A rare earth Y-doped easily-activated Ti-Fe-based AB type hydrogen storage alloy and a preparation method thereof

Abstract

This invention relates to a rare-earth Y-doped easily activated Ti-Fe-based AB-type hydrogen storage alloy and its preparation method, belonging to the field of solid-state hydrogen storage material technology. Its composition is: (Ti 1‑x‑y Y x Zr y ) 1.15 Fe 1.1‑z‑ m Mn z Ni m In the formula, x, y, z, and m are atomic ratios, and 0.01≤x≤0.04, 0.02≤y≤0.1, 0.1≤z≤0.3, and 0.05≤m≤0.12. The hydrogen storage alloy exhibits good activation performance, hydrogen absorption and desorption kinetics, and cycle stability. The hydrogen storage capacity is ≥1.74 wt.%; it can be fully activated in a single charge-discharge cycle at 30℃ and an initial hydrogen pressure of 3 MPa. It exhibits good cycle stability; after 200 charge-discharge cycles, the capacity retention rate is ≥98.9%, and it retains its activation performance even after 72 hours of exposure to air. This preparation method involves obtaining alloy strips through vacuum induction melting and rapid quenching under an argon atmosphere. The rapidly quenched strips are then mechanically crushed and ball-milled to obtain alloy powder with a nanocrystalline structure. This method solves the problems of high activation conditions, long activation cycles, and high energy consumption in batch preparation processes of Ti-Fe based hydrogen storage alloys prepared by traditional processes.

Description

Technical Field

[0001] This invention belongs to the field of solid-state hydrogen storage alloy materials technology, and in particular provides a rare-earth Y-doped easily activated Ti-Fe-based AB-type hydrogen storage alloy and its preparation method. Background Technology

[0002] Hydrogen is a clean energy source with no carbon dioxide emissions, produced by splitting water. As the most abundant element on Earth, hydrogen has attracted increasing attention. The application of hydrogen energy is considered a crucial strategy for achieving the "dual carbon" goal, and the Chinese government has listed it as a new priority development direction and a new economic growth point during the 15th Five-Year Plan period. The challenges of hydrogen energy application lie in its multi-stage nature, encompassing hydrogen production, storage, transportation, refueling, and application. These stages are interconnected and mutually restrictive, each with its own characteristics. The intertwining of various technical issues creates significant technological bottlenecks for hydrogen energy application.

[0003] In the various stages of hydrogen energy applications mentioned above, hydrogen production has several relatively mature solutions, including obtaining green hydrogen through water electrolysis, blue hydrogen as a chemical byproduct, and gray hydrogen derived from coal and petrochemicals. However, there is still no satisfactory solution for efficient and safe hydrogen storage. Hydrogen storage includes storage in the form of high-pressure gas, cryogenic liquid, or solid hydrides. High-pressure gaseous hydrogen storage is currently the most mature method, with its main advantage being convenient hydrogen filling and releasing. However, its volumetric hydrogen storage density is very low. The most effective method is to increase the hydrogen storage pressure to increase the hydrogen storage density, but this inevitably increases compression energy consumption and places extremely high demands on the pressure resistance of the hydrogen storage container, leading to a sharp increase in cost. Cryogenic liquid hydrogen storage is a method of storing liquid hydrogen in an insulated container, characterized by a high volumetric hydrogen storage density (greater than 50 kg H2 / m³). 3 However, its liquefaction process is energy-intensive, requires extremely high insulation performance of the container, and poses safety hazards such as hydrogen evaporation and leakage. Unlike gaseous and liquid hydrogen storage, solid-state hydride hydrogen storage methods have high volumetric density (greater than 50 kg H2 / m³). 3 It offers advantages such as safety and reliability. Furthermore, solid hydride hydrogen storage operates at lower pressures (less than 10 MPa) and produces hydrogen with higher purity. Therefore, solid hydride hydrogen storage has unique advantages in hydrogen energy applications.

[0004] Solid-state hydride hydrogen storage technology relies primarily on the hydrogen absorption and desorption performance of the storage materials. Therefore, the research and development of high-performance hydrogen storage materials is a major focus of solid-state hydrogen storage technology development. After decades of effort, researchers have developed various hydrogen storage materials, including hydrogen storage alloys, composite metal hydrides (metal propionates, metal borohydrides, metal nitrides, etc.), chemical hydrides, and light metal hydrides (magnesium hydrides, aluminum hydrides, etc.). Among these proven hydrogen storage materials, hydrogen storage alloys have received more attention due to their unique advantages and have already achieved large-scale applications in practice. There are many types of hydrogen storage alloys, and those that have been extensively studied include rare-earth-based AB5 type, Ti-Mn-based AB2 type, Ti-Fe-based AB type, Mg-Ni-based A2B type, magnesium-based alloys, and vanadium-based solid solution alloys. Among these, rare-earth-based AB5 type alloys and Ti-Mn-based AB2 type alloys have achieved industrial applications. Ti-Fe-based alloys have attracted attention due to their high hydrogen storage capacity, excellent room-temperature hydrogen absorption and desorption kinetics, and cycle stability. In particular, its low raw material cost (only 1 / 3 that of rare earth-based AB5 alloys) is considered highly competitive in the field of solid-state hydrogen storage. Unfortunately, compared to other alloys, Ti-Fe-based AB alloys are more difficult to activate, requiring multiple cycles of hydrogen absorption and desorption at 450 °C and 6 MPa H₂ for activation, which is unacceptable for large-scale practical applications. Therefore, overcoming the final hurdle for the practical application of this alloy is improving its room-temperature activation capability.

[0005] Studies have confirmed that the main reason for the difficulty in activating Ti-Fe based AB-type alloys is the rapid formation of a dense TiO2 oxide film on the surface of the alloy upon contact with air during preparation. This oxide film blocks hydrogen from contacting the alloy, preventing hydrogen absorption. Many studies have focused on alloy design to eliminate the thermal activation process and / or improve the initial hydrogen absorption performance of Ti-Fe based hydrogen storage alloys. Research shows that Ti-Fe alloys with added Zr that do not require thermal activation can absorb hydrogen, with the Zr-rich phase acting as a hydrogen absorption channel. Some scholars believe that the initial hydrogen absorption behavior is significantly influenced by the chemical composition of the alloy surface oxide layer. Replacing Ti with RE (La, Ce, Pr, Nd, Sm, Y, etc.) and Hf, Zr, Mg, Ca, V metals, and partially replacing Fe with transition metals such as Mn, Cu, Ni, Cr, Co, Al, etc., can alter the phase structure of the alloy matrix and the composition and structure of the surface oxide layer, thereby improving the alloy's hydrogen storage performance, especially its activation performance.

[0006] Besides element substitution, the preparation process also significantly affects Ti-Fe based AB-type alloys. In particular, rapid quenching, ball milling, and annealing processes can significantly improve the overall hydrogen storage performance of the alloys. The particle and grain size of TiFe alloys are strongly related to their activation performance. The smaller the particle size and grain size, the easier the activation, the lower the required hydrogen pressure and temperature, and the shorter the incubation period. The microcrystals, nanocrystals, grain boundaries, and internal defects generated by ball milling and rapid quenching are considered to be preferential nucleation sites for TiFe hydrides during hydrogen absorption. This is because grain boundaries or defects have higher storage energy, providing a driving force for the nucleation of TiFe hydrides.

[0007] This invention employs an A-side overstoichiometry in its composition design, using rare earth element Y and transition metals Mn, Zr, and Ni for alloying. In terms of preparation technology, it utilizes an integrated vacuum induction melting-rapid quenching technique to avoid compositional segregation during solidification and obtain a nanocrystalline microstructure. The rapidly quenched alloy is then subjected to short-term ball milling, altering its surface state. Testing and analysis of the alloy's microstructure, activation properties, thermodynamics, and kinetics revealed that the alloy of this invention possesses excellent activation properties and retains these properties even after prolonged exposure to air, fully meeting the requirements of practical applications. Summary of the Invention

[0008] Based on the above analysis, the present invention aims to provide a rare earth Y-doped easily activated Ti-Fe-based AB-type hydrogen storage alloy and its preparation method, so as to solve the problems of high activation conditions and poor product performance consistency of existing Ti-Fe-based AB-type hydrogen storage alloys.

[0009] On one hand, this invention provides a rare-earth Y-doped, easily activated Ti-Fe-based AB-type hydrogen storage alloy, specifically a Ti-Zr-Y-Fe-Mn-Ni AB-type hydrogen storage alloy, wherein the chemical formula of the hydrogen storage alloy is: (Ti 1-x-y Y x Zr y ) 1.15 Fe 1.1-z-m Mn z Ni m In the formula, x, y, z, and m are atomic ratios, and 0.01≤x≤0.04, 0.02≤y≤0.1, 0.1≤z≤0.3, and 0.05≤m≤0.12.

[0010] A further technical solution of the present invention: the preferred ratio of x, y, z, m is x:y:z:m = 0.02:0.04:0.22:0.08.

[0011] In this invention, the hydrogen storage alloy comprises a multiphase structure containing a main phase TiFe phase (CsCl structure) and a second phase C14 Laves phase (MgZn2 structure), which respectively have cubic CsCl type and MgZn2 type structures.

[0012] A further technical solution of the present invention: the hydrogen storage alloy has a nanocrystalline structure with an average grain size of 30-100 nm.

[0013] In this invention, the hydrogen storage alloy can be fully activated in one activation at 30 °C and 3 MPa hydrogen pressure, with a hydrogen absorption plateau pressure ≥0.43 MPa and a hydrogen release plateau pressure ≥0.36 MPa. The reversible hydrogen absorption capacity is ≥1.74 wt.%, and the capacity retention rate is ≥98.9% after 200 hydrogen absorption / desorption cycles.

[0014] The activated hydrogen storage alloy remains activated even after being exposed to air for 72 hours.

[0015] On the other hand, the present invention also provides a method for preparing a rare earth Y-doped easily activated Ti-Fe-based AB-type hydrogen storage alloy, which is used to prepare the above-mentioned easily activated Ti-Fe-based AB-type hydrogen storage alloy, specifically including the following steps: S1: The dosage is calculated according to the chemical formula composition and the purity of the raw material metal is ≥99.5%, of which Mn and rare earth Y elements are added with appropriate burn-off amount; S2: Place the prepared raw materials in an Al2O3 crucible, cover the furnace, evacuate to the preset vacuum level, and then fill with high-purity Ar gas at a certain pressure as a protective gas. Use electric arc melting, induction heating melting or other heating methods to melt the raw materials. After holding the temperature for a period of time, a liquid master alloy with uniform composition is obtained. S3: Liquid master alloy is directly injected into an tundish with NB nozzles embedded at the bottom. The liquid master alloy (approximately 1400 ℃) is continuously sprayed from the narrow slits (0.3 mm wide) of the NB nozzles onto the surface of a rotating water-cooled copper roller (surface linear velocity of 3–20 m / s), resulting in a rapidly quenched alloy strip (100–300 μm thick). This method enables rapid solidification of liquid metals or alloys; therefore, in the metallurgical industry, this process is called rapid solidification or rapid quenching, and the surface linear velocity of the rotating copper roller is called the quenching rate.

[0016] S4: After mechanically crushing and sieving the fast-quenched alloy strip, the undersize alloy powder is obtained. It is then loaded into a stainless steel ball mill jar together with stainless steel grinding balls. After vacuuming, high-purity argon gas is introduced, and the mixture is ball-milled in an all-around planetary high-energy ball mill to obtain the ball-milled alloy powder, which is the easily activated Ti-Fe-based AB-type hydrogen storage alloy.

[0017] A further technical solution of the present invention: the amount of burn-off of Mn and Y added in step S1 is 3% to 5% of the calculated dosage.

[0018] Further technical solution of the present invention: The placement and specific operation of each raw material in step S2 are as follows: All raw materials Ti, Zr, Y, Fe, Mn, and Ni are cleaned with a special raw material cleaning device to remove the oxide scale from their surfaces. The crucible is cleaned with high-purity liquid Fe before melting the alloy to remove impurities from the crucible surface; this operation is called furnace cleaning in the metallurgical industry. All raw materials are placed in the Al2O3 crucible in sequence, with rare earth Y placed at the bottom, sponge Ti and Zr placed above rare earth Y, industrial pure Fe and electrolytic Ni placed on top of sponge Ti and Zr, and electrolytic Mn placed on top.

[0019] Further technical solution of the present invention: The specific process and parameters of step S2 are as follows: After covering the furnace lid, evacuate to 1×10 -2 ~5×10 -4 Pa, filled with pure argon gas at a pressure of 0.01 to 0.1 MPa as a protective gas.

[0020] A further technical solution of the present invention: In step S2, the smelting process adopts a power-incrementing melting and refining process. Specific parameters are as follows: Melting is carried out for approximately 10 minutes when the power supply is set to 10 kW, and then the power supply is increased by 5 kW every 5 minutes until it reaches 40 kW. After all raw materials have melted, the alloy temperature in the molten state is maintained at 1500–1550 °C for 3–5 minutes to ensure complete homogenization of the liquid alloy composition. The power-incrementing melting process creates a varying electromagnetic stirring force in the molten raw materials, enhancing the electromagnetic stirring effect, improving melting efficiency, promoting alloy composition homogenization, and most importantly, significantly improving the yield of the smelted alloy.

[0021] A further technical solution of the present invention: The rapid quenching operation in step S3 is as follows: the power supply is adjusted to approximately 30 kW to pre-cool the liquid alloy to approximately 1400 °C (1400 °C or close to 1400 °C). The liquid master alloy is directly injected into the tundish with NB nozzles embedded at the bottom. Under the action of its own gravity, the liquid alloy is continuously sprayed from the narrow nozzle slit (slit width 0.3±0.005 mm) at the bottom of the tundish onto the surface of the rotating water-cooled copper roller. The linear velocity of the rapid quenching roller surface is 3–20 m / s. The water inlet flow rate of the copper roller is ≥20 m³ / s. 3 The quenching rate is 18–20 °C. The rapidly quenched sheet detaches from the copper roller and falls into a collection frame below the copper roller. After the alloy cools to room temperature, a rapidly quenched alloy sheet with a thickness of 100–300 μm is obtained. The quenching rate (linear velocity of the copper roller surface) is a key technical parameter. Selecting an appropriate quenching rate can ensure that the rapidly quenched alloy strip has an almost complete nanocrystalline structure with a grain size between 30 and 100 nm.

[0022] A further technical solution of the present invention: In step S4, a 200-mesh sieve is used for sieving, and the diameter of the alloy powder passing through the sieve is ≤75 μm. The specific parameters of the ball milling process in step S4 are: ball-to-material ratio 20:1, rotation speed 250 rpm, and ball milling time 10-100 minutes.

[0023] Compared with existing technologies, the beneficial effects of this technical solution are as follows: In terms of composition design, this invention employs an A-side overstoichiometry, which significantly improves the hydrogen absorption capacity and activation performance of the alloy. Alloying with metallic elements Zr, Mn, Ni, and rare earth element Y significantly improves the alloy's activation performance. Regarding the preparation technology, rapid quenching combined with mechanical ball milling is used. This technology significantly reduces the grain size and improves the surface condition, disrupting the continuity of the dense oxide film on the alloy particles, allowing hydrogen to contact the fresh alloy surface and achieve activation. During rapid quenching and ball milling, numerous crystal defects are formed within and on the surface, increasing hydrogen nucleation sites and diffusion channels, further reducing the alloy's thermal stability and improving its hydrogen absorption and desorption kinetics. A particularly significant advantage is that the activated patented alloy is insensitive to air, maintaining its activation performance even after 72 hours of exposure to air, a characteristic that fully meets the needs of practical applications. This preparation process is simple and easy to operate, making it particularly suitable for the large-scale preparation of Ti-Fe based AB-type alloys.

[0024] In this invention, the above-described technical solutions can be combined with each other to achieve more preferred combinations. Other features and advantages of this invention will be set forth in the following description, and some advantages may become apparent from the description or be learned by practicing the invention. The objects and other advantages of this invention can be realized and obtained from what is particularly pointed out in the description and drawings. Attached Figure Description

[0025] The accompanying drawings are for illustrative purposes only and are not intended to limit the invention. Throughout the drawings, the same reference numerals denote the same parts.

[0026] Figure 1 These are binary phase diagrams for different alloy systems, mainly used to assist in alloy composition design. Figure 2 This is a medium-frequency induction melting-rapid quenching integrated furnace, mainly used for alloy melting and rapid quenching. Figure 2 a is an integrated furnace for medium-frequency vacuum induction melting and rapid quenching. Figure 2 b is a fast-quenching water-cooled copper roller; Figure 3 This is a schematic diagram of the fast-quenched alloy strip in Example 1; Figure 4The images show the SEM morphology of the cross-section (roller side) and longitudinal section (thickness direction) of the rapidly quenched thin strip in Example 1. Figure 4 a is the cross-section. Figure 4 b is the longitudinal section; Figure 5 The SEM morphology, EDS elemental distribution map, and energy dispersive spectroscopy (EDS) of the rapidly quenched thin strip in Example 1 are shown below. Figure 5 a represents the SEM morphology of the rapidly quenched thin strip. Figure 5 bh is the EDS element distribution map. Figure 5 ij represents the EDS energy spectrum of different regions; Figure 6 The SEM morphology of the ball-milled alloy powders in Examples 1-6 is shown. Figure 7 The XRD patterns of the ball-milled alloy powders in Examples 1-6 are shown below. Figure 8 The HRTEM morphology of the ball-milled alloy powders in Examples 1-6 is shown.

[0027] The design concept and mechanism of the present invention will be further described in detail with reference to the accompanying drawings and embodiments, so as to make the technical solution of the present invention clearer.

[0028] Ti-Fe based AB-type alloys possess higher hydrogen storage capacity (1.86 wt.%) than rare-earth-based AB5-type alloys, better room-temperature hydrogen storage performance and kinetics than magnesium-based alloys, and better cycle stability than Ti-Mn-based AB2-type alloys, while also being made from inexpensive raw materials. Therefore, they are considered highly promising solid-state hydrogen storage materials. Although significant progress has been made in the research of Ti-Fe based AB-type alloys, they are still a step away from practical application. This is mainly due to the stringent activation conditions, requiring high temperatures (450 °C) and high hydrogen pressures (6 MPa), as well as a long activation incubation period. Furthermore, activated alloys quickly lose their activity when exposed to air. These drawbacks are unacceptable for most applications. In addition, alloys prepared by traditional casting processes tend to have Ti enrichment and poor compositional uniformity. While subsequent high-temperature diffusion annealing can effectively eliminate segregation, the annealing process is lengthy, inefficient, energy-intensive, and significantly increases production costs.

[0029] Studies have confirmed that the activation difficulty of Ti-Fe based AB type alloys is mainly due to the presence of highly reactive Ti elements. During the preparation process, upon contact with air, a dense TiO2 oxide film forms on the alloy surface, severely hindering hydrogen contact and resulting in the alloy being unable to absorb hydrogen or absorbing very little hydrogen. Generally, prolonged exposure to high temperature and high pressure hydrogen atmosphere is required to break the oxide film on the alloy surface, restoring its activity and enabling successful hydrogen absorption. This process is known in the industry as activation. In summary, the alloy activation process includes the following steps: 1) breaking and removing the dense oxide film on the surface; 2) dissociating hydrogen molecules into hydrogen atoms on the alloy surface; 3) the dissociated hydrogen atoms diffuse through the surface into the alloy interior and interact with the metal to form hydrides. Among these steps, the dissolution and removal of the oxide film is usually the most difficult, generally requiring prolonged exposure to high pressure hydrogen and / or high temperature. Numerous studies have found that through scientific composition design and the application of advanced preparation techniques, the above activation steps can be accelerated, the required temperature and hydrogen pressure reduced, and the activation incubation period significantly shortened. The main mechanisms of this action include the following aspects: 1. Improving activation performance by adding or substituting elements to form new intermetallic compounds. Some argue that adding or substituting elements can form new intermetallic compounds, i.e., second phases, in alloys. These phases are mainly distributed at the phase boundaries of Ti-Fe alloys, where hydrogen absorption and desorption reactions are more likely to occur. The phase boundary between the second phase and the TiFe phase can serve as a channel for hydrogen diffusion within the alloy. Some second phases can absorb hydrogen before the TiFe phase, and this hydrogen absorption expansion promotes the pulverization of alloy particles, thereby improving activation performance.

[0030] 2. The formation of stable hydrides through element addition or substitution is beneficial for improving the activation performance of alloys. Studies have found that some elements, when added to Ti-Fe alloys, neither dissolve into the matrix alloy nor form new intermetallic compounds, but exist independently within the alloy. These elements have a strong affinity for hydrogen and preferentially form stable hydrides in the alloy; rare earth elements are typical examples of this characteristic. Adding trace amounts of rare earth elements to Ti-Fe based AB-type alloys can significantly reduce the activation conditions of the alloy.

[0031] 3. By adding or substituting elements, the composition of oxides on the alloy surface can be altered, increasing the reactivity of hydrogen with the alloy surface and improving the alloy's activation performance. Some argue that most metallic elements can form oxides under certain conditions. While these elements do not form a second phase when added to Ti-Fe based alloys, they can change the composition of oxides on the alloy surface, increasing the reactivity of the alloy surface with hydrogen. Typical examples include Mn, Cr, Co, and Cu.

[0032] 4. Refining grain size and altering surface conditions through alloying and preparation techniques can improve the activation performance of alloys. Some argue that by using non-stoichiometric methods, adding certain elements, or undergoing special preparation processes, highly active regions can be generated on the alloy surface, promoting hydrogen absorption reactions. Typical examples of alloying alloys include Pd, Ni, V, and Mn. In terms of preparation processes, ball milling and rapid quenching can significantly reduce the size of alloy particles and grains, improve their surface conditions, and enhance the alloy's activation performance.

[0033] The above are the main approaches and methods for improving the activation performance of Ti-Fe based AB type alloys. In particular, the effects of alloying and preparation processes on the activation performance of alloys are comprehensive, which provides a reliable basis and route for the design concept of this invention.

[0034] As is well known, when hydrogen atoms enter the crystal lattice of a metal or alloy, lattice expansion inevitably occurs because the diameter of the hydrogen atom is larger than the size of the interstitial space, resulting in lattice stress. This stress increases rapidly with the increase of hydrogen absorption capacity. When this stress exceeds the fracture strength of the metal or alloy, it leads to cracking of the matrix. In the solid-state hydrogen storage industry, this hydrogen-induced cracking phenomenon is called pulverization. The activation process of Ti-Fe based AB-type alloys is a hydrogen-induced cracking process. These alloys are generally brittle materials, so pulverization is more likely to occur. Introducing Zr and rare earth Y on the A side and Mn and Ni on the B side increases the microcracks and lattice stress in the Ti-Fe based alloy. Under the action of internal stress and a certain hydrogen pressure, these cracks can provide deeper diffusion channels. This provides conditions for the initial dissociation and adsorption of hydrogen on the relatively fresh surface at the crack tip, leading to lattice expansion of the material at the crack tip. Subsequently, new cracks are generated, and a large number of fresh surfaces appear explosively. Hydrogen diffuses more rapidly through the new channels, and then hydrogen dissociates and adsorbs on the crack surface deeper in the alloy. The room-temperature activation performance is affected by the reactivity between hydrogen molecules and the surface oxide layer, and the hydrogen reactivity is closely related to the composition of the oxide layer. Trace amounts of rare earth elements, such as La, Ce, Y, and Sc, and Zr partially substituting for Ti, as well as Mn, Cr, Cu, Co, Al, and Ni partially substituting for Fe, alter the composition of the alloy surface oxide layer, promote the reaction between hydrogen molecules and the oxide layer, accelerate the cracking, dissolution, and removal of the surface oxide layer, thereby improving the activation performance of the alloy at room temperature.

[0035] In terms of composition design, the addition of trace amounts of rare earth element Y can significantly improve the activation performance of the alloy, reduce the requirements for activation temperature and hydrogen pressure, and shorten the activation incubation period. Since rare earth element Y is itself a hydrogen-absorbing element, its addition is beneficial to increasing the hydrogen storage capacity of the alloy. Partial substitution of Ti with Zr can significantly increase the lattice constant and cell volume of the alloy, reduce the diffusion activation energy of hydrogen atoms, and promote alloy activation. Furthermore, Zr itself is a hydrogen-absorbing element, and its addition is beneficial to the alloy's capacity. Mn and Ni are considered excellent catalytic elements; their addition can improve the surface activity of the alloy, significantly reduce the dissociation energy of hydrogen molecules decomposing into hydrogen atoms on the alloy surface, accelerate the H2→2H reaction process, and promote alloy activation and hydrogen absorption kinetics.

[0036] In terms of alloy preparation technology, the vacuum induction rapid quenching + ball milling process can give the alloy a special microstructure, significantly improving its activation performance. Rapid quenching generates significant lattice stress within the alloy, accelerating fracture during activation and promoting the formation of numerous fresh surfaces, thus greatly enhancing its activation performance. The advantage of rapid quenching lies in obtaining ultrafine grains—nanocrystalline grains—and generating numerous grain boundaries, providing channels for hydrogen atom diffusion and significantly improving the alloy's hydrogen absorption and desorption kinetics. Due to the extremely rapid cooling rate of rapid quenching (>1000 ℃ / s), component segregation during solidification and cooling is suppressed, resulting in a more uniform elemental distribution within the alloy and improved hydrogen storage performance. Simultaneously, rapid quenching increases the cell volume of the alloy, lowering the hydrogen atom diffusion activation energy and accelerating the activation process. The grain refinement caused by rapid quenching leads to explosive growth of grain boundaries, which is beneficial for increasing the alloy's hydrogen storage capacity, as grain boundaries have a larger hydrogen storage capacity than grain interiors. It is important to note that the most crucial step in activation is the surface reaction of the alloy; therefore, the surface state of the alloy has a decisive influence on its activation performance. There are various methods to improve the surface condition of hydrogen storage alloys, but mechanical ball milling is the best choice considering factors such as process operability and economy. Ball milling directly breaks down the continuously distributed dense oxide film on the alloy surface, exposing a fresh surface that allows the alloy to directly absorb hydrogen. Simultaneously, mechanical ball milling significantly reduces the grain size and greatly increases the density of grain boundaries and defects, providing channels for the rapid diffusion of hydrogen atoms within the alloy. Therefore, it significantly improves the hydrogen absorption / desorption kinetics of the alloy while simultaneously improving activation performance. Detailed Implementation

[0037] Preferred embodiments of the present invention will now be described in detail with reference to the accompanying drawings, which form part of this application and are used together with the embodiments of the present invention to illustrate the principles of the present invention, but are not intended to limit the scope of the present invention.

[0038] Please see Figure 1-8This invention discloses a rare-earth Y-doped, easily activated Ti-Fe-based AB-type hydrogen storage alloy, specifically a Ti-Zr-Y-Fe-Mn-Ni AB-type hydrogen storage alloy, wherein the alloy's specific composition is: (Ti 1-x-y Y x Zr y ) 1.15 Fe 1.1-z- m Mn z Ni m , where x, y, z, and m are atomic ratios, and 0.01≤x≤0.04, 0.02≤y≤0.1, 0.1≤z≤0.3, and 0.05≤m≤0.12.

[0039] Preferably, x:y:z:m = 0.02:0.04:0.22:0.08.

[0040] The roles, contents, and atomic ratios of each component / element are determined based on the following: Overall, an overstoichiometric ratio on the A side is used to ensure the alloy's high hydrogen absorption capacity. While increasing the content of elements on the A side can improve the alloy's hydrogen absorption capacity, an excessively high A / B ratio will significantly reduce the alloy's activation performance. If the A / B value deviates significantly from 1:1, the alloy's hydrogen storage capacity will decrease markedly. It is worth noting that Ti-Fe based AB-type alloys are sensitive to changes in composition, and the synergistic effects between elements have a complex impact on the alloy's performance. Therefore, the content of each element in the alloy must be strictly controlled.

[0041] The content of element A is determined as follows: The elements on side A include Ti, Zr, and Y. Among them, Ti: In Ti-Fe based AB-type alloys, Ti is the main alloying element, and the Ti / Fe mass ratio has a decisive influence on the alloy's hydrogen storage performance. From the Ti-Fe binary phase diagram ( Figure 1 a) It can be observed that two stable intermetallic compounds exist in Ti-Fe alloys: TiFe phase and TiFe2 phase. The formation of TiFe and TiFe2 phases is affected by the proportion of titanium-iron alloy and the preparation process. The TiFe phase occupies a very small area and its formation conditions are quite demanding. Only when the Ti content is higher than 49.5 wt.% and lower than 52.5 wt.% can a body-centered cubic CsCl structure TiFe phase be formed through a eutectic reaction. When the Ti concentration is greater than 52.5 wt.%, α-Ti or β-Ti precipitates will appear; however, when the Ti concentration is less than 49.5 wt.%, a small amount of TiFe2 will form. TiFe alloys can react with H2 to form TiFeH. 0.10 ( α Phase), TiFeH 1.04 ( β (phase) and TiFeH1.95 ( γ Three alloy hydrides (phase), of which, α The phase is a hydrogen solid solution phase. β Harmony γ The phases are tetragonal and cubic, respectively. The phases that can stably exist in the alloy are... β Phase (TiFeH) 1.04 )and γ Phase (TiFeH) 1.95 The PCT curves of Ti-Fe alloys generally contain two hydrogen absorption / desorption plateaus, corresponding to the formation and decomposition of these two hydrides, respectively. TiFe2 does not possess hydrogen storage properties under normal conditions; therefore, the content of the TiFe2 phase in the alloy should be controlled during alloy design to ensure high capacity. Studies have found that appropriately increasing the Ti content in the alloy can reduce or even eliminate the TiFe2 phase, and the presence of a small amount of α-Ti in the alloy is beneficial to its hydrogen storage capacity and activation performance. Since Ti itself is a hydrogen-absorbing element, it readily reacts with hydrogen to form TiH2 hydrides. The formation of these hydrides leads to volume expansion of the alloy, creating microcracks and generating a fresh, oxide-free surface, thus initiating the hydrogenation process of the TiFe alloy. There are two viewpoints regarding the mechanism of the effect of excessive Ti on the alloy's activation performance. One viewpoint suggests that metallic Ti dispersed on the surface of the Ti-Fe alloy facilitates the dissociation of H2 molecules and the jumping of H atoms to the subsurface layer of the TiFe matrix. When hydrogen atoms on the Ti surface are transferred to TiFe, TiH2 hydrides are formed. x The first viewpoint suggests that the phase can serve as a rapid diffusion channel for hydrogen atoms. The second viewpoint posits that the transfer of hydrogen atoms during the hydrogenation / dehydrogenation process of Ti-Fe alloys occurs through the TiFe / Ti interface. Due to the higher storage energy at the interface, it provides the driving force for hydrogen atom transfer, meaning the phase interface plays a crucial role. It is important to note that the excess Ti must be strictly controlled; the presence of excessive α-Ti inevitably leads to a reduction in the main TiFe phase of the alloy, resulting in a significant decrease in hydrogen storage capacity.

[0042] Zr: From the Ti-Zr binary phase diagram ( Figure 1(b) It can be seen that Ti-Zr can be dissolved indefinitely, and the substitution of Ti and Zr in any proportion will not change the main phase structure. The role of Zr includes the following aspects: With increasing Zr content, the lattice constant and cell volume of Zr dissolved in the main TiFe lattice increase, and the activation energy for hydrogen atom diffusion decreases, which is beneficial to the diffusion of the alloy. This is because the atomic radius of Zr is larger than that of Ti, and the partial substitution of Ti by Zr increases the cell volume. Furthermore, since Zr itself is a hydrogen-absorbing element and has a greater affinity for hydrogen than Ti, it can form more stable hydrides during hydrogen absorption, thus improving the hydrogen absorption capacity of the alloy. It is worth noting that Zr hydrides are more stable, and excessive addition of Zr has an adverse effect on the hydrogen desorption performance of the alloy. Besides being dissolved in the matrix lattice, Zr may also form intermetallic compounds with Mn and Ni; their role in the alloy will be discussed later.

[0043] Recent studies have shown that adding trace amounts of rare earth elements (especially Y) to Ti-Fe based AB-type alloys can significantly improve their activation performance, requiring almost no incubation period. The mechanism by which rare earth elements affect the activation performance of Ti-Fe based alloys is still under investigation, but it is confirmed that rare earth elements have extremely low solid solubility in the TiFe phase and generally exist in a free state within the alloy. Y reacts with hydrogen to form stable hydrides, and the addition of Y undoubtedly benefits the alloy's activation ability and hydrogen storage capacity. This is because the atomic radius of Y (181 pm) is much larger than that of Ti (170 pm), and the addition of Y significantly increases the lattice constant and cell volume, thus significantly improving the alloy's hydrogen storage capacity. Simultaneously, the increased cell volume reduces the hydrogen diffusion activation energy, improving the alloy's activation performance. Furthermore, appropriate amounts of Y can significantly improve the alloy's activation performance and resistance to poisoning by impurity gases, and can effectively suppress elemental segregation, thereby significantly improving the hydrogen absorption / desorption hysteresis effect. Y is considered a very effective catalyst, and the hydrogen absorption process preferentially forms a highly stable YH2 / YH3 phase. Figure 1 c) It can serve as a nucleus for heterogeneous nucleation during the hydrogen absorption and desorption phase transition of the alloy, significantly improving the alloy's activation performance and hydrogen absorption and desorption kinetics. It should be noted that due to the atomic radius and specific gravity of Y, its solid solubility in the alloy is very low. Excessive addition of Y will reduce the alloy's mass hydrogen storage density and increase the raw material cost.

[0044] Determination of the content of element B: The elements on the B side include Fe, Mn, and Ni. Fe: Fe is the main alloying element in Ti-Fe based AB type alloys, primarily playing a catalytic role. The function of Fe in the alloy is to shrink the cell volume of the main TiFe phase, reducing the stability of the hydride phase, thereby satisfying the thermodynamic requirements for solid-state hydrogen storage. This allows the Ti-Fe based AB type alloy to reversibly absorb and release hydrogen at room temperature. However, the Fe content must be strictly controlled to ensure the alloy contains as much TiFe phase as possible. Figure 1 As can be seen, TiFe phase can only be formed through eutectic reaction when the iron content is higher than 51.3 wt.% and lower than 54.1 wt.%, and the temperature is within the range of 1358-1590 K. When the Fe content exceeds 68.2 wt.%, the TiFe phase continues to undergo eutectic reaction with Fe to form TiFe2 phase. Excessive Fe can also lead to the formation of α-Fe phase in the alloy, resulting in a significant reduction in the main TiFe phase and a decrease in the hydrogen storage capacity of the alloy. A noteworthy phenomenon is that the presence of an appropriate amount of α-Fe phase can improve the activation performance of Ti-Fe alloys. Some argue that adding an excessive amount of Fe to the alloy will create Fe atom-enriched regions on the alloy surface. Similar to the effect of Pa, Fe enriched on the alloy surface can also promote the decomposition of hydrogen molecules. In alloy design, the presence of TiFe2 phase and α-Fe phase is generally avoided as much as possible, as the appearance of these two phases inevitably reduces the alloy's capacity.

[0045] Mn and Ni: In solid-state hydrogen storage materials, Mn and Ni have similar roles, primarily exhibiting significant catalytic effects on the hydrogen absorption and desorption reactions of the alloy. Especially in rare-earth-based AB5 alloys, the role of Ni is almost irreplaceable. The improvement of the thermodynamics and kinetics of hydrogen absorption and desorption in magnesium-based alloys by Ni is also widely recognized in the academic community; the combined addition of Mn and Ni often achieves twice the result with half the effort. In Ti-Fe-based AB alloys, the main role of Mn and Ni is to improve the alloy's activation performance and enhance its hydrogen absorption and desorption kinetics. Mn and Ni have high solid solubility in the matrix; at low concentrations, they exist in the matrix in solid solution form. Because the atomic radii of Mn and Ni are very close to those of Fe, their influence on the solvent cell volume is limited. The contribution of adding Mn and Ni to the alloy's activation performance is mainly attributed to their improvement of the alloy's surface activity and changes in oxide composition. Besides being able to dissolve into the matrix phase lattice, Mn has a strong affinity for Zr and Ti, and in the alloy, some Mn can form the ZrMn2 phase with Zr and Ti. Figure 1 d) and TiMn2 phase ( Figure 1e) These phases, collectively known as the C14 Laves phase, have a significant catalytic effect on the activation performance of the alloy. Another important function of Mn and Ni partially substituting for Fe is to suppress the formation of the TiFe2 phase, increasing the content of the TiFe main phase in the alloy, which is beneficial to the alloy's hydrogen storage capacity. Another important function of Mn and Ni is to reduce the plateau pressure difference in the PCT dual-plateau, acting as a smoother for the dual-plateau, which is crucial for many applications. It should be noted that Ni is a precious metal, and excessive use of Ni will significantly increase the raw material cost of the alloy.

[0046] After comparing the hydrogen storage capacity, activation performance, and kinetic properties of a series of alloys, the optimal ratio of x, y, z, m was determined to be x:y:z:m = 0.02:0.04:0.22:0.08. Specifically, the hydrogen storage alloy has a nanocrystalline structure with a nanocrystalline content greater than 95%. By substituting rare earth elements and metallic elements for elements on the A and B sides, and combining this with a rapid quenching + ball milling process, the alloy can acquire activation performance that meets the requirements of practical applications.

[0047] Specifically, the Ti-Zr-Y-Fe-Mn-Ni type AB hydrogen storage alloy provided by this invention exhibits excellent activation and hydrogen absorption / desorption performance, with a hydrogen storage capacity ≥1.74 wt.%. It can be activated in one cycle at 30 °C and 3 MPa. After 200 hydrogen absorption / desorption cycles, the capacity retention rate is ≥98.9%, making it suitable for widespread application as a solid-state hydrogen storage material in various scenarios. It is particularly noteworthy that the activated patented alloy retains its activated state even after 72 hours of exposure to air.

[0048] This invention also provides a method for preparing a rare-earth Y-doped easily activated Ti-Fe-based AB-type hydrogen storage alloy, which is used to prepare the above-mentioned easily activated Ti-Fe-based AB-type hydrogen storage alloy, specifically including the following steps: S1: The ingredients are prepared according to the chemical formula composition calculation, with appropriate amounts of Mn and rare earth Y added for burn-off. S2: Place the prepared raw materials in an Al2O3 crucible in sequence, cover the furnace, evacuate to the preset vacuum level, and then fill with high-purity Ar gas at a certain pressure as a protective atmosphere. Use induction heating to heat the raw materials to a molten state and hold for a period of time to obtain a liquid master alloy with uniform composition. S3: Liquid master alloy (approximately 1400 ℃) is directly injected into an intermediate ladle with NB nozzles embedded at the bottom. The liquid alloy is continuously sprayed from the narrow slits of the NB nozzles at the bottom of the intermediate ladle onto the surface of a rotating water-cooled copper roller to obtain a fast-quenched alloy strip. S4: After mechanically crushing the fast-quenched alloy strip and passing it through a 200-mesh sieve, the alloy powder obtained from the sieve is loaded into a stainless steel ball mill jar together with stainless steel grinding balls. After vacuuming, high-purity Ar gas is introduced, and the mixture is ball-milled in an all-around planetary high-energy ball mill to obtain the ball-milled alloy powder, which is the hydrogen storage alloy.

[0049] In some embodiments, the burn-off amount of Mn and rare earth Y added in step S1 is 3% to 5% of the calculated addition amount. Since Mn and rare earth Y are volatile, an appropriate additional burn-off amount is needed to compensate for the burn-off during the smelting process. Based on multiple experiments, it has been found that a burn-off amount of 3% to 5% of the calculated addition amount is preferable.

[0050] In some embodiments, the order and specific operation of placing the raw materials in step S2 are as follows: The cleaned raw materials are prepared in proportion and placed sequentially in an Al2O3 crucible. Rare earth Y is placed at the bottom of the crucible, sponge Ti and Zr are placed above the rare earth Y, industrial pure Fe and electrolytic Ni are placed on top of the sponge Ti and Zr, and electrolytic Mn is placed on top. This placement order is determined based on the melting point and electromagnetic induction efficiency of the raw material metals. This arrangement of raw materials can effectively improve metal smelting efficiency, make the alloy composition more stable, the element distribution more uniform, and increase the yield of the finished product.

[0051] In some embodiments, the specific operation of melting the raw material in step S2 is as follows: evacuating to a vacuum of 5×10 -2 At pressures above Pa, high-purity argon gas at 0.01–0.1 MPa is introduced as a protective gas. The alloy is then melted using medium-frequency induction heating to ensure that all raw materials are completely melted. The temperature of the molten alloy is controlled at approximately 1500–1550 °C and held for 3–5 minutes to achieve complete homogenization of the alloy composition.

[0052] In some embodiments, the rapid quenching operation in step S3 is as follows: the power of the melting furnace is adjusted to 30 kW, and liquid alloy at a temperature of approximately 1400 ℃ is directly injected into the tundish with NB nozzles embedded at the bottom. Under the action of gravity, the liquid alloy is continuously sprayed from the narrow nozzle slits (slit width 0.3±0.005 mm) at the bottom of the tundish onto the surface of a rotating water-cooled copper roller with a linear velocity of 3 to 20 m / s. The water inlet to the copper roller is ≥20 m³ / s. 3 / h, water temperature 18~20 ℃. After the alloy cools to room temperature, a rapid quenching alloy strip is obtained, with a strip thickness of 100~300 μm and a grain size between 30~100 nm.

[0053] The following rapid quenching process parameters are particularly critical: Controlling the initial temperature of the rapid quenching of the liquid alloy is crucial. If the temperature is below 1400 ℃, the fluidity of the liquid alloy will be very poor, resulting in a slow flow rate through the nozzle and even solidification clogging the nozzle. Another problem with excessively low casting temperatures is poor adhesion between the liquid alloy and the rapid quenching roller, which significantly reduces the cooling rate of the alloy, making it impossible to obtain a nanocrystalline structure.

[0054] Choosing the appropriate nozzle slit width is a key technical parameter. If the nozzle slit is too large, it will reduce the cooling effect during the alloy solidification process and prevent the acquisition of an ultra-fine columnar crystal structure. If the nozzle slit is too small, the liquid alloy will solidify at the nozzle, resulting in discontinuous rapid quenching. In addition, if the nozzle is too small, it will increase the rapid quenching time, leading to a decrease in production efficiency.

[0055] Quenching rate is a key technical parameter and the most important factor determining the microstructure of an alloy. Appropriate quenching rates allow for rapid quenching, resulting in thin alloys with an almost completely nanocrystalline structure. Too high a quenching rate leads to the appearance of more amorphous phases in the alloy, resulting in a decrease in hydrogen storage capacity; too low a quenching rate causes significant grain coarsening, affecting the uniform distribution of alloy components and reducing activation performance and hydrogen absorption / desorption kinetics.

[0056] The flow rate and temperature of the cooling water for the copper rollers must also meet the requirements; otherwise, it will affect the microstructure of the rapidly quenched thin strip and thus the hydrogen storage performance of the alloy.

[0057] In some embodiments, step S4 specifically involves the following steps: After mechanically crushing the rapidly quenched alloy strip, it is passed through a 200-mesh sieve; the sieved alloy powder is then loaded into a grinding jar along with grinding balls; the grinding jar is evacuated and then filled with high-purity argon gas, and ball-milled in an omnidirectional planetary high-energy ball mill for 10–100 minutes, preferably 20 minutes. The ball-to-material ratio is 20:1, and the rotation speed is 250 rpm.

[0058] In step S4, using a suitable ball-to-particle ratio and appropriate rotation speed is to obtain the optimal ball milling efficiency. A reasonable ball milling process can obtain powder materials with finer particle size, better dispersion, and nanocrystalline structure within the same ball milling time. The surface state of the material is an important factor affecting its activation performance. The hydrogen absorption and desorption kinetics are controlled by the diffusion rate of hydrogen atoms within the material matrix. Grain boundaries provide good channels for hydrogen atom diffusion, and nanocrystalline structures have the highest grain boundary density, playing a unique role in hydrogen atom diffusion. Choosing the shortest possible ball milling time while disrupting the continuous distribution of oxides on the alloy surface and improving the alloy surface state is required for industrial-scale material preparation. Longer ball milling times result in higher energy consumption and lower material preparation efficiency. It should be noted that appropriately extending the ball milling time is beneficial to the material's performance. However, excessively long ball milling times can lead to the formation of a large number of amorphous phases in the alloy and can easily cause the aggregation or even welding of the milled material particles, significantly reducing hydrogen storage performance. Furthermore, with longer ball milling times, the contamination of the material by the milling balls and the container increases, leading to a decrease in the material's hydrogen storage performance.

[0059] The Ti-Zr-Y-Fe-Mn-Ni AB-type hydrogen storage alloy and its preparation method provided by this invention achieve the corresponding technical effects through scientific composition optimization and vacuum rapid quenching + ball milling. Appropriate rapid quenching + ball milling can obtain alloy powder with a special structure. A significant feature is that rapid quenching produces an ultrafine nanocrystalline structure within the alloy particles, while ball milling further refines the grains and disrupts the continuously distributed dense oxide film on the particle surface, greatly improving the surface condition of the particles and thus giving the alloy excellent activation performance. In particular, rapid quenching suppresses component segregation during solidification, and the rapid cooling significantly reduces the grain size of the alloy, providing a good channel for the rapid diffusion of hydrogen atoms within the alloy. Therefore, while improving activation performance, it significantly improves the hydrogen absorption and desorption kinetics of the alloy. Especially noteworthy is the elimination of the complex traditional casting + annealing process in the preparation process, which significantly reduces the preparation cost.

[0060] Based on the above component design and preparation method, the specific components of the embodiments and comparative examples of this invention are as follows: Example 1: (Ti 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 Example 2: (Ti) 0.92 Y 0.04 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22Ni 0.08 Example 3: (Ti) 0.88 Y 0.02 Zr 0.10 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 Example 4: (Ti 0.96 Y 0.02 Zr 0.02 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 Example 5: (Ti) 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.72 Mn 0.30 Ni 0.08 Example 6: (Ti 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.76 Mn 0.22 Ni 0.12 Example 7: (Ti 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 (3 m / s rapid quenching + ball milling) Example 8: (Ti) 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 (10 m / s rapid quenching + ball milling) Example 9: (Ti) 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 (20 m / s rapid quenching + ball milling) Example 10: (Ti 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80Mn 0.22 Ni 0.08 (Rapid quenching + 10-minute ball milling) Example 11: (Ti 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 (Rapid quenching + 50-minute ball milling) Example 12: (Ti) 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 (Rapid quenching + 100-minute ball milling) Comparative Example 1: (Ti 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 (Copper mold casting) Comparative Example 2: (Ti) 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 (Copper mold casting + annealing at 1000 ℃ for 8 hours) Comparative Example 3: (Ti 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 (quick quenching) Comparative Example 4: (Ti 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 (Copper mold casting + ball milling) Comparative Example 5: (Ti 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 (30 m / s rapid quenching + ball milling) Comparative Example 6: (Ti0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 (Rapid quenching + 5-minute ball milling) Comparative Example 7: (Ti 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 (Rapid quenching + 300-minute ball milling) According to the chemical formulas of each embodiment and comparative example, sponge Ti and Zr, rare earth Y, metallic Fe, electrolytic Mn and Ni were selected. The oxide layer on the surface of the raw materials was removed using specialized raw material cleaning equipment. Mn and rare earth Y were added to increase the burn-off rate by 3%–5% during batching. The technical parameters for each stage were as follows: the vacuum induction melting furnace was evacuated to 5 × 10⁻⁶ m³ / h before heating. -2 Pa; then, inert gas Ar gas at 0.01–0.1 MPa is introduced into the furnace as a protective gas; induction heating is used to raise the temperature of the liquid alloy to approximately 1500–1550 °C, and the temperature is held for 3–5 minutes to fully homogenize the alloy composition; the rapid quenching initiation temperature is approximately 1400 °C, and the quenching rate is 3–20 m / s. The rapidly quenched alloy strip is mechanically crushed and passed through a 200-mesh sieve, with a particle size of approximately 75 μm. The sieved alloy powder is then loaded into a stainless steel ball mill jar along with stainless steel grinding balls and ball-milled using an omnidirectional planetary ball mill for 10–100 minutes.

[0061] It should be emphasized that all process parameters can be appropriately selected within the above range to prepare the hydrogen storage alloy powder described in the patent. Therefore, although only one typical embodiment has been given in this invention, this embodiment is applicable to preparation methods with different parameters.

[0062] Process parameters for Example 1: S1: According to the chemical formula (Ti) 0.94 Y 0.02 Zr 0.04 ) 1.15 Fe 0.80 Mn 0.22 Ni 0.08 The raw materials selected were bulk sponge Ti and Zr, rare earth Y, industrially pure Fe, and electrolytic Mn and Ni. All raw materials had a metal purity ≥99.5%. They were weighed according to stoichiometric ratios, with each furnace containing 5 kg of materials: 2166 g of sponge Ti, 176 g of sponge Zr, 90 g of rare earth Y, 1870 g of industrially pure Fe, 531 g of electrolytic Mn, and 197 g of electrolytic Ni.

[0063] S2: Place the weighed bulk metal into the Al2O3 crucible of the medium-frequency induction furnace according to the designed process. Rare earth Y is placed at the bottom of the crucible, sponge Ti and Zr are placed above the rare earth Y, industrial pure Fe and electrolytic Ni are placed on top of the sponge Ti and Zr, and electrolytic Mn is placed on top. Then cover the furnace and evacuate for approximately 30 minutes until the vacuum level reaches 5 × 10⁻⁶. -2 After reaching a pressure of 0.04 MPa, high-purity Ar protective gas is introduced until the pressure reaches 0.04 MPa. The alloy is then melted using medium-frequency induction heating to ensure that all raw materials are completely melted. The temperature of the molten alloy is controlled at 1500–1550 °C and held for about 5 minutes to ensure that the alloy composition is completely homogenized.

[0064] S3: After adjusting the power supply to 30 kW, pre-cool the liquid alloy to about 1400 ℃, and directly inject it into the intermediate ladle with NB nozzles embedded at the bottom. The liquid alloy is continuously sprayed from the narrow gap of the NB nozzles at the bottom of the intermediate ladle (nozzle gap width 0.3±0.005 mm) onto the surface of a rotating (255 rpm) water-cooled copper roller with a line speed of 5 m / s. After the alloy cools to room temperature, a fast-quenched alloy sheet with an average thickness of 205 μm and an average grain size of about 55 nm is obtained.

[0065] S4: Mechanically crush the fast-quenched alloy sheet and pass it through a 200-mesh sieve. Weigh 100 grams of the sieved alloy powder and mix it with 2000 grams of stainless steel grinding balls. Put the mixture into a 1.0-liter stainless steel ball mill jar, evacuate it, fill it with high-purity argon gas, seal it, and then ball mill it in an all-around planetary high-energy ball mill for 20 minutes to obtain the hydrogen storage material described in the patent.

[0066] The raw materials for all examples and comparative examples were weighed according to their chemical formulas, and the other preparation process parameters were the same as those in Example 1 unless otherwise specified.

[0067] Examples 1 to 6 employ the exact same rapid quenching and ball milling processes, the difference being that the alloy composition varies. Through the study of the above examples, the optimal composition range of the alloy was determined.

[0068] The alloy composition and ball milling time of Examples 7-9 are exactly the same as those of Example 1, the difference being the quenching rate used, which is 3 m / s, 10 m / s, and 20 m / s, respectively. Through the study of the above examples, the quenching rate range of the rapid quenching process was determined.

[0069] The alloy composition and rapid quenching process of Examples 10-12 are exactly the same as those of Example 1, the difference being the ball milling time, which is 10 minutes, 50 minutes, and 100 minutes respectively. Through the study of the above examples, the shortest ball milling time corresponding to obtaining excellent hydrogen storage performance was determined.

[0070] The alloy composition of Comparative Example 1 is exactly the same as that of Example 1. The difference is that the Comparative Example did not undergo rapid quenching and ball milling. This Comparative Example confirms the role of rapid quenching and ball milling in improving the hydrogen storage performance of the material.

[0071] Comparative Example 2 involves subjecting Comparative Example 1 to high-temperature annealing. This comparative example reveals the effect of annealing on the activation performance and overall hydrogen storage performance of the alloy.

[0072] Comparative Example 3 has the same alloy composition and rapid quenching process as Example 1. The difference is that the comparative example did not undergo ball milling. This comparative example reveals the role of ball milling in improving the hydrogen storage performance of the material.

[0073] Comparative Example 4 has the same alloy composition and ball milling process as Example 1. The difference is that the comparative example did not undergo rapid quenching treatment. This comparative example reveals the role of rapid quenching in improving the hydrogen storage performance of the material.

[0074] Comparative Example 5 has the same alloy composition and ball milling time as Example 1. The difference is that the quenching rate of the rapid quenching is much greater than the upper limit of the quenching rate given in this patent. The purpose is to reveal the effect of selecting a suitable quenching rate on improving the hydrogen storage performance of the alloy.

[0075] Comparative Examples 6 and 7 have the same alloy composition and rapid quenching process as Example 1, except that the ball milling time exceeds the range given in this patent. The purpose is to reveal the effect of selecting an appropriate ball milling time on improving the hydrogen storage performance of the alloy.

[0076] The phase structure and phase composition of the rapidly quenched and ball-milled alloy were tested by XRD. The morphology and microstructure of the cross-sections of the rapidly quenched alloy strip were observed by scanning electron microscopy (SEM). The elemental distribution and energy spectrum of the alloy were tested by SEM and EDS.

[0077] Figure 1 The phase diagrams for binary alloys with different compositions are shown, namely Ti-Fe, Ti-Zr, YH, Zr-Mn, and Ti-Mn systems. These phase diagrams can be used to assist in alloy composition design and to identify potential new phases in the alloy.

[0078] Figure 2 Type A is an integrated vacuum melting and rapid quenching furnace, with a rated power of 40 kW and an ultimate vacuum of 5×10⁻⁶. -4 Pa. Figure 2 b is a water-cooled copper roller for a rapid quenching furnace, with a diameter of 375 mm and a maximum quenching speed of 30 m / s (rotation speed of 1500 rpm).

[0079] Figure 3 The fast-quenched alloy strip used in Example 1 has an average thickness of approximately 205 μm, which meets the design requirements.

[0080] Figure 4 The images show the SEM morphology of the horizontal (roller-facing) and longitudinal (thickness direction) sections of the rapidly quenched strip in Example 1. It can be observed that the horizontal section of the rapidly quenched strip exhibits a relatively uniform equiaxed crystal structure, while the longitudinal section shows a roughly parallel columnar crystal structure. This is due to the temperature gradient along the strip thickness direction (from the roller-facing to the back roller-facing) during rapid quenching.

[0081] Figure 5 The images show the SEM morphology, EDS elemental distribution map, and energy dispersive spectroscopy (EDS) spectrum of the rapidly quenched thin strip from Example 1. It can be seen that Ti, Zr, Y, Fe, Mn, and Ni elements are uniformly distributed, indicating that rapid quenching treatment can suppress component segregation during solidification. EDS analysis results show no significant compositional changes between different regions, which is impossible with traditional casting processes.

[0082] Figure 6 The SEM morphology of the ball-milled powders in Examples 1-6 shows that the particle dispersion of the alloy after ball milling is very good, and no obvious agglomeration is formed, which is obviously related to the short ball milling time.

[0083] Figure 7 The XRD patterns of the ball-milled alloys in Examples 1-6 show that ball milling significantly broadens the diffraction peaks of the alloys. This is due to the lattice stress and grain refinement generated after ball milling.

[0084] Figure 8 The HRTEM morphology of the ball-milled alloys in Examples 1-6 shows that the alloys have a nanocrystalline structure with an average grain size of about 55 nm.

[0085] The activation performance, hydrogen storage capacity, hydrogen absorption / desorption plateau pressure, hydrogen absorption / desorption kinetics, and cycle stability of the alloys in the examples and comparative examples were tested using a fully automated Sieverts system. The hydrogen absorption temperature was 30 °C, and the initial hydrogen pressure was 3 MPa; the hydrogen desorption temperature was 30 °C, and the cycle stability was tested at 1 × 10⁻⁶ MPa. -4 The test was conducted at a pressure of MPa. Using S 200 S represents the capacity retention of the alloy after 200 charge-discharge cycles. 200 =C 200 / C max ×100%. C max - Saturated reversible hydrogen absorption capacity; C 200 - Reversible hydrogen absorption capacity after 200 charge-discharge cycles. The activated alloy was exposed to air and allowed to stand for 72 hours before its activation state was tested again. The results are shown in Table 1.

[0086] Table 1. Solid-state hydrogen storage performance of different embodiments and comparative alloys The Ti-Zr-Y-Fe-Mn-Ni type AB hydrogen storage alloy provided by this invention exhibits excellent activation performance, hydrogen absorption and desorption kinetics, and cycle stability, with a hydrogen storage capacity ≥1.74 wt.%. Activation can be completed in a single hydrogen absorption and desorption cycle at 30 °C and an initial hydrogen pressure of 3 MPa. After 200 hydrogen absorption and desorption cycles, the alloy retains a capacity S0. 200 ≥98.9%.

[0087] Examples 1-6 show that, under identical conditions of rapid quenching and ball milling processes, changes in the alloy composition caused significant changes in its properties, indicating that the hydrogen storage performance of the material is closely related to its chemical composition.

[0088] Examples 7-9 used the same master alloy composition as Example 1, but the change in quenching rate during rapid quenching resulted in a significant change in the hydrogen storage performance of the alloy, indicating that selecting an appropriate quenching rate is extremely important for optimizing the hydrogen storage performance of the alloy.

[0089] Examples 10-12 used the same master alloy composition and rapid quenching process as Example 1, but with different ball milling times, resulting in significant changes in the hydrogen storage performance of the alloy. This indicates that selecting an appropriate ball milling time is extremely important for optimizing the hydrogen storage performance of the alloy.

[0090] Comparing Comparative Example 1 with Example 1, the alloy composition was exactly the same, but a copper mold casting process was used. The results showed that the overall hydrogen storage performance of the alloy, especially the activation performance, was significantly reduced. This indicates that only by combining scientific composition design with preparation process can good overall hydrogen storage performance, especially activation performance, be obtained.

[0091] Comparative Example 2 involved subjecting the alloy from Comparative Example 1 to prolonged high-temperature diffusion annealing. The results showed that annealing significantly improved the hydrogen storage performance, particularly the activation performance, of the as-cast alloy. This indicates that the compositional uniformity of the alloy has a significant impact on hydrogen storage performance. It also demonstrates that annealing can achieve good results even when rapid quenching is not feasible.

[0092] Compared with Example 1, Comparative Example 3 had the same alloy composition and rapid quenching process, but no ball milling was performed. The results showed that the activation performance of the alloy decreased significantly, demonstrating the key role of ball milling in the hydrogen storage performance of the alloy.

[0093] Compared with Example 1, Comparative Example 4 had the same alloy composition and ball milling process, but no rapid quenching treatment was performed. The results showed that the overall hydrogen storage performance of the alloy decreased significantly, demonstrating the key role of rapid quenching treatment in the hydrogen storage performance of the alloy.

[0094] Compared with Example 1, Comparative Example 5 had the same alloy composition and ball milling process, but the quenching rate of the rapid quenching was much greater than the quenching rate range given in this invention. The results showed that the overall hydrogen storage performance of the alloy, especially the activation performance, was significantly reduced, demonstrating the important role of quenching rate selection in improving the hydrogen storage performance of the alloy.

[0095] Comparative Example 6 and Example 1 have the same alloy composition and rapid quenching process, but the ball milling time is less than the minimum value given in this invention. The results show that the overall hydrogen storage performance of the alloy, especially the activation performance, is significantly reduced, indicating that the selection of ball milling time plays an important role in improving the hydrogen storage performance of the alloy.

[0096] Compared with Example 1, Comparative Example 7 had the same alloy composition and rapid quenching process, but the ball milling time was much longer than the maximum value given in this invention. The results showed that the overall hydrogen storage performance of the alloy, especially the activation performance, was significantly reduced. The main reason for this result was that the long ball milling time formed a certain number of amorphous phases in the alloy.

[0097] The above results demonstrate that the alloy of the present invention exhibits excellent activation performance, hydrogen absorption capacity, and cycle stability. Compared with similar alloys both domestically and internationally, the hydrogen storage performance of the alloy of the present invention is significantly improved. Furthermore, compared with similar alloys both domestically and internationally, the alloy of the present invention has lower raw material costs, giving it a clear advantage.

[0098] Obviously, compared with the traditional casting and annealing process, the alloy preparation process of the present invention is simple and easy to operate, and is fully suitable for large-scale production, with obvious advantages.

[0099] Of particular note is that all alloys in the embodiments maintained a fully activated state even after 72 hours of exposure to air. Furthermore, it was found that this property of the material is more closely related to its preparation process.

[0100] The above description is only a preferred embodiment of the present invention, but the scope of protection of the present invention is not limited thereto. Any changes or substitutions that can be easily conceived by those skilled in the art within the scope of the technology disclosed in the present invention should be included within the scope of protection of the present invention.

Claims

1. A rare-earth Y-doped, easily activated Ti-Fe-based AB-type hydrogen storage alloy, characterized in that, It contains the rare earth element Y and transition metals Mn, Ni and Zr, and its chemical formula is: (Ti 1-x-y Y x Zr y ) 1.15 Fe 1.1-z-m Mn z Ni m In the formula, x, y, z, and m are atomic ratios, and 0.01≤x≤0.04, 0.02≤y≤0.1, 0.1≤z≤0.3, and 0.05≤m≤0.

12.

2. The easily activated Ti-Fe-based AB-type hydrogen storage alloy according to claim 1, characterized in that, The preferred atomic ratio of the chemical formula is x:y:z:m = 0.02:0.04:0.22:0.

08.

3. The easily activated Ti-Fe-based AB-type hydrogen storage alloy according to claim 1, characterized in that, The easily activated Ti-Fe-based AB-type hydrogen storage alloy has a nanocrystalline main structure, with nanocrystalline content exceeding 95% and an average grain size in the range of 30–100 nm.

4. The easily activated Ti-Fe-based AB-type hydrogen storage alloy according to claim 1, characterized in that, The easily activated Ti-Fe-based AB-type hydrogen storage alloy has a multiphase structure containing a main TiFe phase and a second C14 Laves phase. The main TiFe phase and the second C14 Laves phase have cubic CsCl type and MgZn2 type structures, respectively.

5. A method for preparing a rare-earth Y-doped, easily activated Ti-Fe-based AB-type hydrogen storage alloy, characterized in that, The method for preparing the easily activated Ti-Fe-based AB-type hydrogen storage alloy according to any one of claims 1-4 specifically includes the following steps: S1: The ingredients are prepared according to the chemical formula composition calculation, with appropriate amounts of Mn and rare earth Y added for burn-off. S2: Place the prepared raw materials in an Al2O3 crucible, cover the furnace, evacuate the furnace, and then fill it with high-purity Ar gas as a protective gas. Heat the raw materials to a molten state and keep them at that temperature to obtain a liquid master alloy with uniform composition. S3: The liquid master alloy is directly injected into the intermediate jar with the NB nozzle embedded at the bottom. The liquid master alloy is continuously sprayed from the slit onto the surface of the rotating water-cooled copper roller to obtain an alloy strip with a thickness of 100-300μm. S4: The alloy strip is mechanically crushed and sieved to obtain alloy powder. The sieved alloy powder and stainless steel grinding balls are loaded into a stainless steel ball mill jar, vacuumed, and filled with high-purity argon gas. The jar is then ball-milled in an all-around planetary high-energy ball mill to obtain ball-milled alloy powder. The ball-milled alloy powder corresponds to an easily activated Ti-Fe-based AB-type hydrogen storage alloy.

6. The preparation method according to claim 5, characterized in that, In step S1, the amount of Mn and rare earth Y added is 3% to 5% of the calculated amount added.

7. The preparation method according to claim 5, characterized in that, In step S2, the order of placing the raw materials is as follows: rare earth Y is placed at the bottom of the crucible, sponge Ti and Zr are placed on top of rare earth Y, industrial pure Fe and electrolytic Ni are placed on top of sponge Ti and Zr, and electrolytic Mn is placed on top.

8. The preparation method according to claim 5, characterized in that, The temperature of the liquid master alloy at the start of the rapid quenching in step S3 is 1400 ℃. and / or In step S3, the slit width of the NB nozzle is 0.3 ± 0.005 mm. and / or The surface linear velocity of the rotating water-cooled copper roller in step S3 is 3 to 20 m / s.

9. The preparation method according to claim 5, characterized in that, The specific parameters for the ball milling process in step S4 are: ball-to-material ratio 20:1, ball mill speed 250 rpm, and ball milling time 10-100 minutes.

10. The preparation method according to claim 5, characterized in that, The specific process and parameters for step S2 are as follows: After closing the furnace lid, evacuate to 1×10⁻⁶. -2 ~5×10 -4 Pa, pure Ar gas at a pressure of 0.01 to 0.1 MPa is introduced as a protective gas.

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