# Structural Evolution and Hydrogen Sorption Properties of YxNi2-yMny (0.825 ≤ x ≤ 0.95, 0.1 ≤ y ≤ 0.3) Laves Phase Compounds

Hao Shen, Valerie Paul-Boncour, Ping Li, Lijun Jiang, Junxian Zhang

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Hao Shen, Valerie Paul-Boncour, Ping Li, Lijun Jiang, Junxian Zhang. Structural Evolution and Hydrogen Sorption Properties of YxNi2-yMny (0.825 ≤ x ≤ 0.95, 0.1 ≤ y ≤ 0.3) Laves Phase Compounds. Inorganics, 2024, 12 (2), pp.55. 10.3390/inorganics12020055. [https://doi.org/10.3390/inorganics12020054]

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Article

# Structural Evolution and Hydrogen Sorption Properties of YₓNi₂₋ᵧMnᵧ (0.825 ≤ x ≤ 0.95, 0.1 ≤ y ≤ 0.3) Laves Phase Compounds

Hao Shen <fnref n="1" />,<fnref n="2" />,<fnref n="3" />, Valerie Paul-Boncour <fnref n="1" />, Ping Li <fnref n="2" />, Lijun Jiang <fnref n="3" />,* and Junxian Zhang <fnref n="1" />,*

**Abstract:** The YₓNi₂₋ᵧMnᵧ system was investigated in the region 0.825 ≤ x ≤ 0.95, 0.1 ≤ y ≤ 0.3. The alloys were synthesized by induction melting and corresponding annealing. The substitution of Mn for Ni (y = 0.1) favors the formation of a C15 structure with disordered Y vacancies against the superstructure of Y₀.₉₅Ni₂. For y = 0.2 and 0.3, Mn can substitute in both Y and Ni sites. Single-phase compounds with a C15 structure can be formed by adjusting both the Y and Mn contents. Their hydrogen absorption-desorption properties were measured by pressure-composition isotherm (PCI) measurements at 150 °C, and the hydrides were characterized at room temperature by X-ray diffraction and TG-DSC experiments. The PCIs show two plateaus corresponding to the formation of crystalline and amorphous hydrides. The heating of the amorphous hydrides leads to an endothermic desorption at first and then a recrystallization into Y(Ni, Mn)₃ and YHₓ phases. At higher temperatures, the Y hydride desorbs, and a recombination into a Y(Ni, Mn)₂ Laves phase compound is observed. For y = 0.1, vacancy formation in the Y site and partial Mn substitution in the Ni site enhance the structural stability and suppress the hydrogen-induced amorphization (HIA). However, for a larger Mn content (y ≥ 0.2), Mn substitutes also in the Y sites at the expense of Y vacancies. This yields worse structural stability upon hydrogenation than for y = 0.1, as the mean ratio r(Y, Mn)/r(Ni/Mn) becomes larger than for y = 0.1 r(Y, □)/r(Ni/Mn).

**Keywords:** Y-based alloys; Mn substitution; structural properties; hydrogen-induced amorphization

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**Citation:** Shen, H.; Paul-Boncour, V.; Li, P.; Jiang, L.; Zhang, J. Structural Evolution and Hydrogen Sorption Properties of YₓNi₂₋ᵧMnᵧ (0.825 ≤ x ≤ 0.95, 0.1 ≤ y ≤ 0.3) Laves Phase Compounds. *Inorganics* **2024**, *12*, 55. https://doi.org/10.3390/inorganics12020055

Academic Editors: Rainer Niewa, Torben R. Jensen, Richard Dronskowski, Christian M. Julien, Guido Kickelbick, Alexander S. Novikov, Gary Hix and Hans-Conrad zur Loye

Received: 24 November 2023

Revised: 19 January 2024

Accepted: 4 February 2024

Published: 7 February 2024

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**Copyright:** © 2024 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (https://creativecommons.org/licenses/by/4.0/).

## 1. Introduction

Hydride-forming ABₙ-type alloys (A = rare earth, B = transition metal) are widely used for hydrogen storage [1], as well as negative electrode materials for nickel-metal hydride (Ni–MH) batteries [2]. Recently, (A, Mg)Ni-based ABₙ alloys (3 ≤ n ≤ 4) with stacking structures have been developed for hydrogen storage [3,4], with the capacity increasing from 1.5 wt.% for AB₅-type alloys to 1.8 wt.% for ABₙ (3 ≤ n ≤ 4)-type alloys [5–8]. AB₂-type alloys with a C15 Laves phase structure possess more tetrahedral sites than AB₅ and ABₙ alloys (3 ≤ n ≤ 4) to accommodate hydrogen atoms; thus, a higher hydrogen absorption capacity is foreseen [9].

Laves phases belong to the class of Frank–Kasper phases showing topologically close-packed structures. The closest packing of hard spheres can be obtained in C15 AB₂ alloys, where the A and B atoms show an ideal radius ratio r<sub>A</sub>/r<sub>B</sub> = 1.225. Hydrogen storage in AB₂ Laves phase compounds has been investigated and is still a hot research subject as it presents a larger hydrogen capacity compared with AB₅ and bcc compounds. Many studies have been performed on ZrB₂ compounds (B = V, Cr, Mn, Fe, Co), whose thermodynamic properties can be adjusted by chemical substitution and which can absorb up to 2 wt.% of hydrogen [10–15]. YB₂ hydrides (B = Mn, Fe, Co, Ni) were also widely studied for

<footnote n="1">¹ Univ. Paris Est Creteil, CNRS, ICMPE, UMR 7182, 2 Rue Henri Dunant, 94320 Thiais, France</footnote>

<footnote n="2">² Institute for Advanced Materials and Technology, University of Science and Technology, Beijing 100083, China</footnote>

<footnote n="3">³ National Engineering Research Center of Nonferrous Metals Materials and Products for New Energy, GRINM Group Co., Ltd., Beijing 100088, China</footnote>

<footnote>* Correspondence: jlj@grinm.com (L.J.); junxian.zhang@cnrs.fr (J.Z.)</footnote>
their hydrogenation properties [16,17]. However, depending on the transition metal and the hydrogenation conditions, they can form either crystalline or amorphous hydrides. Disproportionation into transition metal and Y hydride can also occur. In a recent review, all the criteria concerning the AB₂ hydride stability were investigated, and it appears that one critical geometric factor is the rA/rB ratio, where rA and rB are the atomic radii of the A and B elements, respectively [18]. When this ratio becomes larger than 1.37, hydrogen-induced amorphization (HIA) is favored at the detriment of the hydrogen sorption reversibility. To reduce this ratio towards the ideal ratio rA/rB = 1.225 for topologically close-packed structures, it is necessary to either decrease rA, increase rB, or play on both elements by appropriate substitutions. Several studies have already been performed on the role of chemical substitutions in the hydrogenation properties of Y0.95Ni2. As Y has a smaller weight than lanthanides, higher weight capacities are expected. Y0.95Ni2 crystallizes in a superstructure with a doubling of the cubic lattice parameter a and a lower symmetric space group Fd3m (No. 216) compared with the cubic Laves phase (Fd3m space group). The superstructure is related to the formation of ordered Y vacancies in the 4a sites and has also been observed in other ANi2 compounds [19,20]. Hydrogen absorption in Y0.95Ni2 yields a crystalline hydride up to 2.7 H/f.u. with the same superstructure type as the parent compound, whereas hydrogen-induced amorphization (HIA) occurs for a larger H content [19,21,22]. The substitution of Y by Sc, which is isoelectronic and has a smaller size than Y, allowed for the reduction of rA, and HIA was fully eliminated in the Y0.25Sc0.7Ni2 compound [23]. On the other hand, the influence of Al for Ni substitution in Y0.95Ni2 was investigated [21], and for an appropriate Al content, the reversibility was improved.

Another possibility to increase r<sub>B</sub> is to substitute Ni by a larger Mn atom. Among various Y<sub>2</sub>B<sub>2</sub> compounds, YMn<sub>2</sub> is a C15 Laves phase adopting the cubic MgCu<sub>2</sub>-type structure (space group Fd&#773;3m, No. 227). At room temperature, the YMn<sub>2</sub> hydride remains cubic up to 3.5 H/f.u. and undergoes rhombohedral distortion for a larger H content with a progressive cell volume increase [24]. Upon cooling, various structural distortions are observed, depending on the H content [16,25]. For a hydrogen pressure above 0.5 GPa, a complex YMn<sub>2</sub>H<sub>6</sub> hydride, with a fluorite-type structure similar to that of Mg<sub>2</sub>FeH<sub>6</sub>, is obtained. This structure is different from that of the parent compound with two different sites for the Mn atoms, one sharing the same position as the Y atom, with a disordered substitution, and the other one forming MnH<sub>6</sub> octahedra [26,27]. Furthermore, a recent work by machine learning shows that Mn contributes to increasing the hydrogen storage capacity of AB<sub>2</sub> compounds [28].

In a recent study [29], the structural properties of Y<sub>x</sub>Ni<sub>2-y</sub>Mn<sub>y</sub> intermetallic compounds were investigated by both experimental work (X-ray and neutron diffraction) and theoretical DFT calculations in order to determine the homogeneity range of Laves phase compounds. It was observed that the rate of Y vacancies varies with the Mn content and that Mn can substitute not only on the Ni but also in the Y site.

This previous work has been completed by the synthesis and characterization of new Y<sub>x</sub>Ni<sub>2-y</sub>Mn<sub>y</sub> compounds (y ≤ 0.3), and will be detailed in the first section of this paper.

In addition, the hydrogenation properties of the selected Y<sub>x</sub>Ni<sub>2-y</sub>Mn<sub>y</sub> (0.825 ≤ x ≤ 0.95, 0.1 ≤ y ≤ 0.3) sample forming single-phase compounds have been investigated. Their thermodynamic and structural properties have been studied by measuring pressure-composition isotherms (PCIs), X-ray diffraction patterns, and thermal properties. The influence of the Mn substitution on the metal hydride stability, from both thermodynamic and crystallographic points of view, will be presented and discussed.

## 2. Results

### 2.1. Structural Characteristics of Y<sub>x</sub>Ni<sub>2-y</sub>Mn<sub>y</sub> Compounds

The Y0.95Ni2 compound crystallizes in a single TmNi2-type superstructure type with a doubling of the cubic cell parameter and vacancies in the Y site (space group Fd3m, No. 216) [29–31]. The Y0.9Ni2-yMn-y (0.1 ≤ y ≤ 0.3) alloys contain Y(Ni, Mn)2 with a C15 structure (space group Fd3m, No. 227) as the main phase, as previously observed for
Al-substituted AB₂ compounds [21]. YNi (AB-type, space group *Pnma*, No. 62) appears as a secondary phase, and its abundance increases proportionally with the Mn content. The compounds with *y* = 0.4 and 0.5 contain Y(Ni, Mn)₃ (AB₃-type, space group *R3m*, No. 166) as the main phase, whose abundance increases at the expense of the AB₂-type one, and will not be treated in the present paper. Single C15 Laves phases were obtained for Y₀.₉Ni₁.₉Mn₀.₁, Y₀.₈₆Ni₁.₈Mn₀.₂, and Y₀.₈₂₅Ni₁.₇Mn₀.₃ compounds. In the previous work, neutron diffraction analysis and DFT calculation showed that, at a low Mn content, YₓNi₂₋ᵧMnᵧ (x = 0.95, 0.90, 0.86, 0.825, y = 0, 0.1, 0.2, 0.3, 0.4, 0.5), the C15 structure is stabilized by vacancies in the A site. Meanwhile, for a larger Mn content, stabilization occurs through Mn substitution for Y in the A site and a reduction of Y vacancies [30]. The phase abundances and the lattice parameters of YₓNi₂₋ᵧMnᵧ (0.825 ≤ x ≤ 0.95, 0.1 ≤ y ≤ 0.3) obtained by XRD Rietveld analyses are gathered in Table 1. Note that the traces of Y₂O₃ found in the alloys can be either still present in the initial Y ingot, despite surface cleaning, or due to the oxidation of unreacted Y upon thermal treatment.

<table><caption>**Table 1.** Phase compositions determined by electron probe microanalysis (EPMA) and crystallographic data from Rietveld analyses of YₓNi₂₋ᵧMnᵧ (0.825 ≤ x ≤ 0.95, 0.1 ≤ y ≤ 0.3) compounds; * stands for the data that are published in reference [30].</caption><thead><tr><th rowspan="2">Sample</th><th rowspan="2">Phase</th><th rowspan="2">Space Group</th><th rowspan="2">Phase Composition (EPMA) (±0.01)</th><th rowspan="2">Abundance (wt.%)</th><th colspan="3">Lattice Parameters (Å)</th></tr><tr><th>a</th><th>b</th><th>c</th></tr></thead><tbody><tr><td>Y<sub>0.95</sub>Ni<sub>2</sub></td><td>Y<sub>0.95</sub>Ni<sub>2</sub></td><td>F<span style="text-decoration: overline;">4</span>3m</td><td>Y<sub>0.93</sub>Ni<sub>2</sub></td><td>99 (1)</td><td>14.3557<br>(1)</td><td>-</td><td>-</td></tr><tr><td rowspan="2">Y<sub>0.9</sub>Ni<sub>1.9</sub>Mn<sub>0.1</sub> *</td><td>Y<sub>1-v</sub>(Ni, Mn)<sub>2</sub></td><td>F<span style="text-decoration: overline;">d</span>3m</td><td>Y<sub>0.86</sub>Ni<sub>1.87</sub>Mn<sub>0.13</sub></td><td>96 (1)</td><td>7.1733 (1)</td><td>-</td><td>-</td></tr><tr><td>Y<sub>2</sub>O<sub>3</sub></td><td>Ia<span style="text-decoration: overline;">3</span></td><td>-</td><td>4 (1)</td><td>-</td><td>-</td><td>-</td></tr><tr><td rowspan="2">Y<sub>0.9</sub>Ni<sub>1.8</sub>Mn<sub>0.2</sub></td><td>Y<sub>1-v</sub>(Ni, Mn)<sub>2</sub></td><td>F<span style="text-decoration: overline;">d</span>3m</td><td>Y<sub>0.87</sub>Ni<sub>1.79</sub>Mn<sub>0.21</sub></td><td>95 (1)</td><td>7.1900 (1)</td><td>-</td><td>-</td></tr><tr><td>YNi</td><td>Pnma</td><td>-</td><td>4 (1)</td><td>7.131 (2)</td><td>4.1505 (1)</td><td>5.4858 (1)</td></tr><tr><td rowspan="2">Y<sub>0.9</sub>Ni<sub>1.7</sub>Mn<sub>0.3</sub></td><td>Y<sub>2</sub>O<sub>3</sub></td><td>Ia<span style="text-decoration: overline;">3</span></td><td>-</td><td>trace</td><td>-</td><td>-</td><td>-</td></tr><tr><td>Y<sub>1-v</sub>(Ni, Mn)<sub>2</sub></td><td>F<span style="text-decoration: overline;">d</span>3m</td><td>Y<sub>0.83</sub>Ni<sub>1.68</sub>Mn<sub>0.32</sub></td><td>90 (1)</td><td>7.1951 (1)</td><td>-</td><td>-</td></tr><tr><td rowspan="2">Y<sub>0.86</sub>Ni<sub>1.8</sub>Mn<sub>0.2</sub> *</td><td>YNi</td><td>Pnma</td><td>YNi<sub>0.99</sub>Mn<sub>0.02</sub></td><td>9 (1)</td><td>7.134 (2)</td><td>4.1389 (1)</td><td>5.5064 (1)</td></tr><tr><td>Y<sub>2</sub>O<sub>3</sub></td><td>Ia<span style="text-decoration: overline;">3</span></td><td>-</td><td>trace</td><td>-</td><td>-</td><td>-</td></tr><tr><td rowspan="2">Y<sub>0.825</sub>Ni<sub>1.7</sub>Mn<sub>0.3</sub> *</td><td>Y<sub>1-v</sub>(Ni, Mn)<sub>2</sub></td><td>F<span style="text-decoration: overline;">d</span>3m</td><td>Y<sub>0.85</sub>Ni<sub>1.79</sub>Mn<sub>0.21</sub></td><td>99 (1)</td><td>7.1826 (1)</td><td>-</td><td>-</td></tr><tr><td>Y<sub>2</sub>O<sub>3</sub></td><td>-</td><td>-</td><td>trace</td><td>-</td><td>-</td><td>-</td></tr><tr><td></td><td>Y<sub>1-v</sub>(Ni, Mn)<sub>2</sub></td><td>F<span style="text-decoration: overline;">d</span>3m</td><td>Y<sub>0.82</sub>Ni<sub>1.69</sub>Mn<sub>0.31</sub></td><td>99 (1)</td><td>7.1912 (1)</td><td>-</td><td>-</td></tr><tr><td></td><td>Y<sub>2</sub>O<sub>3</sub></td><td>-</td><td>-</td><td>trace</td><td>-</td><td>-</td><td>-</td></tr></tbody></table>

## 2.2. Hydrogen Sorption Properties

### 2.2.1. Pressure-Composition Isotherms

After activation pretreatment at 150 °C under vacuum, the PCIs were measured at 150 °C for all YₓNi₂₋ᵧMnᵧ (0.825 ≤ x ≤ 0.95, 0 ≤ y ≤ 0.3) alloys. As shown in Figure 1a, for Y₀.₉Ni₂₋ᵧMnᵧ (y = 0.1, 0.2, 0.3) compounds, the hydrogen storage capacity of the Mn0.1 (y = 0.1) compound is about 3.3 H/f.u., slightly lower than 3.8 H/f.u. for Mn0.2 (y = 0.2) and Mn0.3 (y = 0.3). YNi absorbs hydrogen at room temperature (RT) and forms very stable YNiH₃ hydrides [21,32]. The small amount of YNi in the Mn0.2 and Mn0.3 alloys should be responsible for the gap of the first measurable point from zero capacity. All the absorption PCI curves present a first plateau between 0.25 and 1.2 H/f.u., with a relatively flat plateau around a pressure of 5 kPa, and a second sloping plateau between 2.0 and 2.5 H/f.u., with a significant reduction in the plateau pressure with an increasing Mn content (from 0.3 MPa for the Mn0.2 compound to 0.08 MPa for the Mn0.3 compound at hydrogen concentrations of around 2.5 H/f.u.).
<figure><img src="image_4.png" /><figcaption>**Figure 1.** P–C isotherms measured at 150 °C for Y<sub>0.9</sub>Ni<sub>2−y</sub>Mn<sub>y</sub> (y = 0.1, 0.2, 0.3) (a) and Y<sub>x</sub>Ni<sub>2−y</sub>Mn<sub>y</sub> samples with a single superstructure or C15 Laves phase structure (b); full symbols stand for absorption, and empty symbols stand for desorption; “a–” stands for hydrogen absorption, and “d–” stands for hydrogen desorption.</figcaption></figure>

A similar behavior has been observed for the single-phase compounds with a super-
structure (y = 0) and a single C15 structure (Mn-containing compounds) (Figure 1b). The
first plateaus overlap each other, showing an equilibrium pressure of 0.005 MPa. In addition,
a second plateau can be observed for the Y<sub>0.95</sub>Ni<sub>2</sub>, Y<sub>0.86</sub>Ni<sub>1.8</sub>Mn<sub>0.2</sub>, and Y<sub>0.825</sub>Ni<sub>1.7</sub>Mn<sub>0.3</sub>
compounds, and the Mn for Ni substitution slightly lowers the second plateau pressure.
Y<sub>0.9</sub>Ni<sub>1.9</sub>Mn<sub>0.1</sub> shows different behaviors compared with others as a sloped branch is ob-
served between 1.5 and 2.5 H/f.u. Due to the limit of low-pressure measurements, full
desorption at 150 °C cannot be achieved.

2.2.2. X-ray Diffraction Results after Hydrogenation

XRD measurements have been performed on Y<sub>0.9</sub>Ni<sub>2</sub>&minus;yMn<sub>y</sub> (y = 0.1, 0.2, 0.3) hydrides
hydrogenated at room temperature up to various hydrogen concentrations, with a final
pressure below 0.1 MPa. The XRD patterns of their hydrides are shown in Figure 2.
The crystallographic results obtained from Rietveld refinement for Mn0.1 compounds
are summarized in Table 2. For Mn0.2 and Mn0.3 compounds due to the broadness of
the diffraction peaks and significant background bumping, Rietveld analyses could not
be performed.
<figure><img src="image_5.png" /><figcaption>**Figure 2.** XRD patterns of the hydrides of Y<sub>0.9</sub>Ni<sub>1.9</sub>Mn<sub>0.1</sub> (a), Y<sub>0.9</sub>Ni<sub>1.8</sub>Mn<sub>0.2</sub> (b), and Y<sub>0.9</sub>Ni<sub>1.7</sub>Mn<sub>0.3</sub> (c) with different hydrogen concentrations prepared at room temperature.</figcaption></figure>

<table><caption>**Table 2.** Crystallographic data of the Y<sub>0.9</sub>Ni<sub>1.9</sub>Mn<sub>0.1</sub> alloy and its hydrides.</caption><thead><tr><th>Hydrogen Content(H/f.u.)</th><th>Phase</th><th>Abundance (wt.%)</th><th>Space Group</th><th>Lattice Parameters(Å)</th><th>Cell VolumeV (Å<sup>3</sup>)</th><th>ΔV/V(%)</th></tr></thead><tbody><tr><td rowspan="2">0</td><td>Y<sub>1-v</sub>(Ni, Mn)<sub>2</sub></td><td>96 (1)</td><td rowspan="2">Fd<span style="text-decoration: overline">3</span>m</td><td>7.1733 (1)</td><td>369.11 (1)</td><td rowspan="2">15.8</td></tr><tr><td>Y<sub>2</sub>O<sub>3</sub></td><td>4 (1)</td><td>/14.3466 (2)</td><td>/2952.8 (1)</td></tr><tr><td rowspan="2">1.1</td><td>Y<sub>1-v</sub>(Ni, Mn)<sub>2</sub>H<sub>w</sub></td><td>47 (1)</td><td>F<span style="text-decoration: overline">4</span>3m</td><td>14.318 (1)</td><td>2935.3 (5)</td><td rowspan="2">15.5</td></tr><tr><td>Y<sub>2</sub>O<sub>3</sub></td><td>52 (1)</td><td>F<span style="text-decoration: overline">4</span>3m</td><td>15.065 (1)</td><td>3419.3 (7)</td></tr><tr><td rowspan="2">1.6</td><td>Y<sub>1-v</sub>(Ni, Mn)<sub>2</sub>H<sub>w'</sub></td><td>trace</td><td>F<span style="text-decoration: overline">4</span>3m</td><td>14.320 (2)</td><td>2936.7 (7)</td><td rowspan="2">19.6</td></tr><tr><td>Y<sub>2</sub>O<sub>3</sub></td><td>35 (1)</td><td>F<span style="text-decoration: overline">4</span>3m</td><td>15.051 (2)</td><td>3409.7 (9)</td></tr><tr><td rowspan="2">2.0</td><td>Y<sub>1-v</sub>(Ni, Mn)<sub>2</sub>H<sub>w''</sub></td><td>21 (1)</td><td>F<span style="text-decoration: overline">4</span>3m</td><td>14.357 (1)</td><td>2959.4 (4)</td><td rowspan="2">15.5</td></tr><tr><td>Y<sub>2</sub>O<sub>3</sub></td><td>78 (1)</td><td>F<span style="text-decoration: overline">4</span>3m</td><td>15.229 (1)</td><td>3532.5 (5)</td></tr></tbody></table>

For Mn0.1, up to 2.0 H/f.u., the sample contains two phases crystallizing in the TmNi₂ structure, the intermetallic and its corresponding hydride with an expanded lattice parameter. For the two samples with hydrogen contents of 1.2 and 1.6 H/f.u., respectively, the lattice parameter of the hydride remains almost constant, and the fraction of the hydride phase increases. This means that these compounds are located in the two-phase region that corresponds to the first plateau (see PCI at 150 °C in Figure 1a). Between 1.6 and 2.0 H/f.u., the lattice parameter of the hydride phase increases from 15.051 (1) Å to 15.229 (1) Å, and the corresponding cell volume increases from 2952.8 (1) Å³ to 3532.5 (5) Å³ (with a
cell expansion of 19.6%, as shown in Table 2), indicating the onset of the solid-solution β branch of the hydride phase. Remarkably, in this sample, there is still a nonhydrogenated intermetallic phase, which may be due to the partial desorption during the sample handling.

For the compounds with a larger Mn content ($y = 0.2, 0.3$), the PCIs show a “second plateau” (Figure 1). Their XRD patterns show the formation of an intermediate hydride between the intermetallic compounds and the $AB_2H_w$ hydride with a superstructure. The intensity of the Bragg peaks of this intermediate phase became markedly important for $y = 0.3$ at hydrogen concentrations up to 1.9 H/f.u., indicating that the formation of such a phase occurs more easily for compounds with a larger Mn content and a moderate H concentration. It is also obvious that for Mn0.3, the decrease in peak intensity compared with the background and the broadness of the diffraction peaks with H content may indicate a loss of crystallinity for all phases after hydrogen absorption. For the Mn0.2 and Mn0.3 compounds, the contribution of the bump in the background increases with H content. This shape of the background suggests a partial amorphization of the hydride, as it has been already observed in $An_2$ compounds [33]. The second plateau can also be related to a transformation from crystalline to amorphous hydride.

### 2.2.3. Structural Evolution upon Dehydrogenation

XRD characterization was performed after PCI measurements and subsequent hydrogen desorption under dynamic vacuum at 150 °C (Figure 3) for the single-phase compounds. As shown in Figure 3, a large bump with a maximum of around $2\theta = 42^\circ$ is observed for all samples. After dehydrogenation, the Bragg peaks of the initial $Y_{0.9}Ni_{1.9}Mn_{0.1}$ intermetallic are still present, as well as that of $Y_{0.86}Ni_{1.8}Mn_{0.2}$ but with a smaller intensity. The XRD pattern of desorbed $Y_{0.9}Ni_{1.7}Mn_{0.3}$ hydride displays mainly a broad bump in the background and a few peaks belonging to $Y_2O_3$ that were already present in the initial alloy. This indicates the full HIA occurrence of this compound, as observed for desorbed $Y_{0.95}Ni_2$ hydride. The presence of broad crystalline diffraction peaks of the C15 phase in equilibrium with an amorphous contribution after PCI measurement in the $Y_{0.9}Ni_{1.9}Mn_{0.1}$ compound indicates better structural stability than in other Mn compounds with a partial HIA.

<figure><img src="image_6.png" /><figcaption>**Figure 3.** XRD patterns of the compounds with single-phase structures after PCI measurements and hydrogen desorption under primary vacuum at 150 °C.</figcaption></figure>

### 2.3. TG-DSC Studies of the Amorphous Hydrides

To understand the hydrogen desorption properties of amorphous hydrides, thermal analysis was performed for $Y_xNi_{2-y}Mn_y$ ($x = 0.95, 0.90, 0.86, 0.825, y = 0, 0.1, 0.2, 0.3$) compounds after hydrogen desorption at 150 °C under primary vacuum. The TG-DSC curves are shown in Figure 4.
<figure><img src="image_7.png" /><figcaption>**Figure 4.** DSC (a) and TG (b) curves of YxNi2-yMnγ compound (x = 0.95, 0.9, 0.86, 0.825, y = 0, 0.1, 0.2, 0.3) samples upon PCI tests and hydrogen desorption under dynamic vacuum. T1 and T2 are the beginning and end temperatures for hydrogen desorption of amorphous hydrides, T4 is the end temperature of crystallization, and T5 is the end temperature of the second endothermic peak which can correspond to the hydrogen desorption from YH2.</figcaption></figure>

All the DSC curves of single-phase compounds with either a superstructure or a C15 structure show a large broad endothermic peak between T1 and T2, followed by a sharp exothermic peak at T3 (Figure 4a). In the temperature region between T1 and T2, a weight loss is observed in the TG curves. It should correspond to the hydrogen desorption of amorphous hydrides (between T1 and T2) and phase crystallization (T3). The shaded area below the broken line should correspond, therefore, to enthalpy changes during the hydrogen desorption from the amorphous hydrides. This temperature does not vary with the Mn content. On the contrary, the hydrogen desorption content measured by TG increases with the Mn content (Mn0.1 to Mn0.3, Figure 4b). The maximum amount of hydrogen desorption can be observed for Mn0.3, indicating a larger initial amount of amorphous hydrides in the sample, in agreement with previous results (Figure 3). The desorbed H2 per mol of sample was calculated from the sample weight loss (TG curves). Then, by integrating the endothermic peaks (shadow area on the DSC curves), the enthalpy changes for one mole of released H2 were calculated. The enthalpy changes are listed in Table 3. The Mn0.1 and Mn0.2 compounds show the minimum and maximum enthalpies, respectively.
<table><caption>**Table 3.** The hydrogen content released from the amorphous hydride phase (where *a* stands for amorphous, and *w* stands for the remaining hydrogen in the amorphous hydrides) and the enthalpy changes per mol H₂.</caption><thead><tr><th>Samples</th><th>Hydrogen Desorption (mol H<sub>2</sub>/mol Sample)</th><th>ΔH (kJ/mol H<sub>2</sub>)</th></tr></thead><tbody><tr><td>a-Y<sub>0.95</sub>Ni<sub>2</sub>H<sub>w</sub></td><td>0.39</td><td>41.58</td></tr><tr><td>a-Y<sub>0.90</sub>Ni<sub>1.9</sub>Mn<sub>0.1</sub>H<sub>w</sub></td><td>0.34</td><td>35.16</td></tr><tr><td>a-Y<sub>0.86</sub>Ni<sub>1.8</sub>Mn<sub>0.2</sub>H<sub>w</sub></td><td>0.39</td><td>52.80</td></tr><tr><td>a-Y<sub>0.825</sub>Ni<sub>1.7</sub>Mn<sub>0.3</sub>H<sub>w</sub></td><td>0.59</td><td>39.48</td></tr></tbody></table>

Just after hydrogen desorption, the following exothermic peak (shown in Figure 4a) should correspond to the recrystallization of the amorphous phase, T4 being the temperature at which the crystallization ends. At around 650 °C, there is another endothermic peak for all samples with a shift to a higher temperature as the Mn content increases.

To interpret the DSC curves, XRD analysis was carried out for the samples after thermal analysis at temperatures T4 and T5, and then they were cooled down to room temperature. As shown in Figure 5a, after T4, all the XRD patterns contain crystalline AB₃-type Y(Ni, Mn)₃ and YH₂ phases, and small amounts of AB₂ and Y₂O₃ are also present. This may indicate that during the hydrogen desorption of the amorphous hydride, a disproportionation reaction occurs; as YH₂ is a very stable hydride, the freshly formed Y catches hydrogen and forms yttrium hydride. As a result of the balance, the crystalized phase is richer in the *B* element forming Y(Ni, Mn)₃ instead of Y(Ni, Mn)₂. After heating to T5, the absence of the YH₂ phase and the presence of the C15 or superstructure Y (Ni, Mn)₂ phase indicates a recombination reaction between the YH₂ and AB₃-type Y(Ni, Mn)₃ phase (Figure 5b) to form the initial Laves phase compound. This reaction is accompanied by a weight loss due to the hydrogen release from YH₂ (Figure 4b). The increase in T5 with the Mn content reflects the relative stability of the AB₂/AB₃ phase: the larger the Mn content, the more stable the AB₃ phase as observed at room temperature on the alloys. The presence of an Y₂O₃ oxide and the remaining AB₃ phase means an incomplete recombination reaction due to partial oxidation of Y during DSC measurement.

<figure><img src="image_8.png" /><figcaption>**Figure 5.** *Cont.*</figcaption></figure>
<figure><img src="image_9.png" /><figcaption>**Figure 5.** The XRD patterns of YₓNi₂₋ᵧMnᵧ (x = 0.95, 0.90, 0.86, 0.825, y = 0, 0.1, 0.2, 0.3) compounds upon thermal analysis at different temperatures: T4, the end temperature of crystallization (a); T5, the end temperature of the second endothermic peak (b).</figcaption></figure>

Similar results have been reported in CeNi₂: the recrystallization upon heating of the amorphous CeNi₂.₁₆H₃.₆₅ hydride yields first a mixture of CeNi₅ and CeHₓ, and then a progressive transformation between CeNi₅ and CeHₓ towards the CeNi₂ intermetallic occurs [34].

**3. Discussion**

**3.1. Mn Substitution Effects on Phase Occurrence and Structures**

The binary Y₀.₉₅Ni₂ compound crystallizes in an earlier reported TmNi₂-type superstructure with the space group F\bar{4}3m [29,31,35]. Upon Mn for Ni substitution in YₓNi₂₋ᵧMnᵧ (y ≥ 0.1), a structural transformation from a TmNi₂-type superstructure to a C15 Laves phase structure with disordered Y vacancies occurs. Such structural evolution has been reported in Y₀.₉₅Ni₂₋ᵧBᵧ (B = Cu, Fe, Al) systems [21,36]. Meanwhile, single-phase compounds with a C15 structure can be obtained by adjusting both the Y and Mn contents (Y₀.₉₀Ni₁.₇Mn₀.₃ → Y₀.₈₂₅Ni₁.₇Mn₀.₃, e.g.,). With the Mn content increasing from 0 to 0.3, Y vacancy formation was only observed in Y₀.₉₅Ni₂ and Y₀.₉₀Ni₁.₉Mn₀.₁ compounds, whereas Y₀.₈₆Ni₁.₈Mn₀.₂ and Y₀.₈₂₅Ni₁.₇Mn₀.₃ showed Mn occupations in both the Y and Ni sites without vacancy formation [30].

**3.2. Hydrogenation Properties**

Due to Y vacancy formation or bilateral distribution of Mn atoms, the PCIs of Y₀.₉₅Ni₂, Y₀.₈₆Ni₁.₈Mn₀.₂, and Y₀.₈₂₅Ni₁.₇Mn₀.₃ compounds show a similar behavior but are different from that of Y₀.₉₀Ni₁.₉Mn₀.₁ [30]. After a short plateau at low pressure, Mn0.1 shows a sloped branch, whereas the other three compounds show a second plateau (Figure 1b). The change of equilibrium pressures of the first plateau is small for all hydrides up to 1.2 H/f.u. (Figure 1). However, by comparing the XRD patterns of Y₀.₉₀Ni₁.₉Mn₀.₁H₂, Y₀.₈₆Ni₁.₈Mn₀.₂H₂.₁, and Y₀.₈₂₅Ni₁.₇Mn₀.₃H₁.₉, it is clear that for the Mn0.1 compound, no amorphous hydride is formed while a bump in the background for Mn0.2 and Mn0.3 indicates partial amorphization for an H content of around 2 H/f.u. (Figure 2).

Regarding the hydrogen desorption on the PCI, most of the hydrogen cannot be fully released under PCI measurement conditions (Figure 1). The XRD patterns (Figure 3) of samples dehydrogenated under vacuum show that all samples contain amorphous hydrides, indicating hydrogen-induced amorphization (HIA). As reported in other rare-earth-based
$AB_2$ compounds, the occurrence of HIA depends on the hydrogenation pressure, time, and temperature [37–40]. It was reported that $Y_{0.95}Ni_2$ forms a crystalline hydride at 100 °C under a pressure up to 3 MPa and becomes amorphous under 5 MPa of hydrogen [23]. We have measured the PCI at 150 °C with a maximum pressure of 10 MPa. After hydrogenation at RT up to 2 MPa, a crystalline hydride was still observed, but a disproportion occurred after PCI desorption. We can therefore assume that the amorphization occurs in the solid-solution branch between 2 and 5 MPa for Mn0. For Mn0.2 and Mn0.3 compounds, we interpret that the second plateau (Figure 1b) compounds the HIA, where the crystalline hydride transforms to an amorphous hydride with an increasing hydrogen content.

Another feature of Figure 3 is a partially recovered intermetallic phase for Mn0.1; we can conclude that Mn0.1 shows a better structural stability upon hydrogenation than Mn0, Mn0.2, and Mn0.3 compounds. As shown in a previous work [30], in the Mn0.1 compound, the Mn atoms are substituted only in the Ni site, whereas the vacancies are in the Y site. For compounds with a larger Mn content (Mn0.2 and Mn0.3), the Mn atoms occupy both the Y site and the Ni site, and the number of vacancies in the Y site is reduced [30]. Considering that for Mn0.1 ($Y_{0.90}Ni_{1.9}Mn_{0.1}$), Y vacancies ($r_{vacancy} = 0$) are in the A site, whereas the Ni and Mn atoms are in the B site, its average atomic radius ratio $r_A/r_B = 1.29$ is smaller than that for two other Mn compounds where Mn is located in both A and B sites (1.40 for Mn0.2 and 1.39 for Mn0.3). The larger $r_A/r_B$ ratio for Mn0.2 and Mn0.3 compared with the ideal value of 1.225 and the critical value of 1.37 observed for the onset of HIA can explain their structural stability compared with Mn0.1. This agrees with the observation that the atomic radius ratio is a dominating factor for the HIA in $AB_2$ Laves phase compounds [39,41]. To reach the ideal atomic radius ratio, the A site vacancies and the antisite occupations (Mn occupying the A site) should be considered.

However, changes in the electronic structure caused by Mn substitution cannot be fully excluded. It was shown that for $AB_2$ ($A$ = rare earth; $B$ = Mn, Fe, Co, and Ni), in the case of the same A element, the stability of $AB_2$ compounds depends on the B element; it increases from Mn, Fe, and Co to Ni [42–44], and the geometric effect plays a reverse role.

Here, since Mn has fewer 3d-electrons than Ni, Mn occupation in Ni sites will cause a reduction in electron number in the Ni d-shell. The literature [45–49] reported that Mn occupation in C15-type $RNi_2Mn$ ($R$ = Tb, Dy, Ho, and Er) compounds prominently impacts their electronic structures. The calculation of $RNi_2Mn$ indicated an increase in the density of the state at the Fermi level due to electronic transfer [49]. Indeed, the formation of R-Ni compounds relies on the charge transfer and filling of the 3d-electron band by valence electrons of rare-earth atoms [50]; Mn doping in R sites will cause a decrease in R atom number. Such changes in the electronic structure may have a direct impact on the structural stability during hydrogen sorption, which deserves further exploration.

Upon thermal analysis, the phase occurrence with increasing temperature can be summarized into three steps as follows (a-: amorphous, c-: crystalline):

$$a-Y_{1-v}(Ni, Mn)_2H_w \xrightarrow[\text{1st endothermic}]{\text{YH}_2+\text{H}_2} a-Y_{1-v}(Ni, Mn)_2H_{w'}+\text{H}_2 \xrightarrow[\text{endothermic}]{\text{c-Y}(Ni, Mn)_3+c-} c-Y_{1-v}(Ni, Mn)_2+\text{H}_2$$

Considering the better performance in hydrogen sorption of $Y_{0.9}Ni_{1.9}Mn_{0.1}$ than other high Mn-containing compounds, the geometric effect caused by Y vacancies and Mn exclusive occupation in Ni sites is more important. The formation of Y vacancies shifts the surrounding Y atoms closer to the vacancy center [19]; thus, a Y-Y distance is contracted near the vacancy, whereas the other distances are expanded. As a result, the site with a Y vacancy is much smaller than those occupied by Y atoms, and the average $r_{(Y+vacancy)}/r_{(Ni+Mn)}$ ratio becomes close to or smaller than 1.37. Consequently, in Mn0.1, HIA is partially suppressed by the synergistic effect of Y vacancy formation and Mn for Ni substitution. These results may be used to guide the research for more stable Laves phases for hydrogen storage.
**4. Materials and Methods**

The synthesis of YₓNi₂⁻ᵧMnᵧ (0.825 ≤ x ≤ 0.95, 0.1 ≤ y ≤ 0.3) compounds is described in [30]. Pressure–composition isotherm (PCI) curves were measured at 150 °C by an experimental apparatus based on the Sieverts method using 100 µm crushed powder. Before PCI measurements, the samples were activated at 150 °C under vacuum for 1 h. After PCI measurements, a 1 h dynamic vacuum at 150 °C was adopted for hydrogen desorption. Hydrides with different hydrogen concentrations were prepared at room temperature for XRD measurements. The sample holders were opened in the air to favor surface passivation to avoid hydrogen desorption. To observe the structural behaviors during hydrogen desorption and crystallization processes, the dehydrogenated samples were subsequently analyzed by DSC and thermogravimetric analysis (ATG Setsys Evolution 16 SETARAM) with a heating rate of 5 °C/min from 25 to 900 °C under argon flow. XRD measurements were performed after DSC at different temperatures.

XRD measurements of as-cast alloys were carried out with a D8 DAVINCI diffractometer from Bruker with Cu-Kα radiation in the 2θ range of 10°–90°. Diffraction patterns were collected at 40 kV and 40 mA for the anode of the X-ray tube. Rietveld analyses of the experimental data were performed with the program FullProf [51]. The phase compositions were examined by electron probe microanalysis (EPMA) using CAMECA (Gennevillier, France) SX100.

**5. Conclusions**

The structural and hydrogen sorption behaviors of intermetallic compounds with the nominal composition YₓNi₂⁻ᵧMnᵧ (x = 0.95, 0.90, 0.86, 0.825, y = 0, 0.1, 0.2, 0.3) were investigated systematically. The formation of the C15 Laves phase structure with disordered vacancies is favored against the superstructure one of Y₀.₉₅Ni₂. The single C15 phase can be formed in a composition region by adjusting both the Y and Mn components.

Y vacancy formation and Mn occupation in Ni sites (Mn0.1) enhance the structural stability and reduce the HIA effect. Mn occupation in Y sites shows the possibility to further decrease rA/rB, but a larger Mn content (y ≥ 0.2) yields less structural stability during hydrogenation, i.e., an easier tendency towards HIA.

**Author Contributions:** H.S.: data curation, formal analysis, investigation, writing—original draft, writing—review and editing; V.P.-B.: writing—original draft, writing—review and editing; P.L.: editing; L.J.: funding acquisition, supervision; J.Z.: conceptualization, formal analysis, writing—original draft, writing—review and editing, funding acquisition, supervision. All authors have read and agreed to the published version of the manuscript.

**Funding:** This work is supported by Campus France under Cai Yuanpei project 44027 WH.

**Data Availability Statement:** Data will be made available on request.

**Conflicts of Interest:** The authors declare no conflicts of interest. The funder was not involved in the study design, collection, analysis, interpretation of data, the writing of this article or the decision to submit it for publication.

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