TiB2 / Al interface modification method based on synergistic regulation of alloy composition and solidification cooling rate

Abstract

The application provides a TiB2 / Al interface modification method based on alloy composition-solidification cooling rate synergistic regulation, which comprises the steps of phase diagram thermodynamic calculation, nucleation theory calculation, drawing of crystal nucleus density diagram, determination of alloy composition, preparation of solidification forming process and the like. The internal principle of the method is that: through composition design and cooling rate regulation, the low-mismatch structure primary phase uniformly modifies the surface of TiB2 particles in the solidification process, forms a low-mismatch "TiB2 / L12 / alpha-Al" multi-level interface, and further improves the solidification fine-grain efficiency and mechanical strengthening effect of the TiB2 particles. With the developed titanium diboride reinforced aluminum-magnesium-scandium-zirconium alloy powder for additive manufacturing as raw material, the bulk material prepared by adopting the laser directed energy deposition technology has excellent mechanical properties, the yield strength is 257 MPa, the tensile strength is 385 MPa, and the elongation is 13.8%, which is better than other directed energy deposition aluminum alloy materials.

Description

Technical Field

[0001] This invention belongs to the field of aluminum alloy smelting and solidification preparation, and relates to a TiB2 / Al interface modification method based on the synergistic regulation of alloy composition and solidification cooling rate. Specifically, it relates to a method for regulating the titanium diboride / aluminum interface in aluminum alloys, as well as a method for designing, preparing and rapidly solidifying high-strength aluminum alloy powder for additive manufacturing. Background Technology

[0002] Adding ceramic particles to aluminum alloys is an effective way to optimize solidification structure and enhance strength and toughness, and it has been widely used in traditional casting and emerging additive manufacturing, where solidification is the main forming process. Examples include grain refinement techniques using TiB2 and TiC as nucleation inoculants, and the preparation of lightweight, high-strength aluminum-based composite materials using TiB2, SiC, and SiN as reinforcements. The compatibility between ceramic particles and the aluminum matrix, i.e., the properties of the "ceramic / aluminum" interface, determines the effects of solidification grain refinement, particle dispersion, and enhancement and toughening. However, the mismatch degree at the "ceramic / aluminum" interface is generally greater than 4%, far exceeding the <1% required for low-stress coherent interfaces, significantly weakening the strengthening effect of ceramic particles. Currently, interfacial element doping is commonly used to reduce the "ceramic / aluminum" interfacial energy and improve its stability. However, this method cannot change the atomic arrangement at the interface, thus having limited effect on improving interfacial compatibility. By utilizing the segregation and ordering of liquid-phase alloying elements at the interface, a coherent two-dimensional interfacial phase can be obtained, which can significantly reduce the mismatch at the ceramic / aluminum interface. However, this two-dimensional interfacial phase is thermodynamically unstable, easily dissolving and returning to the liquid phase, or being poisoned by other alloying elements, leading to structural instability. Therefore, there is currently a lack of a stable and reliable interfacial modification method that can significantly improve the compatibility of the ceramic / aluminum interface.

[0003] With L12, D0 22 and D0 23The Al3M phase (where M is a group IIIB-VB element such as Sc, Zr, Ti, Nb, V, B, Ti, etc.) has a similar crystal structure to FCC aluminum, with a mismatch degree as low as 0.09-1%. According to phase diagram analysis, adding M elements to the melt and appropriately increasing their concentration will cause the primary Al3M phase to precipitate first during solidification, followed by the formation of α-Al. More importantly, TiB2 is a strong heterogeneous nucleation substrate for Al3M, making Al3M an ideal low-mismatch intermediate phase connecting TiB2 and α-Al. However, in conventional casting, the cooling rate is slow, the nucleation rate of primary Al3M is low, and its liquid phase precipitation is mainly a growth process, resulting in large final size, making it impossible to universally obtain a "TiB2 / Al3M / α-Al" low-mismatch composite interface. According to heterogeneous nucleation theory, the nucleation rate of primary Al3M is closely related to the nucleation work and the element diffusion rate. The effects of these two parameters on the nucleation rate show opposite trends with cooling rate (supercooling), leading to a maximum nucleation rate as the cooling rate increases. Therefore, there must exist an optimal cooling rate range that allows Al3M to coat as many TiB2 particles as possible for nucleation and limited growth, maximizing the transformation of the easily poisoned and mismatched TiB2 / α-Al interface into a "TiB2 / Al3M / α-Al" low-mismatch composite interface, which can effectively improve the strengthening effect of TiB2 ceramic particles.

[0004] Determining the optimal amount of element M and the corresponding optimal cooling rate in a scientific and efficient manner to precisely control the nucleation and growth of the Al3M primary phase on the surface of TiB2 particles is the key to obtaining a low-mismatch composite interface of "TiB2 / Al3M / α-Al". Summary of the Invention

[0005] The purpose of this invention is to overcome the shortcomings and deficiencies of existing technologies and provide a TiB2 / Al interface modification method based on the synergistic regulation of alloy composition and solidification cooling rate. This invention is based on the phase diagram calculation method (CALPHAD), which can perform high-throughput calculations of the phase composition and thermodynamic properties of the system in composition space, thus providing strong fundamental data support for solidification nucleation theory calculations. Based on this, this invention couples phase diagram thermodynamic calculations and solidification theory calculations, proposing a synergistic regulation method of "alloy composition-solidification cooling rate" to obtain a "TiB2 / Al3M / α-Al" low mismatch composite interface, significantly improving the interfacial compatibility between TiB2 ceramic particles and the aluminum matrix, and ultimately greatly enhancing the microstructure and toughness of solidified TiB2-reinforced aluminum alloys.

[0006] The underlying principle of this invention is that, through composition design and cooling rate control, the primary phase with low mismatch structure is uniformly modified on the surface of TiB2 particles during solidification, forming a low mismatch "TiB2 / L12 / α-Al" multi-level interface, thereby improving the solidification fine grain efficiency and mechanical strengthening effect of TiB2 particles.

[0007] The objective of this invention can be achieved through the following methods:

[0008] In a first aspect, the present invention provides a method for TiB2 / Al interface modification based on the synergistic regulation of alloy composition and solidification cooling rate, comprising the following steps:

[0009] S1. Phase diagram thermodynamic calculation: The phase diagram thermodynamic results of the Al-Ti-BMS system are calculated using the CALPHAD method.

[0010] S2. Nucleation Theory Calculation: Based on the phase diagram thermodynamic results of the Al-Ti-BMS system obtained in step S1, the nucleation theory is used to calculate the L12 and D0 phases formed in the Al-Ti-BMS system at different temperatures during solidification. 22 D0 23 Phase nucleus density;

[0011] S3. Draw the nucleus density map: Set the concentration step size of M and S, and draw the nucleus density map based on the nucleus density calculated in step S2.

[0012] S4. Determine the alloy composition: Based on the nucleus density diagram drawn in step S3, determine the alloy composition ratio of the Al-Ti-BMS system under the optimal nucleus density;

[0013] S5. Formulate solidification forming process: Based on the crystal nucleus density diagram drawn in step S3, determine the optimal cooling rate under the optimal crystal nucleus density, and then formulate solidification forming process parameters.

[0014] S6. Preparation and Forming: Using Al-Ti-BMS system alloy powder with the alloy composition ratio determined in step S4 as raw material, Al-Ti-BMS system alloy material is prepared according to the solidification forming process parameters in step S5.

[0015] In one embodiment of the present invention, in step S1, M in the Al-Ti-BMS system is L12 and D0. 22 D0 23 One or more of the crystal-forming elements, where S is not L12 or D0 22 D0 23 One or more of the crystal-forming elements.

[0016] Furthermore, M includes one or more of Ti, Sc, Zr, V, Nb, Ta, Hf, Y, Yb, Er, Cr, W, and Li, and S includes one or more of Mg, Zn, Cu, Mn, and Si.

[0017] In some embodiments, the Al-Ti-BMS system is a TiB2 / Al-Mg-Sc-Zr system.

[0018] As one embodiment of the present invention, in step S1, the CALPHAD (CALculation of PHAseDiagram) method includes: using phase diagram thermodynamic calculation software and the Al-Ti-BMS system phase diagram thermodynamic database.

[0019] Furthermore, the phase diagram thermodynamic calculation software includes at least one of Pandat, Thermal-Calc, and FactSage; the Al-Ti-BMS system phase diagram thermodynamic database is the Al-Mg-Sc-Zr-Ti-B system phase diagram thermodynamic database.

[0020] As one embodiment of the present invention, in step S1, the phase diagram thermodynamic results include at least one of phase equilibrium relationship, phase evolution during solidification process, and phase precipitation driving force.

[0021] As one embodiment of the present invention, in step S2, the nucleation theory includes at least one of the classical spherical cap model nucleation theory and the epitaxial nucleation theory based on the ordered stacking of interface atoms.

[0022] In some embodiments, the concentration step size for both M and S is 0.1 wt.%.

[0023] In one embodiment of the present invention, in step S3, the crystal nucleus density diagram is L12, D0. 22 D0 23 Distribution of phase nucleus density in the alloy composition-solidification cooling rate space.

[0024] Furthermore, the nucleus density map includes a compositional spatial distribution map at the maximum nucleus density and a cooling rate distribution map at the maximum nucleus density.

[0025] Furthermore, in step S4, based on the composition space distribution diagram at the maximum nucleation density drawn in step S3, the alloy composition ratio of the Al-Ti-BMS system at the optimal nucleation density is determined.

[0026] In some embodiments, the Al-Ti-BMS alloy composition includes: 0.7 wt.% Sc and 0.2 wt.% Zr; the nucleus density reaches 10. 24 mm -3 .

[0027] As one embodiment of the present invention, in step S5, the optimal cooling rate under the optimal nucleus density is determined based on the cooling rate distribution diagram under the optimal nucleus density drawn in step S3.

[0028] In some embodiments, the optimal cooling rate is 1000°C / s.

[0029] As one embodiment of the present invention, in step S5, the solidification forming process includes at least one of gravity casting, continuous casting, die casting, horizontal casting, and additive manufacturing.

[0030] As one embodiment of the present invention, in step S5, the solidification forming process adopts additive manufacturing, and the solidification forming process parameters include: deposition power of 1200-1600W, scanning rate of 30-60mm / s, and deposition layer thickness of 0.2-0.4mm.

[0031] In some embodiments, the solidification forming process parameters include: power 1200W, scanning rate 40mm / s, and deposition layer thickness 0.3mm.

[0032] Furthermore, the deposition material used in the additive manufacturing is an Al-Ti-BMS system alloy powder.

[0033] As one embodiment of the present invention, the particle size of the Al-Ti-BMS system alloy powder is 53-105 micrometers.

[0034] As one embodiment of the present invention, in step S6, the Al-Ti-BMS alloy powder comprises the following components by weight percentage:

[0035] 0 < Ti ≤ 7%,

[0036] 0 <B≤3.2%,

[0037] 0 <Mg≤6%,

[0038] 0.4 ≤ Sc ≤ 0.8%

[0039] 0.1 ≤ Zr ≤ 0.4%

[0040] Al: Balance.

[0041] In some embodiments, the Al-Ti-BMS alloy powder comprises the following components by weight percentage: 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr.

[0042] As one embodiment of the present invention, the preparation method of Al-Ti-BMS system alloy powder includes: A. melting raw materials to obtain Al-Ti-BMS system alloy preforms; B. preparing Al-Ti-BMS system alloy powder by gas atomization powder preparation method from the alloy preforms.

[0043] As one embodiment of the present invention, step A specifically includes the following steps:

[0044] A1. Weigh out Al blocks, Mg blocks, Al-TiB2 master alloy blocks, Al-Sc master alloy blocks, and Al-Zr master alloy blocks as raw materials according to the alloy composition ratio;

[0045] A2. Heat and melt Al block, Mg block, and Al-TiB2 intermediate alloy block, and stir to obtain melt A;

[0046] A3. Add Al-Sc master alloy block and Al-Zr master alloy block to melt A, heat and melt, and stir to obtain melt B;

[0047] A4. Add refining agent to melt B and degas under vacuum to obtain melt C;

[0048] A4. After removing the slag from the melt C, it is poured into a preheated mold to obtain an alloy preform.

[0049] In one embodiment of the present invention, in step A2 or A3, the melting temperature is 700-800℃ and the stirring time is 1-5 min.

[0050] In one embodiment of the present invention, in step A4, the refining agent includes FOSCO COVERAL GR2510.

[0051] As one embodiment of the present invention, step B specifically includes the following steps:

[0052] B1. The alloy preform obtained in step A is placed in a vacuum environment and heated to melt, resulting in a molten melt;

[0053] B2. The Al-Ti-BMS system alloy powder is prepared by gas atomization of the molten material obtained in step B1.

[0054] In step B2 of this invention, the molten melt flows under the action of gravity, and the outflowing melt is broken into droplets of different sizes under the impact of atomized nitrogen gas. The droplets solidify into powder during the falling process, and the titanium diboride reinforced aluminum magnesium scandium zirconium alloy powder is collected, namely Al-Ti-BMS system alloy powder.

[0055] In one embodiment of the present invention, in step B1, the heating is electromagnetic induction heating, the heating temperature is 850-950℃, and the holding time is 0.5-1h.

[0056] Secondly, the present invention provides an Al-Ti-BMS system alloy material prepared by the TiB2 / Al interface modification method described above, wherein the Al-Ti-BMS system alloy material comprises 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material.

[0057] As one embodiment of the present invention, the Al-Ti-BMS system alloy material further includes Al-Nb-Ti-B rapid cooling foil strip, which comprises the following components by weight percentage:

[0058] 0 < Ti ≤ 7%,

[0059] 0 <B≤3.2%

[0060] 0.1 ≤ Nb ≤ 1.5%

[0061] Al: Balance.

[0062] In one embodiment of the present invention, the thickness of the foil strip is 0.1–0.3 mm. The TiB2 particles inside the foil strip are uniformly dispersed, with no large Nb-rich phases.

[0063] In this invention, the Al-Nb-Ti-B foil strip is prepared by selecting horizontal casting as the solidification forming process in step S5. The parameters of the horizontal casting include: the linear velocity of the copper roller surface is 10-15 m / s, the distance between the copper roller surface and the gate is 5-8 mm, and the gate size is 2.5 × 15 mm; the foil strip thickness is 0.2-0.3 mm; and the cooling rate is estimated to be ≥10 m / s using a heat transfer theory model. 4 ℃ / s.

[0064] Compared with the prior art, the present invention has the following beneficial effects:

[0065] (1) Compared with traditional TiB2 reinforced aluminum alloy materials, the present invention introduces an additional L12Al3M phase that is highly similar to aluminum crystals and controls the cooling rate to uniformly modify the surface of TiB2 particles during solidification, forming a low mismatch “TiB2 / L12Al3M / α-Al” interface. This successfully reduces the interface mismatch between TiB2 particles and aluminum matrix from 4.2% to below 1%, significantly improving the solidification fine grain efficiency and mechanical strengthening effect of TiB2 particles.

[0066] (2) This invention couples phase diagram thermodynamic calculation and nucleation theory calculation, and develops a high-throughput calculation method for evaluating the number density of blanks. It can accurately predict the changing trend of L12 Al3M modified TiB2 particles in the "composition-cooling rate" space, providing scientific guidance for composition design and cooling rate control, and significantly accelerating the development process of new alloys and new solidification forming processes.

[0067] (3) Using the titanium diboride-reinforced aluminum-magnesium-scandium-zirconium alloy powder developed in this invention as raw material, and employing the laser-directed energy deposition process optimized in this invention, the prepared aluminum alloy bulk material exhibits excellent room-temperature mechanical properties. Its deposited yield strength is 257 MPa, which is the highest level among similar materials to date. Simultaneously, this aluminum alloy bulk material maintains good plasticity, with an elongation of 13.8%.

[0068] (4) The titanium diboride reinforced aluminum-magnesium-scandium-zirconium alloy powder and the matching directional energy deposition process of the present invention have outstanding technical advantages in the manufacturing, surface cladding and repair of large-size complex components, and can be widely used in key fields such as aerospace, shipbuilding and marine, and national defense.

[0069] (5) The method for preparing titanium diboride reinforced aluminum-magnesium-scandium-zirconium alloy powder of the present invention is simple and mature, low-cost and high-efficiency, and can realize large-scale industrial production. Attached Figure Description

[0070] Other features, objects, and advantages of the present invention will become more apparent from the following detailed description of non-limiting embodiments with reference to the accompanying drawings:

[0071] Figure 1 The diagram shows the aluminum-rich angle phase diagram of the Al-Mg-Sc-Zr-Ti-B system at 600℃ in Example 1, and the distribution of the driving force for liquid phase precipitation of the primary Al3(Sc,Zr) phase in the composition space.

[0072] Figure 2 This is a distribution diagram of the liquid phase dissolution temperature of the primary Al3(Sc,Zr) phase in the Al-Mg-Sc-Zr-Ti-B system in Example 1 in the composition space.

[0073] Figure 3 The nucleation rate of Al3(Sc,Zr) primary phase on the surface of TiB2 particles and the trend of critical crown nucleus height with temperature in Example 1 are shown.

[0074] Figure 4 The trend of Al3(Sc,Zr) nucleus number density as a function of cooling rate in the 4% TiB2 / Al-4.5Mg-0.7Sc-0.2Zr system in Example 1;

[0075] Figure 5This is a distribution diagram of the maximum nucleus density in composition space for the 4% TiB2 / Al-4.5Mg-xSc-yZr system in Example 1;

[0076] Figure 6 This is a diagram showing the cooling rate distribution at the maximum nucleus density in Example 1;

[0077] Figure 7 The macroscopic morphology of the 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material in the laser-directed energy deposition state in Example 1 is shown.

[0078] Figure 8 The grain structure of the 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material in the laser-directed energy deposition state is shown in Example 1.

[0079] Figure 9 The distribution of TiB2 ceramic particles inside the laser-directed energy deposited 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material in Example 1;

[0080] Figure 10 This refers to the interface state between TiB2 particles and the aluminum matrix inside the laser-directed energy deposited 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material in Example 1.

[0081] Figure 11 The tensile mechanical properties curve of the laser-directed energy deposition 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material in Example 1 is shown.

[0082] Figure 12 This is a comparison of the performance of the directional energy deposition rapidly solidified bulk materials in Example 1;

[0083] Figure 13 The particle morphology of the Al-4.36Mg-0.72Sc-0.22Zr-2.33Ti-1.23B powder in Example 2 is shown.

[0084] Figure 14 The microstructure of the Al-4.36Mg-0.72Sc-0.22Zr-2.33Ti-1.23B powder in Example 2;

[0085] Figure 15 The internal grain structure of the laser-directed energy deposited 4TiB2 / Al-6Mg bulk material in Comparative Example 1 is shown.

[0086] Figure 16 The distribution of TiB2 ceramic particles inside the laser-directed energy deposited 4TiB2 / Al-6Mg bulk material in Comparative Example 1;

[0087] Figure 17 The tensile mechanical properties of the laser-directed energy deposited 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material in Comparative Example 1 are shown.

[0088] Figure 18 The internal grain structure of the laser-directed energy deposited Al-6Mg-0.6Sc-0.3Zr bulk material in Comparative Example 2 is shown.

[0089] Figure 19 The microstructure of the laser-directed energy deposited Al-6Mg-0.6Sc-0.3Zr bulk material in Comparative Example 2 is shown.

[0090] Figure 20 The tensile mechanical properties curves of the laser-directed energy deposition 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material in Comparative Example 2 are shown.

[0091] Figure 21 The microstructure of the Al-10TiB2-1Nb ingot cooled by graphite mold in Example 3;

[0092] Figure 22 The microstructure of the Al-10TiB2-1Nb foil strip cooled by advection casting in Example 3;

[0093] Figure 23 The Nb-rich primary phase on the surface of TiB2 particles in the Al-10TiB2-1Nb foil strip cooled by advection casting in Example 3;

[0094] Figure 24 This is a schematic diagram of the Al3(Sc,Zr), TiB2, and α-Al interface structure of the 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material in Example 1. Detailed Implementation

[0095] The present invention will now be described in detail with reference to the accompanying drawings and specific embodiments. The following examples are implemented under the premise of the technical solution of the present invention, providing detailed implementation methods and specific operating procedures, which will help those skilled in the art to further understand the present invention. It should be noted that the scope of protection of the present invention is not limited to the following embodiments; any adjustments and improvements made under the concept of the present invention are all within the scope of protection of the present invention.

[0096] It should be noted that the phase diagram thermodynamic calculation software used in the embodiments and comparative examples of this invention is Pandat, and the refining agent is FOSCO COVERAL GR2510.

[0097] Example 1

[0098] This embodiment uses the TiB2 / Al-Mg-Sc-Zr system as the calculation object to illustrate the titanium diboride / aluminum interface modification method proposed in this invention based on the synergistic regulation of alloy composition and solidification cooling rate.

[0099] Follow these steps:

[0100] (a) Phase diagram thermodynamic calculations

[0101] Using phase diagram thermodynamics calculation software and the Al-Mg-Sc-Zr-Ti-B system phase diagram thermodynamics database, the driving force for liquid phase precipitation and the liquid phase dissolution temperature of the primary Al3(Sc,Zr) phase were calculated as a function of Sc and Zr element concentrations. Figure 1 and Figure 2 As shown. According to the heterogeneous nucleation theory, when the types and number of nucleation sites are constant, the nucleation rate of nascent Al3(Sc,Zr) is closely related to the nucleation work, which is determined by the driving force of liquid-phase precipitation. Therefore, the two necessary conditions for generally obtaining the "TiB2 / Al3(Sc,Zr) / Al" composite interface are that Al3(Sc,Zr) is completely dissolved before cooling, and the driving force for Al3(Sc,Zr) precipitation is sufficiently large during cooling.

[0102] (b) Nucleation theory calculations

[0103] The cooling rate affects the liquid precipitation behavior of the Al3(Sc,Zr) primary phase on the surface of TiB2 particles, thereby altering the Al3(Sc,Zr) coating state and the ease with which the α-Al solidification interface engulfs TiB2 particles. The quantity and extent of Al3(Sc,Zr) coating on TiB2 are closely related to the heterogeneous nucleation process. According to the classical spherical cap model heterogeneous nucleation theory, the nucleation rate of the Al3(Sc,Zr) primary phase on the surface of TiB2 particles is:

[0104]

[0105] Where N0 is the number density of TiB2 particles, D is the diffusion coefficient of M element in Al melt, a is the liquid / solid interatomic distance, and x l,eff γ is the concentration of atoms (or groups) in the solidified phase, θ is the wetting angle, f(θ) is the angle factor, and γ S / L ΔG is the interfacial energy between the solidified phase and the molten aluminum. * For nucleation work (for Where ΔG V τ is the driving force for the liquid phase precipitation of Al3M, τ is the nucleation and incubation time, T is the local melt temperature, k is the Boltzmann constant, and t is the time it takes for the melt to supercool from the liquidus to T (i.e., in (Refers to the cooling rate). Taking the 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr system as an example, substituting the above physical property parameters, the change in nucleation rate with temperature is as follows: Figure 3 As shown.

[0106] Depend on Figure 3 It can be seen that the nucleation rate first increases and then decreases with decreasing melt temperature, reaching a maximum value around 527℃. However, in reality, heterogeneous nucleation stops when the temperature drops to 733℃. This is because at this temperature, the height of the critical crown nucleus is already less than the interatomic spacing of one layer, allowing Al and Sc atoms to directly attach to the TiB2 substrate or the already formed Al3(Sc,Zr) crystal, no longer generating three-dimensional nuclei. Therefore, Al3(Sc,Zr) nuclei only form between 733 and 890℃, and their number density is determined by the following formula:

[0107]

[0108] Where T L T is the liquidus temperature (890℃). c The nucleation termination temperature, with values ​​ranging from 733 to 890 °C, N n for When the phase fraction precipitated during nucleation reaches a certain value, the subsequent precipitation process will be dominated by growth. The Al3(Sc,Zr) fraction precipitated during the nucleation-growth transition is calculated by the following formula:

[0109]

[0110] Where n total n* represents the total number of atoms (including Al and M) that may participate in the precipitation of Al3M during the entire solidification process; n* represents the number of atoms in the crown-shaped nucleus. total It can be obtained from the following formula:

[0111]

[0112] in ρ represents the total mass fraction of Al3(Sc,Zr) that the system can form after equilibrium solidification. L The density of the melt is n*. This is obtained from the following formula:

[0113]

[0114] From equation (3), we can see that Conversely, this determined T c Therefore, the number of crystal nuclei that can be obtained can be determined by (2). The numbers are intrinsic parameters of the system and need to be calibrated in conjunction with experimental results.

[0115] Figure 4 The calculated Al3(Sc,Zr) nucleus number density as a function of cooling rate is given, where different curves correspond to... The range is 0.5-4.5%. This shows that as the cooling rate increases, the nucleus number density first increases and then decreases, indicating that both excessively slow and excessively fast cooling rates inhibit nucleation. In this case, the Al3(Sc,Zr) nucleus number density is the highest when the cooling rate reaches ~1000℃ / s. Correspondingly, the number of TiB2 particles modified by Al3(Sc,Zr) is the highest.

[0116] (c) Draw the nucleus density diagram

[0117] With the concentration step size of Sc and Zr set to 0.1 wt.%, nucleation theory calculations as described in step (b) were performed on the 4% TiB2 / Al-4.5Mg-xSc-yZr system to obtain the distribution map of the maximum nucleus density in composition space, as shown below. Figure 5 As shown.

[0118] (d) Determine the alloy composition

[0119] Based on the nucleus density diagram, the content of Sc was selected as 0.7 wt.% and the content of Zr as 0.2 wt.%, so that the nucleus number density reached 10. 24 mm -3 Based on this, an alloy composition with a large Al3(Sc,Zr) precipitation driving force and a low liquid phase dissolution temperature was designed: Al-4.5Mg-0.7Sc-0.2Zr-2.72Ti-1.28B (wt.%).

[0120] (f) Develop solidification forming process

[0121] Accordingly, the cooling rate distribution at the maximum nucleus density is obtained, as shown in the figure. Figure 6 As shown, the optimal cooling rate for 0.7 wt.% Sc and 0.2 wt.% Zr was determined to be 1000 °C / s. Using conventional multiphysics calculation methods and commercial software, the temperature field of the laser-fed printing melt pool was simulated under pre-defined printing process parameters to obtain the cooling rate. This rate was then compared with 1000 °C / s, and the printing process parameters were adjusted based on feedback until a combination of printing process parameters that met the cooling rate requirements was designed. The determined printing process combination was: 1200 W power, 40 mm / s scan rate, and 0.3 mm deposition layer thickness.

[0122] (g) Preparation and shaping

[0123] Using 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr powder as raw material, with a particle size range of 53-105 micrometers, a conventional powder-feeding laser-directed energy deposition (EDD) system was employed. The printing process parameters were set as follows: power 1200W, scanning rate 40mm / s, and deposition layer thickness 0.3mm, to prepare 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material. Its Al3(Sc,Zr), TiB2, and α-Al interface structure is as follows... Figure 24 As shown.

[0124] Effects of Example 1:

[0125] (1) Macroscopic Form

[0126] Laser-directed energy deposition (LDED) of 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk materials exhibits high forming quality, with regular shapes, clean surfaces, and no obvious defects. Figure 7 As shown.

[0127] (2) Microstructure

[0128] The laser-directed energy deposited 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material exhibits a uniform internal grain structure with fine grain size, reaching 4.50 micrometers. Figure 8 As shown. Meanwhile, in this embodiment, the TiB2 particles inside the bulk material are uniformly dispersed, as... Figure 9 As shown.

[0129] (3) Interface Structure

[0130] Transmission electron microscopy was used to analyze the interface between TiB2 particles and the aluminum matrix inside a laser-directed energy deposition (LDED) 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material. The results are as follows: Figure 10 As shown. It can be seen that an Al3(Sc,Zr) phase with L12 crystal morphology is formed at the interface, constituting a multi-level interface structure of "TiB2 / L12 Al3(Sc,Zr) / α-Al". Since L12Al3(Sc,Zr) and α-Al have the same lattice structure, the mismatch degree in the (1 1 1)

[110] L12 Al3(Sc,Zr) / (1 1 1)

[110] α-Al orientation is less than 1%. Therefore, the "TiB2 / L12 Al3(Sc,Zr) / α-Al" interface has better matching than the simple "TiB2 / α-Al" interface (lattice mismatch degree is 4.2%), achieving the purpose of interface control and improving the role of TiB2 particles in solidification, fine grains and particle dispersion. For the actual effect, see Figure 8 and Figure 9 .

[0131] (4) Mechanical properties

[0132] According to the conventional tensile test (GB / T 228.1-2021), such as Figure 11 As shown, the laser-directed energy deposited (LDED) 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material exhibits excellent strength and toughness. The yield strength of the LDED 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material reaches 257 MPa, the tensile strength is 385 MPa, and the elongation is 13.8%. After heat treatment at 325℃ for 4 hours, the yield strength reaches 326 MPa, the tensile strength reaches 437 MPa, and the elongation is 11.8%. Using "uniform elongation × (tensile strength × yield strength) / 2" as the strength-toughness matching evaluation index, the value for the LDED 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material is 43.2 mJ / mm². -3 .

[0133] By comparing materials prepared using other directional energy deposition rapid solidification methods, it can be seen that, for example... Figure 12 As shown, the 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material prepared in this embodiment has the best strength and plasticity matching.

[0134] Example 2

[0135] This embodiment illustrates the preparation method of titanium diboride reinforced aluminum-magnesium-scandium-zirconium alloy powder for additive manufacturing.

[0136] The nominal composition of the titanium diboride-reinforced aluminum-magnesium-scandium-zirconium alloy powder was selected as 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr. The preparation steps are as follows:

[0137] (1) Raw material preparation:

[0138] 14.76 kg of Al block, 2.70 kg of Mg block, 20.34 kg of 11.8TiB2-Al master alloy block, 21.00 kg of Al-Sc master alloy block, and 1.20 kg of Al-Zr master alloy block were weighed as raw materials. In addition, according to the standard of 95% recovery rate of Mg, Sc, and Zr, 0.14 kg of Mg block, 1.11 kg of Al-Sc master alloy block, and 0.06 kg of Al-Zr master alloy block were weighed as raw materials to supplement the raw materials lost during burning.

[0139] (2) Smelting of 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr alloy preforms:

[0140] a. Mix Al blocks, Mg blocks, and TiB2-Al master alloy blocks and place them in a graphite crucible. Heat the mixture in a resistance furnace to 750°C until melted, and then stir with a graphite stirring rod for 3 minutes.

[0141] b. Add the Al-Sc master alloy block and Al-Zr master alloy block to the melt and stir with a graphite stirring rod for 3 minutes;

[0142] c. Add the Mg block to the melt and use a graphite rod to press it into the bottom of the melt to dissolve it;

[0143] d. Add FOSCO COVERAL GR2510 for refining, followed by vacuum degassing for 10 minutes;

[0144] e. Remove the surface slag and cast it into a cylindrical mold preheated at 250℃ to obtain a cylindrical ingot;

[0145] f. Remove the oxide scale from the surface of the ingot using mechanical processing methods.

[0146] (3) Gas atomization forming of 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr alloy powder:

[0147] a. Place the 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr preform into a graphite crucible in the melting chamber, close the furnace door, reduce the vacuum level in the melting chamber through the vacuum system, and then introduce nitrogen into the chamber to further replace the air in the chamber and reduce the oxygen content in the chamber.

[0148] b. The cavity is heated by electromagnetic induction to a target temperature of 800℃ and held for 0.5 hours to completely melt the ingot;

[0149] c. The molten melt flows along the nozzle under the action of gravity and is broken into droplets of different sizes under the impact of the rapidly moving atomized nitrogen gas. The droplets solidify into powder as they fall, and the falling powder is collected at the bottom of the cavity.

[0150] d. Vacuum pack the collected powder to prevent it from oxidizing.

[0151] Effects of Example 2

[0152] The composition of the 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr powder prepared in this embodiment was obtained by inductively coupled plasma atomic emission spectrometry as Al-4.36Mg-0.72Sc-0.22Zr-2.33Ti-1.23B, which is consistent with the nominal composition of the design.

[0153] The typical particle morphology and typical microstructure of the Al-4.36Mg-0.72Sc-0.22Zr-2.33Ti-1.23B powder prepared in this embodiment are as follows: Figure 13 and Figure 14 As shown, the powder exhibits high sphericity and contains primary phases of α-Al, TiB2, and L12Al3(Sc,Zr).

[0154] Comparative Example 1

[0155] Comparative Example 1 is a comparative example of Example 1.

[0156] The powder raw material was commercially available 4TiB2 / Al-6Mg alloy powder with a particle size range of 53-105 micrometers. 4TiB2 / Al-6Mg bulk material was prepared using the same powder-feeding laser-directed energy deposition equipment and processing parameters as in Example 1.

[0157] Comparative Example 1 Effect

[0158] (1)Alloy composition

[0159] The chemical composition of the laser-directed energy deposited 4TiB2 / Al-6Mg bulk material was determined to be Al-4.51Mg-2.05Ti-1.20B (wt.%) using conventional inductively coupled plasma atomic emission spectrometry.

[0160] (2) Microstructure

[0161] The grain size of the laser-directed energy deposited 4TiB2 / Al-6Mg bulk material exhibits some fluctuation, with an average grain size of 16.00 micrometers. Figure 15 As shown. Meanwhile, the TiB2 particles inside exhibit significant agglomeration at the grain boundaries, such as... Figure 16 As shown. Since this material does not contain Sc or Zr, the "TiB2 / L12 Al3(Sc,Zr) / α-Al" multi-level interface structure was not obtained.

[0162] (3) Mechanical properties

[0163] The yield strength of the laser-directed energy deposited 4TiB2 / Al-6Mg bulk material is 126 MPa, the tensile strength is 291 MPa, and the elongation is 19.6%. Figure 17 As shown, the strength and toughness evaluation index of the laser-directed energy deposited 4TiB2 / Al-6Mg bulk material is 38.1 mJ / mm². -3 .

[0164] The comparison shows that the microstructure and mechanical properties of the laser-directed energy deposited 4TiB2 / Al-6Mg bulk material are weaker than those of the laser-directed energy deposited 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material in Example 1.

[0165] Comparative Example 2

[0166] Comparative Example 2 is a comparative example of Example 1.

[0167] The powder raw material was commercially available Al-6Mg-0.6Sc-0.3Zr alloy powder with a particle size range of 53-105 micrometers. Al-6Mg-0.6Sc-0.3Zr bulk material was prepared using the same powder-feeding laser-directed energy deposition equipment and processing parameters as in Example 1.

[0168] Comparative Example 2 Effect

[0169] (1) Microstructure

[0170] Analysis using conventional inductively coupled plasma atomic emission spectrometry (ICP-AES) revealed the chemical composition of the laser-directed energy deposited Al-6Mg-0.6Sc-0.3Zr bulk material to be Al-4.21Mg-0.56Sc-0.28Zr (wt.%).

[0171] (2) Microstructure

[0172] The bulk Al-6Mg-0.6Sc-0.3Zr material, obtained through laser-directed energy deposition, consists of columnar and equiaxed crystals, with significant fluctuations in grain size. The average grain size of the equiaxed crystals is 5.49 micrometers. Figure 18 As shown. Furthermore, the material contains a large number of Al3(Sc,Zr) primary phases with dimensions smaller than 1 micrometer, such as... Figure 19 As shown, their distribution is uneven. At the bottom of the molten pool, due to the slow cooling rate, the number density of Al3(Sc,Zr) primary phase particles is higher; near the middle and top of the molten pool, due to the faster cooling rate, the number density of Al3(Sc,Zr) primary phase particles is lower. Since this bulk material does not contain TiB2 particles, a multi-level interface structure of "TiB2 / L12 Al3(Sc,Zr) / α-Al" is not formed.

[0173] (3) Mechanical properties

[0174] The testing method is the same as in Example 1. The yield strength of the laser-directed energy deposited Al-6Mg-0.6Sc-0.3Zr bulk material is 194 MPa, the tensile strength is 318 MPa, and the elongation is 15.2%. Figure 20 As shown. The strength and toughness evaluation index of this bulk material is 38.9 mJ / mm. -3 After heat treatment, the yield strength of this bulk material is 310 MPa, the tensile strength is 403 MPa, and the elongation is 14.4%.

[0175] The comparison shows that the microstructure uniformity and mechanical properties of the laser-directed energy deposited Al-6Mg-0.6Sc-0.3Zr bulk material are weaker than those of the laser-directed energy deposited 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material in Example 1.

[0176] Example 3

[0177] This embodiment uses the Al-10TiB2-1Nb system as an example to illustrate the titanium diboride / aluminum interface modification method proposed in this invention based on the synergistic regulation of alloy composition and solidification cooling rate.

[0178] Follow these steps:

[0179] (1) Raw material preparation:

[0180] 1.53 kg of Al block, 25.41 kg of 11.8TiB2-Al master alloy block, and 3 kg of Al-20Nb master alloy block were weighed as raw materials, totaling 30 kg;

[0181] (2) Smelting of Al-10TiB2-1Nb alloy ingots:

[0182] a. Mix the Al block and TiB2-Al intermediate alloy block and place them in a graphite crucible. Heat the mixture in a resistance furnace to 750°C to melt it, and then stir it with a graphite stirring rod for 3 minutes.

[0183] b. Add the Al-20Nb intermediate alloy block to the melt and stir with a graphite stirring rod for 3 minutes;

[0184] c. Add refining agent for refining, then degas under vacuum for 10 minutes and remove surface slag;

[0185] e. After reheating to 900-950℃, 10kg of melt is poured into a square graphite mold preheated to 250℃, resulting in a sample thickness of 8mm and a corresponding cooling rate of ~50℃ / s.

[0186] d. Another 10 kg of melt was used to prepare foil strips using a horizontal casting method (i.e., strip casting). The process parameters were: linear velocity of the copper roller surface of 12 m / s, distance between the copper roller surface and the gate of 5–8 mm, and gate size of 2.5 × 15 mm. The resulting foil strip thickness was 0.2–0.3 mm. Using a heat transfer theory model, the cooling rate was estimated to be ≥10⁻⁶ m / s. 4 ℃ / s.

[0187] Effects of Example 3

[0188] The composition of the Al-10TiB2-1Nb material prepared in this embodiment was determined to be Al-6.49Ti-3.11B-0.96Nb using inductively coupled plasma atomic emission spectrometry. The microstructures of the two samples prepared in this embodiment are as follows:

[0189] (1) In the samples prepared by graphite molds, Nb exists alone in the form of coarse, flake-like Al3Nb, while TiB2 particles are severely agglomerated, such as Figure 21 As shown, Al3Nb growth at this time was mainly characterized by spontaneous homogeneous nucleation, and the multi-level interface structure of "TiB2 / Al3Nb / α-Al" was not generally formed.

[0190] (2) According to Figure 22 The microstructure of the Al-10TiB2-1Nb foil strip shown is obtained by advection casting cooling. Figure 23 As shown in the diagram, the Nb-rich primary phase on the surface of TiB2 particles in the Al-10TiB2-1Nb foil strip prepared by advection casting indicates that no large Nb-rich phases were observed under scanning electron microscopy for the foil strip sample prepared by advection casting, and the particle dispersion was significantly improved. Transmission electron microscopy and energy dispersive spectroscopy analysis showed that there was an approximately 5 nm Nb-rich layer between the Al matrix and TiB2, indicating that Al3Nb nucleates and grows on the TiB2 particle substrate, completing the coating and forming a beneficial "TiB2 / Al3Nb / α-Al" multi-level interface structure, achieving the design control objective.

[0191] The specific embodiments of the present invention have been described above. It should be understood that the present invention is not limited to the specific embodiments described above, and those skilled in the art can make various modifications or variations within the scope of the claims, which do not affect the essence of the present invention.

Claims

1. A method for TiB2 / Al interface modification based on synergistic regulation of alloy composition and solidification cooling rate, characterized in that, The modification method includes the following steps: S1. Phase diagram thermodynamic calculation: The phase diagram thermodynamic results of the Al-Ti-BMS system are calculated using the CALPHAD method. In step S1, M includes one or more of Ti, Sc, Zr, V, Nb, Ta, Hf, Y, Yb, Er, Cr, W, and Li, and S includes one or more of Mg, Zn, Cu, Mn, and Si. The CALPHAD method includes: performing calculations using phase diagram thermodynamics calculation software and the Al-Ti-BMS system phase diagram thermodynamics database; S2. Nucleation Theory Calculation: Based on the phase diagram thermodynamic results of the Al-Ti-BMS system obtained in step S1, the nucleation theory is used to calculate the L12 and D0 phases formed in the Al-Ti-BMS system at different temperatures during solidification. 22 D0 23 Phase nucleus density; S3. Draw the nucleus density map: Set the concentration step size of M and S, and draw the nucleus density map based on the nucleus density calculated in step S2. In step S3, the nucleus density map is a distribution map of the nucleus density of L12, D022, and D023 phases in the alloy composition-solidification cooling rate space; the nucleus density map includes a composition space distribution map at the maximum nucleus density and a cooling rate distribution map at the maximum nucleus density. S4. Determine the alloy composition: Based on the nucleus density diagram drawn in step S3, determine the alloy composition ratio of the Al-Ti-BMS system under the optimal nucleus density; S5. Formulate solidification forming process: Based on the crystal nucleus density diagram drawn in step S3, determine the optimal cooling rate under the optimal crystal nucleus density, and then formulate solidification forming process parameters. S6. Preparation and forming: Using the Al-Ti-BMS system alloy powder with the alloy composition ratio determined in step S4 as raw material, Al-Ti-BMS system alloy material is prepared according to the solidification forming process parameters in step S5. The preparation method of Al-Ti-BMS system alloy powder includes: A. melting raw materials to obtain Al-Ti-BMS system alloy preforms; B. preparing Al-Ti-BMS system alloy powder by gas atomization powder preparation method from the alloy preforms.

2. The modification method according to claim 1, characterized in that, In step S1, the phase diagram thermodynamic results include at least one of phase equilibrium relationship, phase evolution during solidification process, and phase precipitation driving force; in step S2, the nucleation theory includes at least one of classical spherical cap model nucleation theory and epitaxial nucleation theory based on interface atomic ordered stacking.

3. The modification method according to claim 1, characterized in that, In step S4, the composition ratio of the Al-Ti-BMS system alloy at the optimal nucleation density is determined based on the composition space distribution map at the maximum nucleation density drawn in step S3; in step S5, the optimal cooling rate at the optimal nucleation density is determined based on the cooling rate distribution map at the maximum nucleation density drawn in step S3.

4. The modification method according to claim 1, characterized in that, In step S5, the solidification forming process adopts additive manufacturing. The solidification forming process parameters include: deposition power of 1200~1600 W, scanning rate of 30~60 mm / s, and deposition layer thickness of 0.2~0.4 mm.

5. The modification method according to claim 1, characterized in that, In step S6, the Al-Ti-BMS alloy powder comprises the following components by weight percentage: 0 < Ti ≤ 7%, 0<B≤3.2%， 0 <Mg≤6%, 0.4≤Sc≤0.8%, 0.1≤Zr≤0.4%, Al: Balance.

6. The modification method according to claim 1, characterized in that, Step A specifically includes the following steps: A1. Weigh Al blocks, Mg blocks, Al-TiB2 master alloy blocks, Al-Sc master alloy blocks, and Al-Zr master alloy blocks respectively according to the alloy component ratio as raw materials; A2. Heat and melt and stir the Al blocks, Mg blocks, and Al-TiB2 master alloy blocks to obtain melt A; A3. Add the Al-Sc master alloy blocks and Al-Zr master alloy blocks into melt A, heat and melt and stir to obtain melt B; A4. Add a refining agent to melt B and degas it under vacuum to obtain melt C; A4. After removing the slag from melt C, pour it into a preheated mold to obtain an alloy preform ingot.

7. The modification method according to claim 1, characterized in that, Step B specifically includes the following steps: B1. Put the alloy preform ingot obtained in step A into a vacuum environment and heat it to melt to obtain a molten melt; B2. Obtain Al-Ti-B-M-S system alloy powder from the molten melt obtained in step B1 by gas atomization powder making method.

8. An Al-Ti-BMS alloy material prepared by the TiB2 / Al interface modification method as described in any one of claims 1-7, characterized in that, The Al-Ti-B-M-S system alloy material includes 4TiB2 / Al-4.5Mg-0.7Sc-0.2Zr bulk material.

9. The Al-Ti-BMS alloy material according to claim 8, characterized in that, The Al-Ti-B-M-S system alloy material also includes Al-Nb-Ti-B rapidly solidified foil strip. The Al-Nb-Ti-B rapidly solidified foil strip includes the following components by weight percentage: 0 < Ti ≤ 7%, 0 < B ≤ 3.2%, 0.1 ≤ Nb ≤ 1.5%, and the balance is Al.

10. The Al-Ti-BMS alloy material according to claim 9, characterized in that, The preparation method of the Al-Nb-Ti-B rapid cooling foil strip includes: in step S5, using horizontal casting as the solidification forming process to prepare the foil strip, the process parameters of which include: the surface linear velocity of the copper roller is 10~15m / s, the distance between the surface of the copper roller and the gate is 5~8 mm, the gate size is 2.5×15 mm; the foil strip thickness is 0.2~0.3 mm; and the cooling rate is estimated to be ≥10 using a heat transfer theory model. 4 ℃ / s.

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