An easily-activated high-capacity TiFe-based hydrogen storage alloy and a preparation method thereof

Abstract

The application discloses an easily-activated high-capacity TiFe-based hydrogen storage alloy and a preparation method thereof. The chemical composition of the alloy is Ti a Fe b Mn c N d RE e , and the microstructure of the alloy comprises a TiFe main phase and a second phase dispersedly distributed, the distribution density is not less than 1800 / mm2, and the fluctuation of the counting rate of the Mn element in the TiFe phase within a range of 30 microns in length is not more than ±15% of the average value. The alloy is subjected to first hydrogen absorption at 25 DEG C and 5 MPa hydrogen pressure, the hydrogen absorption amount of the alloy within 10000 seconds reaches more than 90% of the maximum hydrogen storage capacity, the maximum hydrogen storage capacity is greater than 1.85wt.%, the hydrogen release platform slope factor F is lower than 0.5, and the alloy can realize rapid hydrogen absorption without high-temperature activation. The alloy is an ideal solid-state hydrogen storage material with high capacity, low platform slope, excellent activation performance and low manufacturing cost, and is suitable for a solid-state hydrogen storage device for scale energy storage.

Description

Technical Field

[0001] This invention belongs to the field of hydrogen storage alloys, specifically relating to an easily activated, high-capacity TiFe-based hydrogen storage alloy and its preparation method. Background Technology

[0002] TiFe-based hydrogen storage alloys are widely recognized as one of the most promising solid-state hydrogen storage materials for large-scale applications due to their high theoretical hydrogen storage capacity (approximately 1.86 wt%), abundant raw material resources, and low cost. However, the commercial application of this alloy has long been limited by the following technical bottlenecks: First, initial activation is extremely difficult, typically requiring multiple cycles at temperatures above 400°C and high hydrogen pressures, placing stringent demands on hydrogen storage devices and systems; second, the hydrogen absorption and desorption kinetics are sluggish, especially at room temperature where the reaction rate is insufficient for practical applications; and third, they have weak resistance to gaseous impurity poisoning, and their cycle life decreases sharply when exposed to hydrogen containing trace amounts of O2, H2O, etc.

[0003] To overcome the aforementioned challenges, introducing alloying elements, particularly rare earth (RE) elements, has become an effective strategy for optimizing the performance of TiFe alloys. Studies have shown that the addition of rare earth elements can significantly improve the activation performance and reaction kinetics of the alloy through various mechanisms, such as purifying the alloy melt and forming a highly catalytically active second phase at grain boundaries to reduce the hydrogen diffusion barrier. However, this improvement is often accompanied by significant negative effects: on the one hand, the introduction of rare earth elements usually occupies lattice positions of the TiFe main phase in the alloy or forms a non-hydrogen-storage phase, leading to a significant loss in the effective hydrogen storage capacity of the alloy; on the other hand, the addition of different rare earth elements will alter the hydride formation enthalpy and plateau pressure of the alloy to varying degrees, specifically manifested as an increase in the slope of the equilibrium plateau of the pressure-composition-isotherm (PCT) curve and an exacerbation of hysteresis, resulting in a significant reduction in the effective hydrogen storage capacity within the actual operating pressure window.

[0004] Therefore, the core challenge facing current technology lies in how to fully utilize the beneficial effects of rare earth elements on activation and kinetic performance while minimizing their impact on the hydrogen storage capacity and PCT platform characteristics of the TiFe main phase through precise composition design and microstructure control. Current technology lacks a TiFe-based hydrogen storage material solution that can balance easy activation, high capacity, and a flat PCT platform. Developing a hydrogen storage alloy that maintains the low-cost advantage of TiFe alloys while achieving rapid room-temperature activation, high hydrogen storage capacity, and excellent PCT platform characteristics has become crucial for driving the technology towards large-scale application. Summary of the Invention

[0005] To address the aforementioned issues, the inventors, through repeated research, discovered a TiFe-based hydrogen storage alloy with high hydrogen storage capacity, easy activation, and good hydrogen absorption / desorption kinetics. Therefore, the objective of this invention is to provide an easily activated, high-capacity TiFe-based hydrogen storage alloy with the effects described above.

[0006] To achieve the above objectives, the present invention adopts the following technical solution.

[0007] One technical solution of the present invention:

[0008] A readily activated, high-capacity TiFe-based hydrogen storage alloy, characterized in that the chemical composition of the alloy is Ti a Fe b Mn c N d RE e The atomic percentages of each element satisfy the following conditions: 0.76≤a≤52.63, 39.27≤b≤44.27, 2.60≤c≤7.85, d≤1.05, 1.54≤e≤4.10, and in the main phase, a+b+c+d =100; RE is one or more rare earth elements selected from Y, Ce, La, Nd, Pr, and Sm, and RE must contain Y and / or Ce with a content greater than 80% of the total number of RE atoms; N is one or more elements selected from V, Zr, Al, Cu, B, Si, Mo, Sn, and Cr.

[0009] The microstructure of the alloy includes a TiFe main phase and a dispersed second phase, which is a RE-rich phase with a size between 0.5 and 10 micrometers and a distribution density of not less than 1800 particles / square millimeter.

[0010] Line scan analysis of the TiFe phase under a scanning electron microscope showed that the count rate of Mn element fluctuated within any 30-micrometer length range, not exceeding ±15% of its average value.

[0011] In the TiFe-based hydrogen storage alloy, the atomic ratio of Ti to (Fe+Mn) is between 1.03 and 1.11, and the alloy microstructure also contains a small amount of Ti-rich phase.

[0012] In the TiFe-based hydrogen storage alloy of the present invention, if the Ti / (Fe+Mn) ratio of the alloy is between 1.03 and 1.11, the hydrogen storage capacity of the alloy can be increased. However, if it is too high, it will lead to the formation of too much titanium-rich phase in the alloy, which will reduce the reversible hydrogen storage capacity.

[0013] The easily activated high-capacity TiFe-based hydrogen storage alloy was vacuumed at 70°C for 2 hours and then subjected to its first hydrogen absorption at 25°C and 5MPa hydrogen pressure. Within 10,000 seconds, the amount of hydrogen absorbed reached more than 90% of the maximum hydrogen storage capacity, and the maximum hydrogen storage capacity was higher than 1.85 wt.%.

[0014] The hydrogen desorption pressure-composition-isotherm (PCT) curve of the readily activated high-capacity TiFe-based hydrogen storage alloy at 25°C satisfies the following: hydrogen storage capacity w at a hydrogen pressure of 6 MPa. max The hydrogen content is 90% w max and 10% w max The slope factor F between two points is ≤ 0.5. Where F = log(P) 90% / P 10% ) / (90%w max -10%w max ), P 90% and P 10% The hydrogen content on the hydrogen release curve is equal to 90% w. max and 10% w max The equilibrium hydrogen pressure at that time is expressed in MPa.

[0015] Another technical solution of the present invention is as follows:

[0016] A method for preparing the readily activated, high-capacity TiFe-based hydrogen storage alloy as described above comprises the following steps: Following the design of the target composition Ti... a Fe b Mn c N d RE e Weigh the alloy raw materials, which are Ti, Fe, Mn, N and RE with a purity of not less than 99.5%; melt the raw materials under a protective atmosphere using electric arc melting or induction melting technology to ensure uniform composition, and then cool them to room temperature in the furnace to obtain the alloy ingot.

[0017] The protective atmosphere is argon, and the pressure range of the protective gas is 0.01~0.1MPa; during the melting process, the melt temperature is controlled between 1500℃ and 1700℃.

[0018] One application involves using the readily activated, high-capacity TiFe-based hydrogen storage alloy material described above in a large-scale solid-state hydrogen storage device for energy storage.

[0019] The beneficial effects of this invention are:

[0020] According to the preparation method of this invention, the core contradiction of achieving both easy activation and high capacity in the field of TiFe-based hydrogen storage alloys is successfully resolved through precise component design. Furthermore, during the preparation process, rapid activation at room temperature without high-temperature pretreatment is achieved while maintaining a high hydrogen storage capacity (>1.85 wt%) close to the theoretical value. In addition, this invention is the first to discover and confirm that introducing a specific range of Y and / or Ce (atomic percentage 1.54~4.10) into the TiFe-Mn system can significantly improve the uniformity of Mn element distribution in the TiFe main phase, even under conventional melting conditions. Moreover, the role of this specific content of Y and / or Ce is that, due to the intervention of rare earth elements in the solidification process, such as forming fine, dispersed metal or oxide particles in the melt as heterogeneous nucleation nuclei, significantly increasing the nucleation rate, refining grains, or altering solidification interface behavior, these effects fundamentally promote the micro-region homogenization of the overall composition, which is also key to obtaining a flat PCT platform and excellent kinetic performance.

[0021] Furthermore, according to the preparation method of the present invention, the designed alloy system is highly compatible with conventional electric arc melting or induction melting processes, without relying on complex rapid solidification equipment or harsh subsequent heat treatment regimes, which greatly reduces the complexity of the production process and equipment investment costs, and is conducive to its large-scale, low-cost industrial application. Attached Figure Description

[0022] Figure 1 The scanning electron microscope (SEM) image and linear scan of Mn element between two points in Example 1 are shown.

[0023] Figure 2 The scanning electron microscope (SEM) image and linear scan of Mn element fluctuations between two points are shown in Example 2.

[0024] Figure 3 To compare the scanning electron microscope (SEM) image and the fluctuation of Mn element between two points in Example 1.

[0025] Figure 4 To compare the scanning electron microscope (SEM) images and the fluctuations of Mn element between two points in Example 2.

[0026] Figure 5 To compare the scanning electron microscope (SEM) images and the fluctuations of Mn element between two points in Example 3.

[0027] Figure 6 The hydrogen absorption kinetic curves at 298 K for Examples 1 and 2 and Comparative Examples 1 and 2 are shown.

[0028] Figure 7 PCT curves at 298K for Examples 1, 2 and Comparative Examples 1, 2 and 3. Detailed Implementation

[0029] The present invention will be further described below with reference to the accompanying drawings and embodiments, but the scope of protection of the present invention is not limited to the following embodiments.

[0030] Example 1

[0031] First, according to the target component Ti 51.28 Fe 43.59 Mn 5.13 Y 2.56 Weighing was performed, and to compensate for the loss of Mn and rare earth elements during the smelting process, the raw materials were added in a trace excess of 102% of the theoretical value, with all raw materials having a purity of not less than 99.5%. Subsequently, the uniformly mixed raw materials were pressed into electrode blocks and placed in a water-cooled copper crucible in a non-consumable vacuum arc melting furnace. After evacuating to 5 × 10⁻³ Pa, high-purity argon gas was introduced to 0.04 MPa as a protective atmosphere. During smelting, Ti ingots were first melted by arc initiation to further absorb residual oxygen in the furnace, and then the formal alloy material was added for smelting. The melt temperature was 1700℃. To ensure uniform composition, each ingot was repeatedly smelted three times, and turned over after each smelting. Each smelting was held for 3 minutes to ensure that the alloy liquid was fully melted and homogenized. Finally, the obtained alloy ingots were cooled to room temperature in the furnace. The obtained alloy was measured using a scanning electron microscope equipped with an X-ray energy dispersive spectroscopy (EDS) instrument, and the average value of the Mn element intensity fluctuation range and the surface density of the rare earth-rich phase were obtained. The slope factor F, maximum hydrogen absorption capacity, and time required to reach 90% of the maximum hydrogen absorption capacity were obtained through hydrogen storage performance testing, and are shown in Table 1.

[0032] Examples 2-9 and Comparative Examples 1-4 below used different target compositions and were prepared according to the same method as in Example 1 to obtain corresponding alloy ingots. The determinations for each alloy ingot as described above were performed, and the results are shown in Table 1.

[0033] Example 2 According to the target component Ti 51.28 Fe 43.59 Mn 5.13 Ce 2.56 .

[0034] Comparative Example 1: Based on the target component Ti 51.28 Fe 43.59 Mn 5.13 La 2.56 .

[0035] Comparative Example 2: Based on the target component Ti 51.28 Fe 43.59 Mn 5.13 Sm 2.56 .

[0036] Comparative Example 3, based on the target component Ti 51.28 Fe 43.59 Mn 5.13 .

[0037] Example 3 According to the target component Ti 51.28 Fe 43.59 Mn 5.13 Y 1.54 .

[0038] Example 4 According to the target component Ti 51.28 Fe 43.59 Mn 5.13 Y 4.10 .

[0039] Example 5 According to the target component Ti 50.76 Fe 44.16 Mn 5.08 Y 2.03 La 0.51 .

[0040] Example 6 According to the target component Ti 52.63 Fe 42.11 Mn 5.26 Y 1.58 Ce 1.05 .

[0041] Example 7 According to the target component Ti 52.36 Fe 39.27 Mn 7.85 Zr 0.52 Ce 2.62 .

[0042] Example 8 According to the target component Ti 52.08 Fe 44.27 Mn 2.60 V 1.04 Y 2.08 .

[0043] Example 9 According to the target component Ti 52.36 Fe 41.88 Mn 5.24 Co 0.52 Y 3.14 .

[0044] Comparative Example 4: Based on the target component Ti 51.28 Fe 43.59 Mn 5.13 Y 0.51 .

[0045] Comparative Example 5: Based on the target component Ti 51.28 Fe 43.59 Mn 5.13 Ce 5.13 .

[0046] Table 1

[0047]

[0048] Determination of alloy ingots

[0049] The obtained alloy ingot samples were cut to obtain flat observation surfaces, and then the cross-sections were ground and polished using standard metallographic sample preparation procedures. To obtain clear backscattered electron contrast images, a diamond polishing paste with a particle size of 0.1 μm was used for final fine polishing. Subsequently, the prepared samples were observed for microstructure and compositional analysis using a scanning electron microscope equipped with an X-ray energy dispersive spectrometer. First, representative micro-regions were selected in backscattered electron mode, and BSE images were acquired. To quantitatively analyze the elemental concentration gradient at specific locations, a 30 μm long scan path was set at the TiFe main phase location to perform a line scan. The continuous variation curves of X-ray intensity along the scan path for each characteristic element (Ti, Fe, Mn, etc.) were recorded to assess the magnitude of elemental content fluctuations.

[0050] Figures 1 to 5 Scanning electron microscopy (SEM) backscattered electron (BSE) spectra of Examples 1 and 2, and Comparative Examples 1, 2, and 3, along with their corresponding line scan Mn elemental distribution curves, are presented. In each sample's BSE spectrum field of view, three paths were randomly selected for line scan analysis to ensure the representativeness of the results.

[0051] right Figures 1-5 Observation of the BSE spectra showed that, in addition to the matrix phase, all the alloys in the figures contained a small amount of titanium-rich phase (black areas) in their microstructure. However, Figure 1 Example 1 and Figure 2 In Example 2 (with added Y or Ce), in addition to the titanium-rich phase, a clear and diffusely distributed second phase rich in rare earth elements (white area) was also observed. Energy dispersive spectroscopy (EDS) analysis confirmed that it mainly consists of yttrium and cerium oxides, with a size distribution between 0.5 and 10 micrometers and a high distribution density of 1800–3500 particles / mm². In contrast, Figure 3 Comparison Example 1 and Figure 4 In Comparative Example 2 (with added La and Sm), the rare earth-rich phases (mainly oxides) partially overlapped with those of Examples 1 and 2 in size range (0.1 to 10 micrometers), but their distribution density was significantly lower, at 1000–1700 particles / mm², especially sparser in the La-added alloy. Figure 5The microstructure of Comparative Example 3 (without rare earth elements) consists only of a matrix phase and a small amount of titanium-rich phase. This suggests that, under the influence of this specific content of Y and / or Ce, the rare earth elements form fine, dispersed metal or oxide particles in the melt, refining the grains or altering the solidification interface behavior, fundamentally promoting the micro-regional homogenization of the overall composition. Furthermore, the differences in the microstructures of the above figures directly determine the uniformity of Mn element distribution: from Figure 5 The line scan results show that, due to the lack of rare earth elements, the Mn element intensity in Comparative Example 3 fluctuates within the range of ±20% to 24% of the average value, indicating generally poor uniformity. Meanwhile... Figure 1 and Figure 2 In the embodiments of the present invention, the effects of adding specific rare earth elements show significant differentiation: namely, the Mn element intensity curves of Examples 1 and 2 fluctuate smoothly and stably throughout the entire scanning path, fluctuating only within the range of ±12% to 15% of the average value, indicating a highly uniform compositional distribution. In stark contrast, Figure 3 and Figure 4 The Mn element curves of Comparative Examples 1 and 2 both exhibit large and violent fluctuations, with multiple sharp peaks and valleys. Their fluctuation ranges are between ±30%~35% and ±33%~50% of the average value, respectively. This proves that the Mn element has undergone severe microscopic segregation in both examples, and the compositional uniformity is significantly worse than that of Examples 1 and 2 of the present invention.

[0052] The results above indicate that introducing Y and / or Ce within a specific content range (atomic percentage 1.53–4.10) into the TiFe-Mn system can significantly improve the uniformity of Mn distribution in the TiFe main phase, even under conventional smelting conditions. This can be explained by the fact that rare earth elements Y and / or Ce interfere with the solidification process, acting as heterogeneous nucleation sites, significantly increasing the nucleation rate, refining grains, or altering the solidification interface behavior, thereby promoting micro-regional homogenization of the overall composition.

[0053] Furthermore, in the alloy material of this invention, the hydrogen storage performance of the hydrogen storage alloy can be represented by a hydrogen absorption kinetic curve and a pressure-composition-temperature curve (PCT curve). The activation performance of the material can be obtained from the hydrogen absorption kinetic curve, and the hydrogen release plateau pressure and slope of the hydrogen release plateau can be obtained from the PCT curve. The specific testing steps are as follows: Take approximately 2g of alloy powder smaller than 100 mesh, vacuum at 70°C for 2 hours, and then cool to 25°C. Introduce 5MPa hydrogen gas for hydrogen absorption activation, and then vacuum at 70°C for 1 hour. Repeat this cycle 3-4 times to obtain the hydrogen absorption kinetic curve. After cooling to 25°C, perform the PCT curve test (using the Sieverts method). (See attached diagram) Figure 6 and Figure 7 middle.

[0054] like Figure 6 The figure shows the initial hydrogen absorption kinetics curves of Examples 1, 2, and Comparative Examples 1-3 under activation treatment at 70°C with vacuum for 2 hours. It can be seen that Comparative Example 3, without rare earth elements, exhibits poor hydrogen absorption kinetics, with slow growth in hydrogen absorption capacity within 50,000 seconds, and its initial hydrogen absorption amount is only 1.64 wt.%. In contrast, Examples 1 and 2 (with Y or Ce added, respectively) show significantly enhanced hydrogen absorption kinetics, both reaching 90% of their maximum hydrogen absorption capacity within 9,600 seconds, and both achieving an initial maximum hydrogen absorption capacity of over 1.9 wt.%. Comparative Examples 1 and 2 (with La or Sm added, respectively) show significantly lower hydrogen absorption rates, requiring 35,700–36,500 seconds to reach 90% of their maximum hydrogen absorption capacity, with final hydrogen absorption capacities of only 1.84 wt.% and 1.76 wt.%, respectively. These results indicate that the presence of an appropriate amount of rare earth-rich phase is beneficial to improving the activation performance of the alloy, and the higher the content, the more significant the activation-promoting effect.

[0055] Figure 7 The PCT curves at 298K for Examples 1 and 2 and Comparative Examples 1, 2, and 3 are shown below. Figure 7 From the various figures, based on the PCT curves, the hydrogen storage capacity w at a hydrogen pressure of 6 MPa can be obtained. max The hydrogen content on the hydrogen release curve is 90% w. max and 10%w max The slope factor F between two points (F = log(P)) 90% / P 10% ) / (90%w max -10%w max ), where P 90% and P 10% The hydrogen content on the hydrogen release curve is equal to 90% w. max and 10% w max The hydrogen pressure at that time is in MPa. The relevant results are summarized in Table 1.

[0056] As shown in Table 1, the slope factors of Examples 1 and 2 are 0.40 and 0.49, respectively, significantly lower than those of Comparative Example 3 (F=0.61), which exhibits more severe fluctuations in Mn element distribution. The slope factors of Comparative Examples 1 and 2 further increase, reaching 0.62 and 0.65, respectively. This trend clearly indicates that the greater the fluctuation in Mn element distribution (see the average value column for Mn element intensity fluctuation range in Table 1), the higher the slope factor, and the more pronounced the plateau tilt. These results also suggest that while the addition of La and Sm affects the maximum hydrogen absorption capacity in hydrogen absorption kinetics, it not only fails to improve the alloy plateau characteristics but may also have a negative impact. The increase in the slope factor directly leads to a decrease in effective hydrogen storage capacity. Compared to Example 1 with a flat plateau in the figure, Comparative Examples 1 and 2 show a significant reduction in reversible hydrogen storage capacity under the same application conditions. A flat platform ensures that the system pressure remains relatively stable during the main stages of hydrogen absorption or release, which is considered to reduce the requirements for external pressure management systems and improve the controllability and safety of system operation. On the other hand, a tilted platform leads to a wider operating pressure range for the same capacity utilization, increases compression energy consumption, and frequent pressure fluctuations can also exacerbate material pulverization and container fatigue, thus being detrimental to the long-term cycle life of the device.

[0057] Furthermore, as shown in Table 1, in Examples 3 and 4, while maintaining other conditions consistent with Example 1, the amount of Y added was adjusted to an atomic percentage of 1.54 and 4.10, respectively. Scanning electron microscopy and hydrogen storage performance tests showed that both exhibited good activation performance and hydrogen release plateau characteristics, with the Mn element distribution fluctuation range remaining at a low level of ±13.9% and ±13.8% of the average value, respectively. The initial hydrogen absorption kinetics were comparable to those of Example 1, both reaching more than 90% of the maximum hydrogen absorption capacity within 10,000 seconds, with maximum hydrogen absorption capacities of 1.93 wt.% and 1.90 wt.%, respectively; the slope factors F were 0.42 and 0.41, respectively. In contrast, Comparative Example 4 further reduced the Y content to 0.51. The results showed that its Mn element fluctuation range expanded to ±20.2% of the average value, the hydrogen absorption kinetics deteriorated significantly, requiring more than 23,400 seconds to reach 90% of the initial maximum hydrogen absorption capacity, and the slope factor worsened to 0.60. This proves that when the rare earth content is below the lower limit required by the present invention, it is impossible to effectively induce the formation of a sufficient number and size of rare earth-rich phases (second phases), and thus the expected effect of improving activation performance and platform characteristics cannot be achieved. It can be seen that when the Y content is within the range of the present invention, it can effectively improve the alloy microstructure and hydrogen storage performance.

[0058] On the other hand, Comparative Example 5 increased the Ce content to 5.13% based on Example 2. Experiments showed that although its activation performance was acceptable (reaching 90% capacity within 12300 seconds), the maximum hydrogen storage capacity significantly decreased to 1.53%, and the slope factor increased to 0.59. This is because the excessive rare earth elements, on the one hand, reduced the effective hydrogen storage components in the matrix, and on the other hand, formed coarse rare earth oxide phases. This saturated the positive effect of suppressing Mn segregation and even introduced new compositional inhomogeneities, leading to a deterioration in capacity and plateau characteristics. This demonstrates that when the rare earth content is within the range of this invention, the microstructure and hydrogen storage performance of the alloy material obtained by this invention can be effectively improved.

[0059] Furthermore, when the present invention contains two or more rare earth elements, such as in Example 5 (addition of 0.51La + 2.31Y) and Example 6 (addition of 1.05Ce + 1.58Y), it exhibits excellent performance similar to that of adding Y or Ce alone: ​​the Mn element content fluctuates gently (±12.5%~±13.6%), the hydrogen absorption rate is fast (reaching 90% capacity within 9000 seconds), the maximum hydrogen absorption capacity is approximately 1.87~1.93 wt.%, and the slope factor is stable at 0.38~0.40. This indicates that, provided the total rare earth content is controlled within the scope of the present invention, using multiple rare earth elements can also achieve the objectives of the present invention, and in some cases, the synergistic effect between elements may produce even better results.

[0060] In summary, Examples 1 to 6 collectively demonstrate that by controlling the content of the aforementioned rare earth elements within the scope of this invention, an advanced hydrogen storage alloy with uniform Mn element distribution (fluctuation range < ±15%), excellent initial activation performance, high hydrogen storage capacity, and a flat PCT plateau (slope factor F < 0.50) can be successfully obtained. All comparative examples, deviating from the scope of this invention, exhibited significant deterioration in at least one key performance indicator.

[0061] Examples 7–9 further introduced small amounts of Zr, V, and Co elements into the TiFe-Mn-RE system. The Mn content fluctuated between ±13.5% and ±14.1%, the rare earth-rich phase distribution density was 2300–3100 particles / mm², the slope factor F was below 0.43, and the maximum hydrogen absorption capacity was above 1.85 wt.%, indicating that the introduction of appropriate amounts of N element did not significantly reduce the hydrogen storage performance. In other words, the introduction of N element played a "fine-tuning" role in the system of this invention, further optimizing the compositional uniformity, plateau characteristics, and overall hydrogen storage performance without significantly altering the main phase structure, demonstrating the flexibility and scalability of the component design of this invention.

[0062] This invention relates to an easily activated, high-capacity TiFe-based hydrogen storage material and its preparation method. The material successfully resolves the technical contradiction between the difficulty in activation, slow kinetics, and the inability to simultaneously achieve high capacity in traditional TiFe-based hydrogen storage alloys. Furthermore, its preparation process ensures excellent overall performance and low production costs, providing crucial material support for the large-scale application of solid-state hydrogen storage technology.

[0063] The above description is merely a preferred embodiment of the present invention, but the scope of protection of the present invention is not limited thereto. Any variations 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. Therefore, the scope of protection of the present invention should be determined by the scope of the claims.

Claims

1. A readily activated, high-capacity TiFe-based hydrogen storage alloy, characterized in that, The chemical composition of the alloy is Ti a Fe b Mn c N d RE e The atomic percentages of each element satisfy the following conditions: 50.76≤a≤52.63, 39.27≤b≤44.27, 2.60≤c≤7.85, d≤1.05, 1.54≤e≤4.10, and in the main phase, a+b+c+d =100; RE is one or more rare earth elements selected from Y, Ce, La, Nd, Pr, and Sm, and RE must contain Y and / or Ce with a content greater than 80% of the total number of RE atoms; N is one or more elements selected from V, Zr, Al, Cu, B, Si, Mo, Sn, and Cr. The microstructure of the alloy includes a TiFe main phase and a dispersed second phase, which is a RE-rich phase with a size between 0.5 and 10 micrometers and a distribution density of not less than 1800 particles / square millimeter.

2. The TiFe-based hydrogen storage alloy according to claim 1, characterized in that, Line scan analysis of the TiFe phase under a scanning electron microscope showed that the count rate of Mn element fluctuated within any 30-micrometer length range, not exceeding ±15% of its average value.

3. The TiFe-based hydrogen storage alloy according to claim 1, characterized in that, In the TiFe-based hydrogen storage alloy, the atomic ratio of Ti to (Fe+Mn) is between 1.03 and 1.11, and the alloy microstructure also contains a small amount of Ti-rich phase.

4. The TiFe-based hydrogen storage alloy according to claim 1, characterized in that, After being evacuated at 70°C for 2 hours, the TiFe-based hydrogen storage alloy was subjected to its first hydrogen absorption at 25°C and 5MPa hydrogen pressure. Within 10,000 seconds, the amount of hydrogen absorbed reached more than 90% of the maximum hydrogen storage capacity, and the maximum hydrogen storage capacity was higher than 1.85 wt.%.

5. The TiFe-based hydrogen storage alloy according to claim 1, characterized in that, The hydrogen desorption pressure-composition-isotherm curve of the TiFe-based hydrogen storage alloy at 25℃ satisfies the following: hydrogen storage capacity w at a hydrogen pressure of 6MPa. max The hydrogen content is 90% w max and 10% w max The slope factor F between two points is ≤0.5, where F=log(P 90% / P 10% ) / (90%w max -10%w max ), P 90% and P 10% The hydrogen content on the hydrogen release curve is equal to 90% w. max and 10% w max The equilibrium hydrogen pressure at that time is expressed in MPa.

6. A method for preparing an easily activated, high-capacity TiFe-based hydrogen storage alloy according to any one of claims 1 to 5, characterized in that, The steps are as follows: Following the target component Ti designed above... a Fe b Mn c N d RE e Weigh the alloy raw materials, which are Ti, Fe, Mn, N and RE with a purity of not less than 99.5%; melt the raw materials under a protective atmosphere using electric arc melting or induction melting technology to ensure uniform composition, and then cool them to room temperature in the furnace to obtain the alloy ingot.

7. The preparation method according to claim 6, wherein the protective atmosphere is argon, and the pressure range of the protective gas is 0.01~0.1MPa; during the melting process, the melt temperature is controlled between 1500℃ and 1700℃.

8. An application in which the TiFe-based hydrogen storage alloy of any one of claims 1 to 5 or the TiFe-based hydrogen storage alloy obtained by the preparation method of claim 6 or 7 is used in a solid-state hydrogen storage device for large-scale energy storage.

Images