0070] Environmental control: substrate preheating 150 °C, defocus amount +2.5 mm

[0071] Volume energy density: 75 J / mm 3

[0072] IV. Post-processing technology

[0073] After the formed part is peeled from the substrate by wire cutting, no heat treatment is required. The standard tensile sample is made by wire cutting and surface polishing according to the GB / T228.1-2010 standard.

[0074] This solution solves the problem of easy generation of hot cracks in traditional aluminum alloys during additive manufacturing by optimizing the synergistic control of alloy composition and process parameters, and the relative density reaches 99.5%. After tensile testing, the as-formed alloy has a room-temperature tensile strength of 390 MPa, a yield strength of 354 MPa, and an elongation of 18%, which is suitable for the direct forming manufacturing of complex structural parts in the aerospace field.

[0075] Comparative example

[0076] This comparative example is a casting aluminum alloy and its preparation method. The specific implementation steps are as follows:

[0077] The chemical composition of the aluminum alloy is calculated by mass percentage as follows: Mg 5.0%, Mn 0.8%, Er 0.2%, Zr 0.1%, and the balance is Al and unavoidable impurity elements.

[0078] Raw material pretreatment: High-purity Al ingots (99.99%), high-purity Mg ingots (99.95%) and Al-10Mn, Al-5Er, Al-10Zr master alloys are selected, weighed according to the designed ratio and placed in a graphite crucible;

[0079] Melting: Use an electric resistance furnace for melting. First, slowly heat up the electric resistance furnace, then put the graphite crucible containing high-purity aluminum into the melting furnace, keep the melting temperature at 780 °C, add the master alloy after the high-purity aluminum melts, degas with hexachloroethane and stir evenly after all melting, let it stand for 30 min and then pour it into an iron mold for casting, and finally obtain an Al alloy ingot.

[0080] The ingot is made into a standard tensile sample by wire cutting and surface polishing according to the GB / T228.1-2010 standard.

[0081] The comparative examples are not within the scope of protection of the present invention and mainly serve for comparison. This process is already a relatively mature process for manufacturing aluminum alloys.

[0082] Table 1 shows the properties of the examples and comparative examples of the present invention. Compared with the two examples, the current commercial aluminum alloys have achieved significant improvements in yield strength and elongation. The main reasons for such a large improvement in mechanical properties are as follows: 1). Significant grain refinement brings about grain refinement strengthening (Hall-Petch effect). As can be seen from Figure 1 it that a large part of the grain size in the alloy is <1 μm, and these grains are mainly distributed at the melt pool boundary. While the center of the melt pool is mainly composed of coarse grains with a grain size of 1-10 μm. The significant grain refinement contributes significantly to the improvement of strength. 2). The strength improvement brought by nano-scale planar defects such as twins. Among different microstructural defects, twin boundaries and 9R phases become effective barriers for storing dislocations inside grains by hindering the movement of dislocations. Therefore, they can simultaneously improve the strength and strain hardening ability of nano-grained metallic materials. At the same time, twin boundaries can become additional interfaces for dislocation pile-up, contributing to strain shunting and thus improving ductility. Figure 1 The thick line in it represents the twin boundary, and the length of the twin boundary accounts for 4.51% of the total grain boundary length, which is very rare in bulk aluminum alloys. Many scientists at home and abroad have been committed to introducing nano-scale planar defects such as twins into bulk aluminum alloys and have made great efforts in the past few decades. This article provides a means to introduce nano-scale planar defects such as twins into bulk aluminum alloys through composition design combined with preparation processes, and proposes a formation mechanism of Al growth twins with an L12-Al3M buffer structure. Attached Figure 2 , the TEM images of twins and 9R phases in Figure 3 provide direct evidence for our discovery.

[0083] Figure 3 are the stress-strain curves of Example 1 and the comparative example. Compared with the cast alloy, the alloy prepared by additive manufacturing has achieved a synergistic improvement in strength and ductility, indicating that the strength-ductility trade-off has been broken (the strength-ductility trade-off means that in metallic materials, as the strength increases, the elongation will decrease). This shows the advantage of introducing nano-scale planar defects such as twins into additive manufacturing aluminum alloys to achieve a synergistic improvement in strength and elongation.

[0084] Table 1. Room temperature tensile properties of as-formed aluminum alloys

[0085]

Claims

1. A method for preparing an aluminum alloy with high-density twins and a 9R phase, characterized in that, The aluminum alloy is mainly composed of elements Mg, Mn, Er, and Zr, and their mass percentages are as follows: The content of Mg is 3.5 - 9%, the content of Mn is 0.4 - 0.8%, the content of Er is 0.5 - 1.5%, the content of Zr is 0.5 - 2.5%, and the rest is Al and inevitable impurities; Its preparation method adopts 3D printing and heat treatment processes, including the following steps: Step 1: Alloy powder preparation Select high-purity Al, high-purity Mg, Al-Mn master alloy, Al-Er master alloy, and Al-Zr master alloy, and put them into the crucible according to the ratio; after the melting equipment pre-pumps the vacuum, argon is introduced at 750 °C for alloy melting. After the alloy is fully melted, the temperature is raised to 800 - 840 °C; then gas atomization (or vacuum atomization) is carried out to produce powder under an inert atmosphere (such as argon), and the gas pressure is set to 2 MPa; after cooling, the powder is sieved to obtain alloy powder with a particle size range of 20 - 50 μm; Step 2: Powder drying Put the alloy powder obtained in Step 1 into a vacuum drying oven and dry it at 80 °C to remove the moisture and adsorbed gas on the powder surface; Step 3: 3D printing Use a selective laser melting (SLM) metal printer to perform 3D printing on the powder dried in Step 2. The specific process parameters are as follows: Laser power (P): 280 - 420 W, scanning speed (V): 800 - 1800 mm / s, hatch distance (h): 100 - 200 μm, layer thickness (t): 20 - 30 μm, preheating temperature: 80 - 200 °C, volumetric energy density (E): 80 - 500 J / mm³, defocus distance (Δf): 0 - 2.5 mm.

2. The method according to claim 1, wherein The composition of the aluminum alloy is: the content of Mg is 5.5 - 8.0%, the content of Mn is 0.6 - 0.8%, the content of Er is 0.8 - 1.2%, the content of Zr is 1.8 - 2.2%, and the rest is Al and inevitable impurities; The printing process parameters are: laser power: 280 - 300 W, scanning speed: 800 - 1200 mm / s, hatch distance 120 - 140 μm, layer thickness: 20 - 30 μm, preheating temperature: 100 - 150 °C, defocus distance 2.5 mm, volumetric energy density: 55 - 155 J / mm³.

3. The preparation method according to claim 1 or 2, characterized in that, The obtained product has equiaxed ultrafine grains (grain size: 100 nm - 1 μm) and micron-sized coarse grains (grain size ≥ 1 μm), and the center of the melt pool is coarse grains, and the boundary of the melt pool is ultrafine grains, and the ultrafine grains account for a very high proportion.

4. The method according to claim 1, characterized in that, The obtained product has twins and 9R phase, and both twins and 9R phase are contained in the ultrafine grains and coarse grains.

5. The method according to claim 1, wherein The obtained product grows Al twins with the L12-Al3M phase as the core (buffer structure).

6. An aluminum alloy 3D printing product prepared by the method according to any one of claims 1 - 5.

7. The aluminum alloy 3D printing product prepared by the method according to any one of claims 1-5 has a room temperature tensile strength of up to 530 MPa or more, a yield strength of 450 MPa or more, and an elongation of 27% or more.

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