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Article

Effect of Mn Element on the Structures and Properties of A2B7-Type La–Y–Ni-Based Hydrogen Storage Alloys

by
Anqiang Deng
1,2,3,
Yongchun Luo
1,2,*,
Jianfei Zhou
1,
Yunding Xie
1,
Yuan Yuan
1,
Xiaoyan Kang
1,
Bingjin Shen
1 and
Haimin Zhang
1
1
School of Materials Science and Engineering, Lanzhou University of Technology, Lanzhou 730050, China
2
State Key Laboratory of Advanced Processing and Recycling of Nonferrous Metals, Lanzhou University of Technology, Lanzhou 730050, China
3
School of Mechanical Engineering, Ningxia University, Yinchuan 750021,China
*
Author to whom correspondence should be addressed.
Metals 2022, 12(7), 1122; https://doi.org/10.3390/met12071122
Submission received: 1 June 2022 / Revised: 27 June 2022 / Accepted: 27 June 2022 / Published: 30 June 2022
(This article belongs to the Topic Metal Hydrides: Fundamentals and Applications)
Browse Figures
Versions Notes

Abstract

:
The structures, hydrogen storage behaviors and electrochemical properties of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) alloys were analyzed by X-ray diffraction, Neutron powder diffraction, pressure–composition isotherms and electrochemical tests. All alloys have a multiphase structure. With the increase in Mn content, the Gd2Co7-type phase of the alloys gradually transforms into the Ce2Ni7-type phase; the Mn atom mainly occupies the Ni sites in the [AB5] subunit and the interface between the [AB5] and [A2B4] subunits; the V[A2B4]/V[AB5] continuously decreases from 1.045 (x = 0) to 1.019 (x = 0.3), which reduces the volume mismatch between [A2B4] and [AB5] subunits. The maximum hydrogen absorption of the series alloys increases first and then decreases, and the addition of Mn effectively promotes the hydrogen absorption/desorption performance of the alloys. The maximum discharge capacity of the alloy electrodes is closely related to their hydrogen storage capacity at 0.1 MPa and hydrogen absorption/desorption plateau pressure. The cyclic stability of all the Mn-containing alloy electrodes is improved clearly compared to that of Mn-free alloy electrodes, because the volume mismatch between the [AB5] and [A2B4] subunits of the alloys becomes smaller after the addition of Mn, which can improve the structural stability and reduce the corrosion of alloys during hydrogen absorption/desorption cycles. When the Mn content is between 0.1 and 0.15, the Ce2Ni7-type phase of the alloys has high abundance and the alloy electrodes exhibit excellent overall performance.

1. Introduction

The proposal of the goal of “carbon peak and carbon neutrality” will promote the rapid development of green energy vehicles. The large-scale application of new energy vehicles, such as pure electric vehicles (EVs), hybrid electric vehicles (HEVs) and fuel cell vehicles (FCVs), has put forward new requirements for power batteries [1,2,3,4]. Among the large number of rechargeable batteries, the nickel-metal hydride (Ni-MH) battery plays an important role in the battery-powered electric vehicle market, especially in hybrid electric vehicles, because of its excellent high-rate and low-temperature discharge capabilities, resistance to overcharge/discharge ability and significant safety [5,6,7].
Hydrogen storage alloys are used as Ni-MH anode materials, and their structure and properties affect the performance of batteries. Compared with traditional AB5-type (A is a rare earth; B is a transition metal) hydrogen storage alloys, La–Mg–Ni-based superlattice hydrogen storage alloys have the advantages of easy activation, high discharge capacity and high kinetic performance, so they have received considerable attention as negative electrode materials for advanced nickel-metal hydride (Ni-MH) batteries [5,8,9,10,11,12,13]. However, the key element, Mg, in La–Mg–Ni system alloys is volatile at high temperatures. Therefore, it is not only difficult to control the composition but also has potential safety hazards during the preparation of La–Mg–Ni alloys [14,15].
The superlattice structure alloys are formed by stacking the [A2B4] subunit (Laves-type structure) and [AB5] (CaCu5-type structure) subunit along the c-axis [16]. La2Ni7 alloy is prone to hydrogen-induced amorphization, hydrogen-induced phase transformation and disproportionation during the hydrogen absorption/desorption process [17,18,19]. This is mainly due to the large volume mismatch between the [La2Ni4] and [LaNi5] subunits, and hydrogen only enters into the [La2Ni4] subunit and not into the [LaNi5] subunit, which causes the “anisotropic expansion” of the alloy during hydrogenation. After adding Mg to La2Ni7 alloy, since the atomic radius of Mg (RMg = 1.72 Å) is smaller than that of La (RLa = 2.74 Å), and Mg only occupies the La sites in the [A2B4] subunit, with the addition of Mg, the volume mismatch between the [A2B4] and [AB5] subunits reduces, which effectively improves the hydrogen absorption/desorption properties of the alloys [20,21]. Unlike La2Ni7, Latroche et al. [22] found that the hydrogen storage capacity of Y2Ni7 alloy did not decay after several hydrogen absorption/desorption cycles at 0.1 and 0.7 MPa, and the alloy structure remained unchanged, but there were three hydrogen absorption/desorption plateaus at 10 MPa and the plateau pressures were higher (0.055 MPa, 0.5 MPa, 2.9 MPa); therefore, Y2Ni7 alloy was not properly used as an alloy electrode material. Paul-Boncour et al. [23] studied the La2−xYxNi7 (x = 0–2) alloy and found that Y preferentially occupies La sites in the [A2B4] subunit and the atomic radius of Y (RY = 2.27 Å) is smaller than that of La (RLa = 2.74 Å); the addition of Y in La–Y–Ni alloys helps to adjust the volume mismatch between the [A2B4] and [AB5] subunits. Therefore, the effect of Y in La–Y–Ni alloys is similar to the Mg in La–Mg–Ni alloys. In addition, Yuan et al. [14] studied the La3−xYxNi9.7Mn0.5Al0.3 (x = 1, 1.5, 1.75, 2, 2.25, 2.5) alloys and found that the alloy electrodes had better overall electrochemical performance when x = 1.75–2.25.
In traditional AB5-type hydrogen storage alloys, Mn is an essential and key element, which occupies the 3g and 2c positions in the CaCu5 structure, and is more effective in adjusting the hydrogen absorption/desorption plateau pressure than Al for AB5-type alloys [24]. In La–Mg–Ni alloys, the substitution of Mn for Ni can not only decrease the plateau equilibrium pressure and increase the discharge capacity [25], but also increase the catalytic activity of the alloy electrode [26,27]. At the same time, the A2B7-type superlattice hydrogen storage alloys have attracted much attention as new-generation nickel-metal hydride battery anode materials [28,29,30].
In this study, we selected Y0.75La0.25Ni3.5 alloy as the original alloy, and the effect of the Mn element on the structures and properties of A2B7-type La–Y–Ni-based hydrogen storage alloys was investigated systematically. This study will offer a guide for further developing the La–Y–Ni system hydrogen storage alloys.

2. Materials and Methods

2.1. Sample Preparation

Y0.75La0.25Ni3.5−xMnx(x = 0, 0.05, 0.1, 0.15, 0.2, 0.3) hydrogen storage alloys were prepared by arc melting under a 0.05 MPa argon (Ar) atmosphere, and as-cast alloys were melted three times to obtain a homogeneous ingot. Subsequent annealing was performed for 24 h at 1173 K in a 0.2 MPa argon (Ar) atmosphere. The purity of all component elements was above 99 wt.%. In order to compensate for the evaporative loss, an appropriate excess of some component metals (5 wt.% for La, Y and 8% for Mn) was added, respectively.

2.2. Structure Characterization

The annealed alloys were crushed mechanically by hand in an agate mortar, and sieved through a 200–300 size mesh for X-ray diffraction (XRD) measurements, 50 size mesh for neutron powder diffraction (NPD) measurements and 200–300 size mesh for alloy electrodes. XRD measurements were performed on a Bruker D8 Advance diffractometer (Bruker Corporation, Karlsruhe, Germany) with Cu radiation and a power of 40 kV × 40 mA. The patterns were recorded over the range from 8° to 120° in 2θ by steps of 0.02°. NPD was analyzed with λ = 1.8846 Å by a step size of 0.02° in the range 12–150° at room temperature. Then, the XRD and NPD collected data were analyzed by the Rietveld method [31] using the Fullprof program (March 2021, Institut Laue-Langevin, Grenoble, France) [32].

2.3. Hydrogen Absorption and Desorption Measurements

Pressure–composition isotherm (P-C-T) tests were performed using a Sieverts-type apparatus (Beijing Nonferrous Metal Research Institute, Beijing, China) at 298 K. Prior to formal measurements, powder samples were evacuated at 373 K and 1 × 10−4 Pa for at least 2 h in a resistance furnace to remove the impurities. In order to ensure that the sample was completely activated, the sample was hydrided under 8 MPa for 2 h and dehydrided at 0.001 MPa for approximately 2 h, several times.

2.4. Electrochemical Measurements

Alloy electrodes were prepared by cold pressing the mixture of alloy power and carbonyl nickel powder with the weight ratio of 1:3 under 20 MPa pressure to form a pellet of 10 mm in diameter. Electrochemical measurements were performed at 298 K in a standard open tri-electrode electrolysis cell, in which the alloy electrode was used as the working electrode, the Ni(OH)2/NiOOH electrode as the counter electrode, the Hg/HgO electrode as the reference electrode and KOH solution (6 M) as the electrolyte. Each electrode was discharged to cut-off potential −0.6 V Vs. Hg/HgO reference electrode.
Alloy electrodes were charged/discharged at 100 mA g−1 when activated and examined for cyclic stability. The cyclic stability was identified by the capacity retention after the 100th cycle with the following equation [33]:
S 100 = C 100 C max × 100 %
where C100 and Cmax were the discharge capacity at the 100th cycle and the maximum discharge capacity, respectively.
To analyze the corrosion behaviors of the alloy electrodes, a corrosion polarization curve (Tafel polarization curve) was acquired on a CHI660A electrochemical work station (Shanghai Chenhua Instrument Co., Ltd., Shanghai, China) after the alloy electrodes were activated. The Tafel polarization curves were measured by scanning the electrode potential at a rate of 1 mV s−1 from −250 to 250 mV (vs. open circuit potential) at 50% depth of discharge (DOD).

3. Results and Discussion

3.1. Crystal Structure

Figure 1 and Figure 2 show the XRD and the Rietveld refinement pattern of Y0.75La0.25Ni3.5−xMnx (x = 0, 0.05, 0.1, 0.15, 0.2, 0.3) annealed alloys; crystallographic parameters obtained by the Rietveld whole pattern fitting method are tabulated in Table 1. It can be seen that Y0.75La0.25Ni3.5 alloy consists of complex phases, which can be identified as PuNi3-type (SG: R-3m), Ce2Ni7-type (Space group: P63/mmc), Gd2Co7-type (Space group: R-3m), Ce5Co19-type (Space group: R-3m) and CaCu5-type phase (Space group: P6/mmm). When x = 0.05, a small amount of Mn substitution for Ni does not change the phase structure of the alloys, but the Gd2Co7-type phase decreases while the Ce2Ni7- and CaCu5-type phases increase. With further increasing Mn content, the Gd2Co7-type phase transforms to the Ce2Ni7-type phase gradually, the Gd2Co7-type phase continues to decrease and the Ce2Ni7-type phase becomes the main phase of the alloys, which is mainly ascribed to the atomic radius of Mn (RMn = 1.79 Å) being larger than that of Ni (RNi = 1.62 Å). R2Ni7 compounds are polymorphic; they crystallize either in the P63/mmc (2H) or R-3m (3R) space group. Buschow et al. [34] found that the crystal structure of rare-earth nickel-based R2Ni7-type alloys depended on their R atomic radius: when the R atomic radius is large, the R2Ni7 alloy can easily form a 2H-Ce2Ni7 structure, while, when the R atomic radius is small, the R2Ni7 alloy is prone to form a 3R-Gd2Co7 structure. When Mn = 0.1 and 0.15, the abundance of Ce2Ni7-type phase in the alloy reaches 84% and 93%, respectively. As Mn further increases, the PuNi3-type phase increases and Ce2Ni7-type phase abundance decreases again. Therefore, if the content of Mn is too low, the Gd2Co7-type phase exists in the alloy; on the contrary, if the Mn content is too high, the PuNi3-type phase increases. When the content of Mn is between 0.1 and 0.15, the alloys have an approximately single-phase structure, where the abundance of the Ce2Ni7-type phase reaches more than 80%. The variation trend of phase abundance as a function of the Mn content of alloys is presented in Figure 3.
Figure 4 shows the lattice constant variation of each phase as a function of the Mn content of the Y0.75La0.25Ni3.5−xMnx (x = 0, 0.05, 0.1, 0.15, 0.2, 0.3) annealed alloys. The atomic radius of Mn (RMn = 1.79 Å) is larger than that of Ni (RNi = 1.62 Å), so the unit cell volume of each phase becomes larger after Mn is added. The PuNi3-type phase increases from 521.86 Å3 (x = 0) to 534.71 Å3 (x = 0.3), Ce2Ni7-type phase increases from 518.97 Å3 (x = 0) to 529.65 Å3 at (x = 0.3), and Gd2Co7-type phase increases from 781.23 Å3 (x = 0) to 786.22 Å3 at (x = 0.1), respectively. It is worth nothing that the lattice constant a of the three phases increases, whereas the lattice constant c changes differently; the c value increases in the Ce2Ni7-type type phase, and it remains basically unchanged in the Gd2Co7-type phase, but decreases in the PuNi3-type phase. This may be related to the selective occupation in the space lattice of Mn atoms in each phase.
Because the atomic numbers of Ni (25) and Mn (28) are relatively close, it is difficult to accurately distinguish the atomic occupancy of the two elements by XRD. The neutron scattering length of Mn is negative (−0.373 × 10−12 cm) while that of Ni is positive (+1.03 × 10−12 cm) [35]; there is a large difference between them. Therefore, neutron diffraction can easily characterize the occupancy of Ni and Mn in B-site elements of La-Y-Ni-Mn hydrogen storage alloys. Figure 5 shows the Rietveld refinement of the NPD pattern for Y0.75La0.25Ni3.2Mn0.3 alloys, Table 2 lists the crystal structure data of the PuNi3-type (Table 2(a)) and Ce2Ni7-type phases (Table 2(b)), and Figure 6 presents the schematic diagram of the atomic stacking along the c-axis of the two phases. The results show that Mn occupies the 6c site in the [AB5] subunit and a small amount of Mn occupies the 18h site of the interface between the [A2B4] and [AB5] subunits of the PuNi3-type phase, while Mn only occupies the 4e and 6h sites in the [AB5] subunit of the Ce2Ni7-type phase.
The superlattice structure is described as stacking structures made of [A2B4] and [AB5] subunits piled along the c-axis. Their general formula, summarized by Khan [36], can be defined as y = (5n + 4)/(n + 2) (where n is the number of [AB5] subunits). For A2B7 alloy, n is equal to 2 and, consequently, the basic period can be defined as [A2B4] + 2[AB5]. These phases are polymorphs as they crystallize either in hexagonal (2H, Ce2Ni7-type) or rhombohedral (3R, Gd2Co7-type) symmetry. As long as the volume of the [A2B4] subunit is larger than that of the [AB5] subunit, the [A2B4] subunit remains more active with respect to hydrogenation [37]. In other words, for superlattice structure hydrogen storage alloys, the larger the volume mismatch of V[A2B4]/V[AB5] is, the more obvious the “anisotropic” expansion of the lattice will be during hydrogen absorption/desorption. The structural evolution of the hydrogen absorption/desorption of La2Ni7 [17] alloy is the most typical example. Figure 7 shows the variation of the [A2B4] and [AB5] subunits as a function of the Mn content of the Ce2Ni7-type phase. It is clear that the volume of the [AB5] subunit increases almost linearly with increasing Mn content, from 85.22 Å3 (x = 0) to 87.72 Å3 (x = 0.3), while the volume of the [A2B4] subunit is almost unchanged; the results also indicate that Mn mainly occupies Ni sites in the [AB5] subunit, and this is in good agreement with the atomic occupancy (Table 2) obtained from NPD data. In addition, V[A2B4]/V[AB5] decreases gradually, from 1.045 (x = 0) to 1.019 (x = 0.3). This is mainly due to the atomic radius of Mn being larger than that of Ni, and its selective occupation in the superlattice structure. The smaller the volume mismatch in V[A2B4]/V[AB5] is, the better the structural stability of the alloys will be during hydrogen absorption/desorption [38,39].

3.2. Hydrogen Storage Characteristics

Figure 8 shows the P-C isotherm curves (gaseous hydrogen absorption/desorption) of Y0.75La0.25Ni3.5−xMnx (x = 0, 0.05, 0.1, 0.15, 0.2, 0.3) annealed alloy at 298 K. The basic data in the process of the hydrogen absorption/desorption of alloys are presented in Table 3. It can be found that all alloys are activated within three times; with the increase in Mn content, the hydrogen capacity and plateau pressure during hydrogen absorption/desorption change, which may be closely related to the phase structures and the corresponding unit cell volume of the alloys. This will be discussed in detail below.
With the increase in Mn content, the unit cell volumes of all phases of La-Y-Ni-Mn alloys increase, which leads to a decrease in the plateau pressure during hydrogen absorption/desorption [10,40]. The hydrogen absorption plateau pressure decreased from 0.67 MPa (x = 0) to 0.02 (x = 0.3), and the hydrogen desorption plateau pressure decreased from 0.39 MPa (x = 0) to 0.01 MPa (x = 0.3). This shows that Mn can effectively reduce the plateau pressure of the hydrogen absorption/desorption of La-Y-Ni system hydrogen storage alloys, which is consistent with the current research [15,41]. At the same time, with the rise of x, the hydrogen absorption/desorption plateau of the alloy becomes flatter and wider, indicating that the addition of Mn can effectively promote the hydrogen absorption and desorption properties of La-Y-Ni system alloys.
The maximum hydrogen absorption capacity of series alloys is 1.26 wt.%–1.44 wt.% under 8 MPa. When the Mn content increases, the maximum hydrogen absorption of the alloy increases first and then decreases, but the overall trend is increasing. It increases from 1.34 wt.% (x = 0) to 1.44 wt.% (x = 0.15) and then decreases to 1.40 wt.% (x = 0.3); the increase in the maximum hydrogen absorption capacity (x = 0–0.15) may be mainly due to the Mn effectively reducing the hydrogen absorption/desorption plateau pressure of the alloys, while the decrease in the maximum hydrogen absorption capacity (x = 0.15–0.3) is related to the fact that the hydrogen absorption/desorption plateau pressure of the alloy is so low that the hydrogen cannot be released. All alloys have relatively obvious hydrogen absorption/desorption plateaus, indicating that there is no serious hydrogen-induced amorphization during the hydrogen absorption/desorption process of the alloys. When the Mn content is low (x < 0.2), the alloys have multiple hydrogen absorption/desorption plateaus, which is mainly ascribed to the alloys having a multiphase structure, and the unit cell volume of each phase is different. When the Mn content is high (x ≥ 0.2), the alloys have only two types of phases, PuNi3-type and Ce2Ni7-type, and the unit cell volumes of the two types of phases are almost the same, so the alloys have only one hydrogen absorption/desorption plateau.

3.3. Discharge and Cyclic Properties

Figure 9 shows the electrochemical P-C isotherm curves (desorption) of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloy electrodes at 298 K. It is found that the variation law of the electrochemical P-C isotherm curves of the alloy electrodes is in good agreement with the P-C isotherm curves of the alloy gaseous hydrogen desorption between 10−3 and 10−1 MPa. Except for x = 0.05, the hydrogen desorption plateau pressure of the alloy gradually decreases with the increase in Mn content. Figure 10 shows the 100 charge/discharge cycles curve of the alloy electrode at a charge/discharge current density of 100 mA·g−1. Figure 11 presents the relationship between the maximum discharge capacity and cyclic stability of the Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) alloy electrode as a function of Mn content. Table 4 summarizes the electrochemical performance of the alloy electrodes. It can be seen that all of the alloy electrodes exhibit good activation properties; except for x = 0 and 0.05, the alloy electrodes were fully activated within five times, and the activation times of x = 0.05 and 0.1 alloy electrodes exceeded three times, but all reached 90% of the maximum discharge capacity within three times. With the increase in x, the maximum discharge capacity of the alloy electrodes first increases and then decreases, from 231.9 mAh·g−1 (x = 0) to 367.4 mAh·g−1 (x = 0.15), and then decreases to 334.4 mAh·g−1 (x = 0.3).
From the gaseous P-C-T curves of the alloys previously discussed, it can be seen that with the increase in x, the maximum discharge capacity of the alloy electrodes first increases and then decreases. When x = 0–0.15, the gradual increase in discharge capacity is due to the increasing unit cell volume of each phase after the addition of Mn, which effectively reduces the hydrogen absorption and desorption plateau pressure of the alloy electrodes; when x = 0.15–0.3, the gradual decrease in discharge capacity may be mainly due to the fact that the hydrogen absorption/desorption plateau pressure of the alloy was too low, which made it difficult to release hydrogen. The hydrogen absorption/desorption plateau pressure between 0.01 and 0.1 MPa is most suitable for the alloy electrodes to absorb and desorb hydrogen. When the plateau pressure is too high, the hydrogen atoms will not be absorbed by the alloys, and hydrogen will be released; on the other hand, when the plateau pressure is too low, the form of hydrogen in the hydride will change, and the discharge voltage and battery capacity will be reduced.
The capacity degradation of alloy electrodes is mainly ascribed to structural failure (hydrogen-induced amorphization, hydrogen-induced phase transformation, hydrogen-induced defects, stress-strain, etc.) and the electrochemical corrosion of active materials during cycling. Since Mn is mainly distributed in the [AB5] subunit, with the increase in Mn content, the mismatch between the [A2B4] and [AB5] subunits is reduced; meanwhile, the structural stability of the alloy electrodes during electrochemical cycling is enhanced. At the same time, in general, the micro-strain resulting from the lattice expansion/contraction can lead to the pulverization of the alloy particles during the hydrogenation/dehydrogenation cycles; as a result, more and more fresh surfaces of alloy particles are directly exposed to alkaline electrolytes, which will accelerate the corrosion of the alloy electrodes [42]. According to this study, with the increase in Mn, the mismatch between the [A2B4] and [AB5] subunits, micro-strain, and pulverization of alloys will all decrease, which reduces the corrosion of the alloys and improves the cycle stability of the alloy electrodes. Figure 12 shows the Tafel polarization curves of the alloy electrodes. It can be seen that with the increase in Mn content, the corrosion current density (icorr) decreased from 10.1 mA·cm−2 (x = 0) to 3.85 mA·cm−2 (x = 0.3) and the corrosion potential (Ecorr) decreased from −0.934 V (x = 0) to −0.925 V (x = 0.3), respectively. These results show that the addition of the Mn element reduces the corrosion behavior of alloy electrodes; this is consistent with the previous discussion. Hence, Mn-containing alloys will promote the cyclic stability of the alloy electrodes. From x = 0.1 to x = 0.3, the cyclic stability S100 of the alloy electrodes did not change significantly, all between 74.88% and 77.42%, which indicates that the crystalline structure of the alloy has been maintained well in the electrochemical charge/discharge process since x = 0.1.

4. Conclusions

In this paper, the structures and properties of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) alloys were systematically investigated. Some conclusions can be summarized as follows:
(1)
Y0.75La0.25Ni3.5 alloy consisted of complex phases: PuNi3-type, Ce2Ni7-type, Gd2Co7-type, Ce5Co19-type and CaCu5-type phase. A small amount of Mn substitution for Ni does not change the phase structure of the alloys. With the increase in Mn content, the Gd2Co7-type phase turns into the Ce2Ni7-type phase. When the value of Mn is between 0.1 and 0.15, the single-phase property of the alloys is better, and the phase abundance of Ce2Ni7-type reaches more than 80%. Mn mainly occupies Ni sites in the [AB5] subunit and the interface between the [AB5] and [A2B4] subunits.
(2)
The maximum hydrogen absorption capacity of series alloys is 1.260 wt.%–1.444 wt.% under 8 MPa. When the Mn content increases, the maximum hydrogen absorption of the alloy increases first and then decreases. It increases from 1.339 wt.% (x = 0) to 1.444 wt.% (x = 0.15) and then decreases to 1.404 wt.% (x = 0.3). The hydrogen desorption plateau pressure of the alloys gradually decreases, and the hydrogen absorption/desorption plateau pressure of the alloy becomes flatter and wider. The addition of Mn effectively improves the hydrogen absorption/desorption performance of the series alloys.
(3)
With the increase in Mn content, the maximum discharge capacity of the alloy electrodes first increased and then decreased, from 231.9 mAh·g−1 (x = 0) to 367.4 mAh·g−1 (x = 0), and then decreased to 334.4 mAh·g−1 (x = 0). The maximum discharge capacity of the alloy electrodes is closely related to its hydrogen storage capacity at 0.1 MPa and its hydrogen absorption/desorption plateau pressure. The cyclic stability of all the Mn-containing alloy electrodes was improved distinctly compared to that of Mn-free alloy electrodes, because the volume mismatch between the [A2B4] and [AB5] subunits of each phase for series alloys became smaller after the addition of Mn, which improved the structural stability and reduced the corrosion of the alloys during the hydrogen absorption/desorption cycles.
In summary, for Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) alloys, there is an optimum Mn substitution (x = 0.1–0.15) for Ni in terms of the structures, hydrogen storage behaviors, as well as the electrochemical properties of the alloys, in which the Ce2Ni7-type phase abundance is higher and the overall performance is better.

Author Contributions

A.D. and Y.L. conceived and designed the experiments; A.D. performed the experiments; A.D., J.Z., Y.X., Y.Y., X.K., B.S. and H.Z. analyzed the data; A.D. wrote the draft and Y.L. revised it. All authors have read and agreed to the published version of the manuscript.

Funding

This work was funded by the National Natural Science Foundation of China (No. 51761026), Ningxia Natural Science Fund (No. 2022AAC03048 and 2019AAC03003) and the fund of the State Key Laboratory of Advanced Processing and Recycling of Non-Ferrous Metals (No. SKLAB02019004), Lanzhou University of Technology.

Conflicts of Interest

The authors declare no conflict of interest.

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Metals 12 01122 g001 550
Figure 1. XRD pattern of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys: (a) whole pattern; (b) local pattern.
Figure 1. XRD pattern of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys: (a) whole pattern; (b) local pattern.
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Metals 12 01122 g002 550
Figure 2. Rietveld refinement pattern of the XRD data for Y0.75La0.25Ni3.5−xMnx (x = 0, 0.05, 0.1, 0.15, 0.2, 0.3) alloys: (a) x = 0; (b) x = 0.05; (c) x = 0.1; (d) x = 0.15; (e) x = 0.2; (f) x = 0.3.
Figure 2. Rietveld refinement pattern of the XRD data for Y0.75La0.25Ni3.5−xMnx (x = 0, 0.05, 0.1, 0.15, 0.2, 0.3) alloys: (a) x = 0; (b) x = 0.05; (c) x = 0.1; (d) x = 0.15; (e) x = 0.2; (f) x = 0.3.
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Metals 12 01122 g003 550
Figure 3. Phase abundance as a function of Mn content of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys.
Figure 3. Phase abundance as a function of Mn content of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys.
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Metals 12 01122 g004 550
Figure 4. Lattice constant of PuNi3-type (a), Ce2Ni7-type (b), Gd2Co7-type (c)phase as a function of Mn content for Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys.
Figure 4. Lattice constant of PuNi3-type (a), Ce2Ni7-type (b), Gd2Co7-type (c)phase as a function of Mn content for Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys.
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Metals 12 01122 g005 550
Figure 5. Rietveld refinement of the NPD pattern for Y0.75La0.25Ni3.2Mn0.3 alloys.
Figure 5. Rietveld refinement of the NPD pattern for Y0.75La0.25Ni3.2Mn0.3 alloys.
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Metals 12 01122 g006 550
Figure 6. Representative structures of 3R-type PuNi3 (a) and 2H-type Ce2Ni7 (b) stacked by [A2B4] and [AB5] subunits along the c-axis in La-Y-Ni system hydrogen storage alloys with superlattice structure.
Figure 6. Representative structures of 3R-type PuNi3 (a) and 2H-type Ce2Ni7 (b) stacked by [A2B4] and [AB5] subunits along the c-axis in La-Y-Ni system hydrogen storage alloys with superlattice structure.
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Figure 7. Subunit volume of Ce2Ni7-type phase for Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys.
Figure 7. Subunit volume of Ce2Ni7-type phase for Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys.
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Figure 8. P-C isotherm curves (absorption and desorption) of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys at 298 K.
Figure 8. P-C isotherm curves (absorption and desorption) of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys at 298 K.
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Metals 12 01122 g009 550
Figure 9. Electrochemical P–C isotherm curves (desorption) of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys at 298 K.
Figure 9. Electrochemical P–C isotherm curves (desorption) of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys at 298 K.
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Metals 12 01122 g010 550
Figure 10. Cyclic stability curves of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) alloy electrodes.
Figure 10. Cyclic stability curves of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) alloy electrodes.
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Metals 12 01122 g011 550
Figure 11. Maximum discharge capacity and capacity retention ratio after 100 cycles of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) alloy electrodes at 298 K.
Figure 11. Maximum discharge capacity and capacity retention ratio after 100 cycles of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) alloy electrodes at 298 K.
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Metals 12 01122 g012 550
Figure 12. Tafel polarization curves of Y0.75La0.25Ni3.5−xMnx (x = 0, 0.1, 0.3) alloy electrodes.
Figure 12. Tafel polarization curves of Y0.75La0.25Ni3.5−xMnx (x = 0, 0.1, 0.3) alloy electrodes.
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Table
Table 1. Crystallographic data of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys.
Table 1. Crystallographic data of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloys.
SamplePhaseSpace GroupPhase AbundanceLattice Constant
(x) (%)a(Å)c(Å)V(Å3)c/a
x = 0PuNi3R-3m24(1) 4.934(3) 24.755(2) 521.86(4) 5.017
Ce2Ni7P63/mmc19(1) 4.978(1) 24.166(1) 518.97(10) 4.853
Gd2Co7R-3m37(1) 4.975(1) 36.446(2) 781.23(6) 7.326
Ce5Co19R-3m17(1) 4.977(1) 48.327(4) 1036.8(2) 9.710
CaCu5P6/mmm3(1) 4.922(1) 3.9756(1) 83.407(1) 0.808
x = 0.05PuNi3R-3m11(1) 4.956(1) 24.736(1) 526.07(18) 4.992
Ce2Ni7P63/mmc28(1) 4.986(1) 24.202(1) 521.08(6) 4.854
Gd2Co7R-3m14(1) 4.986(1) 36.425(1) 784.09(6) 7.306
Ce5Co19R-3m18(1) 4.968(2) 48.230(3) 1031.0(1) 9.708
CaCu5P6/mmm29(1) 4.927(2) 3.9722(1) 83.523(1) 0.806
x = 0.1PuNi3R-3m6(1) 4.974(1) 24.682(2) 528.83(7) 4.962
Ce2Ni7P63/mmc84(1) 4.993(1) 24.241(5) 523.42(2) 4.855
Gd2Co7R-3m10(1) 4.993(1) 36.418(29) 786.22(10) 7.294
x = 0.15PuNi3R-3m7(1) 5.008(1) 24.581(14) 533.88(4) 4.908
Ce2Ni7P63/mmc93(1) 4.997(1) 24.287(3) 525.15(1) 4.861
x = 0.2PuNi3R-3m30(1) 5.018(1) 24.532(1) 534.93(2) 4.889
Ce2Ni7P63/mmc70(1) 5.003(1) 24.328(1) 527.33(2) 4.863
x = 0.3PuNi3R-3m38(1) 5.020(1) 24.500(6) 534.71(2) 4.880
Ce2Ni7P63/mmc62(1) 5.010(1) 24.369(1) 529.65(1) 4.864
Table
Table 2. a. Crystal structure data of PuNi3-type phase for Y0.75La0.25Ni3.2Mn0.3 annealed alloys. b. Crystal structure data of Ce2Ni7-type phase for Y0.75La0.25Ni3.2Mn0.3 annealed alloys.
Table 2. a. Crystal structure data of PuNi3-type phase for Y0.75La0.25Ni3.2Mn0.3 annealed alloys. b. Crystal structure data of Ce2Ni7-type phase for Y0.75La0.25Ni3.2Mn0.3 annealed alloys.
(a)
SiteAtomxyzBiso(Å2)Occupancy
3aLa0001.9(2)0.28(4)
Y0001.9(2)0.72(4)
6cLa000.1418(7)0.4(1)0.11(2)
Y000.1418(7)0.4(1)0.89(2)
3bNi000.50.6(1)1
6cNi000.3317(9)2.5(2)0.75(3)
Mn000.3317(9)2.5(2)0.25(3)
18hNi0.4999(12)−0.4999(12)0.0830(4)0.3(1)0.96(1)
Mn0.4999(12)−0.4999(12)0.0830(4)0.3(1)0.04(1)
(b)
SiteAtomxyzBiso(Å2)Occupancy
4f1Y000.0282(6)2.7(2)1
4f2La000.1717(6)1.7(2)0.58(5)
Y000.1717(6)1.7(2)0.42(5)
2aNi0001.9(2)1
4eNi000.1671(8)1.1(2)0.77(6)
Mn000.1671(8)1.1(2)0.23(6)
4fNi0.66670.33330.1666(6)0.2(2)1
6hNi0.1621(6)0.3241(12)0.751.6(1)0.68(5)
Mn0.1621(6)0.3241(12)0.751.6(1)0.32(5)
12kNi0.1689(7)0.3377(14)0.9156(3)1.2(1)1
Biso: Temperature factor; Occupancy: Atomic occupation.
Table
Table 3. Hydrogen absorption/desorption characteristics of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloy.
Table 3. Hydrogen absorption/desorption characteristics of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) annealed alloy.
Sample (x)ActivationPlateau PressureHydrogen CapacityHydrogen Capacity
Time(MPa)8 MPa0.1 MPa
NaAbs.Des.(wt.%)(wt.%)
x = 0.0030.05/0.670.01/0.391.340.44
x = 0.0530.25/0.690.04/0.421.260.38
x = 0.1030.03/0.020.01/0.131.420.79
x = 0.1530.03/0.130.02/0.101.440.87
x = 0.2030.030.021.401.16
x = 0.3030.020.011.401.21
Abs: Absorption; Des: Desorption.
Table
Table 4. Electrochemical properties of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) alloy electrodes.
Table 4. Electrochemical properties of Y0.75La0.25Ni3.5−xMnx (x = 0–0.3) alloy electrodes.
Sample (x)NaCmaxS100icorrEcorr
(mAh‧g−1)(%)(mA‧cm−2)(V)
x = 03231.967.110.1−0.934
x = 0.059245.884.0
x = 0.19315.475.59.53−0.93
x = 0.153367.474.9
x = 0.23336.374.1
x = 0.35334.477.43.85−0.925
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Deng, A.; Luo, Y.; Zhou, J.; Xie, Y.; Yuan, Y.; Kang, X.; Shen, B.; Zhang, H. Effect of Mn Element on the Structures and Properties of A2B7-Type La–Y–Ni-Based Hydrogen Storage Alloys. Metals 2022, 12, 1122. https://doi.org/10.3390/met12071122

AMA Style

Deng A, Luo Y, Zhou J, Xie Y, Yuan Y, Kang X, Shen B, Zhang H. Effect of Mn Element on the Structures and Properties of A2B7-Type La–Y–Ni-Based Hydrogen Storage Alloys. Metals. 2022; 12(7):1122. https://doi.org/10.3390/met12071122

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Deng, Anqiang, Yongchun Luo, Jianfei Zhou, Yunding Xie, Yuan Yuan, Xiaoyan Kang, Bingjin Shen, and Haimin Zhang. 2022. "Effect of Mn Element on the Structures and Properties of A2B7-Type La–Y–Ni-Based Hydrogen Storage Alloys" Metals 12, no. 7: 1122. https://doi.org/10.3390/met12071122

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