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Article

Dispersion of Sintered Mg-Ni-Ce Materials for Efficient Hydrogen Storage

by
Nuriya Mukhamedova
1,2,*,
Arman Miniyazov
1,2,*,
Aisara Sabyrtayeva
1,*,
Timur Tulenbergenov
1,2 and
Ospan Oken
1,3
1
Institute of Atomic Energy Branch of NNC RK, Kurchatov 071100, Kazakhstan
2
Department of Technical Physics and Heat Power Engineering, Shakarim University, Semey 071410, Kazakhstan
3
Department of Information and Technology Sciences, Alikhan Bokeikhan University, Semey 070000, Kazakhstan
*
Authors to whom correspondence should be addressed.
Crystals 2025, 15(8), 743; https://doi.org/10.3390/cryst15080743
Submission received: 14 July 2025 / Revised: 15 August 2025 / Accepted: 18 August 2025 / Published: 20 August 2025
(This article belongs to the Section Inorganic Crystalline Materials)
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Abstract

This paper presents the results of the effect of dispersion on the structural-phase state of the material for hydrogen storage of the Mg-Ni-Ce system. X-ray phase analysis and scanning electron microscopy studies have shown that the sequential use of mechanical synthesis and spark-plasma sintering methods ensures the formation of a stable and dense microstructure with a high content of the intermetallic phase Mg2Ni. As a result of dispersion for 1 h, the sintered material was transferred to a finely dispersed state without changing the phase composition. Increasing the duration of dispersion to 2 h led to the formation of large agglomerates and the destruction of the material structure. For the first time, the dispersion technology was applied to materials of the Mg-Ni-Ce system, pre-sintered by spark-plasma sintering.

1. Introduction

Hydrogen energy plays a key role in the strategy for transition to a sustainable and low-carbon economy [1,2,3,4,5]. Currently, large-scale scientific research is being conducted around the world to improve hydrogen storage methods. The main problems of hydrogen storage are related to its low bulk density, the need for high pressure or cryogenic temperatures, leaks, brittleness of materials and the high cost of technologies. One of the promising solutions to these problems is the use of metal hydrides based on intermetallic compounds (IMCs), which provide safe and compact storage of hydrogen compared to traditional methods [6,7,8,9,10,11].
Metal hydrides are characterized by the ability to reversibly absorb and release hydrogen depending on the gas phase pressure and temperature, as well as to form stable hydride phases with high resistance to cyclic use [12,13]. Magnesium (Mg) as a material for hydrogen storage is of particular interest due to its low molecular weight, which provides a high gravimetric capacity (~7.6 wt.%). There are two main obstacles to the practical application of pure magnesium: low kinetics of hydrogen absorption/release and the need for high temperatures for desorption.
When metallic Mg interacts with H2, hydride MgH2 is formed, which in the initially hexagonal lattice of Mg builds a tetragonal crystal structure, occupying interstices in the lattice [14,15,16]. In the α-phase, the hydrogen concentration reaches 0.4 wt.%, and at higher concentrations, the α-phase is transformed into β-MgH2. This process is accompanied by an expansion (by about 30% [15,17]) of the hexagonal close-packed lattice of Mg. During the decomposition of the hydride, the reverse transformation into the α-phase begins at T~350 °C and atmospheric pressure [18], and this is due to the high enthalpy of hydrogen desorption (ΔH = 74.5 kJ/mol H2 [15]), which is due to the ionic–covalent nature of the bond. To reduce the enthalpy, Mg is alloyed with rare earth and transition elements, which form less stable hydrides (for example, Mg2Ni/Mg2NiH4 − ΔH = 62 kJ/mol H2 [15]) [17,18,19].
Therefore, alloying magnesium (Mg) with transition nickel (Ni) and rare earth cerium (Ce) promotes an increase in the kinetics of hydrogen sorption, thereby substantiating the promise of the IMC of the Mg-Ni-Ce system [20,21,22,23]. According to [24], alloys with a Ce content of 10 to 15 wt.% and Ni 5–10 wt.% demonstrate the formation of an intermetallic structure, including Mg2Ni, Mg12Ce phases, which contribute to improved hydride formation kinetics. However, in [25], at a higher Ce content (for example, 15%), a decrease in the fraction of the metallic Mg phase with a hexagonal crystal structure is observed, which may limit the overall hydrogen capacity. In the study [26], it was shown that the Ce5Mg85Ni10 composition demonstrates good hydrogen storage characteristics due to the finely dispersed structure and stabilization of hydride phases.
Materials based on IMC are characterized by structural ordering, characterized by high symmetry and specific distribution of atoms of various elements over sublattices, which determines their physicochemical properties [27]. To form such a structure, methods such as mechanical synthesis (MS) and spark-plasma sintering (SPS) are used. However, with MS, it is difficult to control the formation of the required structure due to magnesium’s tendency to oxidize [28]. Among the existing methods of powder metallurgy, the SPS method attracts attention due to its high heating rate and the use of pulsed direct current, which causes plasma discharges and suppresses grain growth under applied pressure. The use of the SPS method to Mg-based systems reduces the risk of its oxidation due to the vacuum environment and applied pressure, which reduces the temperature dependence of the sintering process. An additional advantage of the method is the localized point effect of pulsed electric current directly on the powder particles, which leads to intense heating of the contact zones without the need to bring the material to the melting point [29,30,31,32].
The sequential application of mechanical alloying (MA) and spark-plasma sintering (SPS) enables the synthesis of materials with enhanced structure and controlled phase composition, including intermetallic compounds. The combined use of MA and SPS was previously examined in detail in our earlier work [33], which presented the results of synthesis, microstructural analysis, and phase composition of the obtained samples. Although SPS leads to the formation of a consolidated structure, it also results in a reduced specific surface area due to increased material density. The specific surface area is one of the key characteristics of materials intended for hydrogen storage, directly influencing the sorption and desorption kinetics [34]. To increase the specific interaction surface and ensure efficient hydrogen uptake, dispersion becomes a mandatory technological step.
When dispersing sintered materials, mechanical separation of the material occurs, usually along the phase boundaries formed during sintering. Due to the different hardness and brittleness of the intermetallic components, the material is easily separated along the phase boundaries. As a result, ultrafine powders with particle sizes in the range from 5 μm to 100 nm can be obtained. This approach allows preserving the phase composition and local structure while significantly increasing the specific surface area of the material [35,36,37].
In modern studies of finely dispersed materials, special attention is paid to a stable microstructure [38]. However, during the dispersion process, mechanical action can activate the material due to the formation of fresh surfaces, which can lead to particle agglomeration and undesirable phase transformations [39,40]. In order to avoid such structural changes, it is especially important to carry out dispersion under controlled conditions, in which the mechanical action is limited and aimed exclusively at reducing the particle size without damaging their structure, which is the novelty of the work. In connection with the observations above, the aim of the work is to study the effect of the dispersion process on the structural and phase characteristics of the Mg-Ni-Ce system obtained by sequential exposure to MS and SPS.

2. Materials and Methods

2.1. Composition and Ratio of Alloying Components

The Mg-10%Ni-5%Ce ratio was chosen because it is the most balanced and can achieve optimal material characteristics for conducting this study. High-purity metal powders were used as starting materials for the preparation of the investigated samples, ensuring minimal impurity content and high reproducibility of the results. Magnesium was used in the form of MPF-3-grade powder with a purity of no less than 99% and a particle size range of 120–180 µm. Cerium was employed as a powder with a purity of 99.99% and an average particle size of approximately 70 µm, which enabled uniform distribution within the mixture and effective interaction with other components during thermal treatment. Nickel was introduced as PNK-UT3-grade powder with ultra-high purity (99.99%) and a particle size of about 20 µm, contributing to increased reactivity and improved microstructural homogeneity of the resulting material. The selection of these powders was driven by their high chemical purity, meeting the requirements for the synthesis of intermetallic compounds and alloys with targeted physicochemical properties.

2.2. Synthesis, Sintering and Dispersion of Materials of the Mg-Ni-Ce System

The initial powders were subjected to MS in a planetary ball mill for 10 h (the mill rotation direction was rotated once per hour during the synthesis) at 450 rpm. During the experiments, a 250 mL grinding jar and 5 mm diameter grinding balls made of AISI 420 stainless steel (material hardness 48–50 HRC) were used. The ratio of the mass of grinding balls to the mass of powder (BPR) was 30:1. It should be noted that in order to prevent oxidation, the preparation and weighing of the powders, as well as loading the powder and ball mixture into the grinding jar, were carried out in a glove vacuum box within an argon atmosphere. Then, the obtained metal powders were sintered by the SPS method on the “ISKRA” installation at a temperature of 500 °C, holding time 5 min, and applied pressure 1 ± 0.01 MPa. Sintering modes were selected based on the conducted analytical review and previously conducted works [41], demonstrating optimal conditions for the formation of a dense structure and minimization of porosity [42]. The sintered material is a cylinder with a diameter of 25 ± 1 mm and a height of 5 ± 1 mm. Figure 1 shows the chamber of the SPS installation with the placed batch (graphite matrix) and the resulting sintered material.
The dispersion of the obtained sintered materials was carried out in a ball crusher using an 80 mL beaker and 5 mm diameter grinding balls made of tungsten carbide (material hardness 89.7 HRA), which contributed to the effective dispersion of the sintered materials. Dispersion was carried out in an argon atmosphere, at a BPR of 30:1, a rotation speed of 350 rpm, and a duration of 1 and 2 h. In order to minimize particle agglomeration and eliminate undesirable phase transformations, the grinding process was interrupted every 30 min. Such parameters were chosen in order to preserve the crystalline structure. As shown in [43,44], it is the processing conditions with a reduced load that make it possible to obtain finely dispersed powders with a preserved phase composition.

2.3. Research Methods

The microstructure of the material obtained under different processing conditions was examined using scanning electron microscopy (SEM) with a TM4000plus microscope equipped with an energy-dispersive spectroscopy (EDS) attachment. Prior to analysis, the material was ground and polished using a MetPrep 3™ grinding and polishing machine. The shooting was carried out in the secondary electron (SE) and backscattered electron (BSE) modes, providing high resolution of the surface relief and contrast by atomic number.
To assess the phase composition, the X-ray diffraction method was used using a D6 Phaser diffractometer in monochromatic Cu Kα radiation (λ = 0.15418 nm) in the angle range of 2θ = 20–90 degrees with a step of Δ2θ = 0.013 degrees and an exposure of 23.4 s. During the shooting, the voltage and current were 40 kV and 30 mA, respectively, and the table rotation was 5 rpm. The diffraction patterns were decoded using a specialized program for processing and searching, “DIFFRAC.EVA” version 7, and the PDF-4 Axiom 2025 [45] database was used to identify reflections. The search for suitable phases was carried out based on the evaluation of angular positions, phase intensities at different angles, and comparison of interplanar distance data. The height and width of the peaks at half maximum (FWHM) were obtained using automated profile analysis (integrated functionality in the DIFFRAC.EVA software) which uses the Voigt function [46]. For a more detailed analysis, the lattice parameters, microstrain (ε), were calculated from the experimental data using the Williamson–Hall method, and the average crystallite size (D) was calculated using the Scherrer equation. The specific surface area of interaction was estimated by the BET (Brunnoyer–Emmett–Teller) method on a V-Sorb analyzer based on the analysis of nitrogen gas adsorption isotherms at 77 K. Before measurements, glass flasks were washed with distilled water and ethanol. Powder and sintered materials were weighed on a high-precision analytical balance, after which they were degassed at 200 °C for 2 h in a vacuum (residual pressure < 1 Pa). Measurements were carried out using nitrogen (as an adsorbate), helium (inert gas), and liquid nitrogen for cooling. Isotherms were plotted in the range of relative pressures Pa/P0 up to 0.35. To improve accuracy, a series of five successive measurements were carried out for each material. The specific surface area (Ssp) was calculated using the linear form of the BET equation using the standard value of the nitrogen molecule area (wo = 16.2 Å2). The relative error of the method did not exceed 10%
Kinetic measurements of hydrogen absorption were carried out using the H-Sorb high-pressure sorption analyzer. Before the measurements, the powder and sintered materials were pre-degassed for 2 h at 350 °C under a vacuum to remove surface-adsorbed gases and moisture. Helium was used as a calibration gas to correct the void volume. High-purity hydrogen gas was used as the sorbed gas. Kinetics measurements were carried out at 300 °C and 50 bar, which ensured sufficient thermodynamic activity of hydrogen. The time dependence of the absorbed hydrogen concentration was recorded in real time until the system reached a steady-state level. The time to reach the threshold hydrogen concentration (1 wt.%) was used as a comparison criterion.
To conduct a comparative analysis, the microstructure, phase composition, and specific surface area were studied for both the solid sintered material obtained by spark plasma sintering and the fine powders of this system after dispersion. To provide a complete picture of the efficiency of the applied technology, the kinetics of hydrogen sorption was studied for activated pure magnesium, sintered material of the Mg-Ni-Ce system, and powder material of the same system after 1 h of dispersion.

3. Results and Discussion

3.1. Microstructural Analysis

The microstructure at ×2000 magnification of the Mg-Ni-Ce system sintered by the SPS method is shown in Figure 2. The morphology of the sintered material is characterized by a dense structure with a high degree of compaction and the absence of pores. This reflects the features of structure formation in sintered materials, due to the use of optimal sintering modes while taking into account the preliminary MS.
The BSE image (Figure 2b) shows a pronounced inhomogeneity of the contrast, caused by the heterogeneity (multiphase structure) of the material: light areas (indicated as 1) correspond to intermetallic compounds, while dark areas (indicated as 2) correspond to the magnesium matrix. Oxide inclusions (indicated as area 3) were observed in the microstructure of the investigated material, as confirmed by the combined results of energy-dispersive X-ray spectroscopy and elemental mapping (Figure 3). A joint analysis of the local composition and spatial distribution of elements revealed a significant oxygen content in the specified regions.
Similar morphological features are observed in the work [47], devoted to the study of sintered materials of the Mg-Ni-Ce system obtained by the SPS method after preliminary MS. The authors recorded the formation of a uniform structure with pronounced phase inhomogeneity, observed by the BSE contrast. In the indicated work, it is shown that the phase components of the system (in particular, Mg2Ni and Ce2Mg17) are distributed throughout the volume of the sintered material in the form of light enriched zones, which are confirmed by elemental analysis.
Figure 4 shows SEM images of the powder material obtained after 1 h of dispersion. According to the results, the bulk of the particles have an irregular, angular shape, characteristic of mechanically dispersed materials.
The particles demonstrate pronounced porosity and surface roughness, which potentially increases their specific surface area. In [48], a similar particle morphology is reported in dispersed composites of the Mg-Ni-Ce system: particles with a torn contour, multiple cracks, and a developed surface are observed in SEM images. The authors attribute the formation of such structures with mechanical action and note that this contributes to the improvement of hydride kinetics due to an increase in the contact area and easier penetration of hydrogen to the active areas of the material.
The uneven distribution of contrast in the SEM image in the BSE mode, due to the heterogeneity of the structure, indicates the preservation of the structure observed in the sintered materials after dispersion. Figure 5 shows mapping of the powder material of the Mg-Ni-Ce system after dispersion for 1 h.
Thus, the morphology we obtained after dispersion is confirmed by the results of SEM analysis from literary sources [48,49], where similar structural characteristics are considered as positive factors contributing to the improvement of the hydrogen properties of magnesium systems.
Figure 6 shows the SEM images of the powder material after dispersion for 2 h. At ×1000 magnification, it can be seen that the particles have increased in size and consist of smaller fragments that have stuck together. According to [43,44], such behavior of the particles is associated with intense mechanical action. Long-term grinding leads to the accumulation of plastic defects and structural changes, which, in turn, change the morphology of the particles and contribute to their agglomeration.
Figure 7 shows mapping of the powder material of the Mg-Ni-Ce system after dispersion for 2 h.
The formation of larger loose particles is also explained by the increased content of oxides, which act as a binder, which promotes the adhesion of particles to each other and causes the effect of cold welding [38,50,51]. The results of elemental mapping for oxygen reveal the presence of oxide inclusions in both powder materials. However, in the material after 2 h of dispersion, a high density of oxygen-containing areas is observed, which indicates more pronounced surface oxidation.

3.2. X-Ray Phase Analysis

The diffraction pattern of the sintered material of the Mg-Ni-Ce system is shown in Figure 8. Table 1 shows the lattice parameters, microdefect values, and average crystallite sizes for the matrix phase (Mg) and intermetallic phase (Mg2Ni) calculated from the obtained data.
The obtained diffraction pattern of the sintered material of the Mg-Ni-Ce system demonstrates narrow and intense peaks (up to 6500 counts), which indicates a high degree of crystallinity of the material. The basis of the phase composition of this material is the hexagonal matrix phase of Mg, with the space group P63/mmc. In the lattice parameters calculated from the experimental data, a = b = 3.2159 Å, c = 5.2208 Å, there is an insignificant increase compared to the reference values (a = b = 3.205 Å, c = 5.215 Å; ICCD PDF-4, 04-015-2580). This difference (˂0.3%) is due to the presence of residual stress after sintering.
According to the X-ray phase analysis of space group P6222 (180), the sintered material has a large amount of the intermetallic phase Mg2Ni with a hexagonal crystal structure. As noted by the authors of [25,33,49,52], this intermetallic phase is one of the hydride-forming phases, capable of reducing the desorption temperature due to the lower decomposition temperature of hydride phases compared to pure magnesium. The average crystallite size of this phase is 61 nm, which indicates a nanocrystalline structure of the material and a developed grain boundary of this phase. Broadening and shifting of diffraction peaks to the region of large angles indicates the presence of defects and dislocations of the crystal lattice, which is also confirmed by the value of microdefects for this phase (0.00262). The presence of such defects plays an important role in hydrogen absorption processes: dislocations can serve as traps for hydrogen atoms, promote accelerated diffusion, and also form local areas with increased free volume.
The phase composition of this material also includes MgO and Ce2Mg17 phases. Despite the insignificant content of the oxide phase, its presence affects the structural state of the material. Magnesium is characterized by high chemical activity and easily interacts with components of the environment, which contributes to the formation of an oxide shell. According to [53], the MgO phase that appears on the surface during synthesis stabilizes the morphology of nanoparticles and prevents their agglomeration and further oxidation in the volume. The authors of this work note that a thin MgO layer acts as a protective barrier, limiting grain growth and maintaining the structural integrity of the active phases during sorption and desorption cycles. Moreover, it has been shown that MgO can have a positive effect on hydrogen capacity and hydride formation kinetics, preventing structural destruction and ensuring stability during multiple hydrogenation–dehydrogenation cycles [53].
Figure 9 shows the results of X-ray phase analysis of the powder material of the Mg-Ni-Ce system after dispersion for 1 h (Figure 5), which is similar in phase composition to the original material after spark-plasma sintering.
The diffraction pattern shows an increase in intensity and broadening of reflections. These differences are due to the transition of the sintered solid material to a powder state: as a result of dispersion, the average size of the crystallites of the Mg and Mg2Ni phases decreased (37 nm and 33 nm, respectively), which led to a broadening of the reflections. The increase in the intensity of the reflections (100) and (002) of the magnesium phase is explained by the reorientation of the crystallites [54,55]. The lattice parameters for these phases (see Table 1) are similar to those of the sintered material, and the microdefect values increase, indicating an increase in lattice dislocation.
After dispersion for 2 h (see Figure 10), the structure of the powder material of the Mg-Ni-Ce system underwent significant changes.
The diffraction pattern shows a significant broadening of the reflections, an increase in the intensity of oxide inclusions, and a complete destruction of the intermetallic phase of Mg2Ni. These changes in the diffraction pattern are a consequence of the intense mechanical action on the sintered material accompanying the dispersion process. The effect of high-energy deformation loads leads to the accumulation of structural defects, the grinding of crystallites and the destruction of the original intermetallic compounds [35,56]. The values of microdeformation of the Mg matrix phase (ε = 0.02365) calculated from the experimental data (Table 1) after dispersion for 2 h indicate a high degree of distortion of the crystal lattice, which exceeds the characteristic values for equilibrium structures [36].
An increase in the content of oxide inclusions was recorded, among which magnesium oxide MgO predominates. Particular attention should be paid to the intense reflexes of the phase identified as MgO4. This oxide phase is considered a non-stoichiometric oxide, which indicates the formation of oxides not only on the surface, but also in the internal volume of the crystal lattice [57]. The appearance of this oxide phase is due to the destruction of the protective layer in the form of the MgO phase during dispersion, resulting in the appearance of a non-stoichiometric oxide inclusion [58]. This is confirmed by the results of microstructure analysis on SEM and leads to deterioration of the structural integrity.

3.3. Specific Surface Area Analysis

Table 2 shows the specific surface area values of sintered and powdered materials of the Mg-Ni-Ce system, determined by the BET method using a low-pressure analyzer.
According to the obtained data, for the sintered material and the powder material of the Mg-Ni-Ce system, an increase in the specific surface area is observed after dispersion for one hour, which indicates effective grinding and an increase in the available active area. As shown in studies [59,60], an increase in the specific surface area of interaction of Mg-based composites is accompanied by an improvement in desorption properties and a decrease in the operating temperature. In addition, in the experimental work [60], it was demonstrated that even at moderate values of the specific surface area (up to 2.5 m2/g), a noticeable increase in the reactivity of Mg powders is observed.
After 2 h of dispersion, the powder material exhibited a specific surface area value falling below the detection threshold, which, in accordance with the literature, is regarded as effectively zero. This indicates the absence of gas sorption and can be associated with a large number of oxide inclusions in the material. Such inclusions block the active surfaces, preventing adsorption, and indicate the degradation of the sorption properties of the material.
Thus, the increase in the specific surface area of the finely dispersed powder material of the Mg-Ni-Ce system by more than five times compared to the initial sintered state confirms not only the efficiency of mechanical dispersion, but also indicates a potential increase in the reactivity of the material. These results are in good agreement with previously published data by other authors and emphasize the importance of textural and morphological characteristics in the development of magnesium materials for hydrogen technologies.

3.4. Hydrogen Sorption Kinetics

Figure 11 shows the kinetic curves of hydrogen sorption up to 1 wt.% for pure magnesium, sintered material of the Mg-Ni-Ce system and after dispersion for 1 h. For the material after dispersion for 2 h, the presence of non-stoichiometric oxides led to a complete loss of the sorption activity of the material, which excluded the possibility of obtaining data on the kinetics of hydrogen absorption.
From the analysis of the experimental curves, it is evident that alloying with Ni and Ce leads to a significant acceleration of the hydrogen sorption process. Pure magnesium demonstrates the lowest rate of hydrogen absorption—the value of 1 wt.% is reached only by the 40th minute. The introduction of Ni and Ce (blue curve) promotes the formation of active phases Mg2Ni and Ce2Mg17, which accelerates the sorption kinetics: the threshold concentration of 1 wt.% is reached already by the 31st minute. These compounds, as established in a number of works [22,24,25,52,61], are effective catalysts for the dissociation of hydrogen molecules and nucleation centers for MgH2.
The greatest improvement in sorption characteristics is observed in the powder material after mechanical dispersion (red curve). In this case, the concentration of 1 wt.% is reached by the 23rd minute, which is associated with the formation of a nanocrystalline structure, an increase in the surface area, and an increase in the density of defects and grain boundaries. It should be noted that the results obtained with high-energy grinding and the introduction of catalysts make it possible to achieve a significant acceleration of kinetics, especially in the initial phase of sorption, which is also noted in [25,41,48].
Thus, it has been experimentally confirmed that complex modification of the magnesium matrix due to alloying and an increase in the specific surface area of interaction makes it possible to significantly increase the rate of hydrogen absorption, which makes such materials promising candidates for stationary and mobile hydrogen energy systems.

4. Conclusions

In this study, the structural and phase state of the Mg-Ni-Ce system was investigated through a sequential application of spark-plasma sintering (SPS) and mechanical dispersion (MD). This approach resulted in the formation of a material predominantly composed of intermetallic phases, particularly Mg2Ni.
The morphology of the sintered material is characterized by a dense structure with no visible porosity. According to X-ray diffraction (XRD) analysis, microdeformations associated with dislocations were observed in the matrix after sintering. These structural defects enhanced the hydrogen absorption rate compared to pure magnesium, confirming the effectiveness of the applied technological strategy for improving the hydrogen-related properties of the material.
Mechanical dispersion for 1 h at a rotation speed of 350 rpm resulted in the formation of irregularly shaped particles with rough surfaces, typical of materials subjected to mechanical milling. The resulting powder preserved its initial phase composition (Mg, Mg2Ni, MgO, Ce2Mg17) and microstructure similar to the sintered sample, indicating that the material was fragmented primarily along phase boundaries. Furthermore, the specific surface area increased by a factor of five, which contributed to a twofold improvement in hydrogen sorption kinetics compared to pure magnesium.
Extending the dispersion time to 2 h led to the degradation of the ordered structure formed during SPS and altered the phase composition. Morphological analysis revealed particle agglomeration, while XRD confirmed the formation of non-stoichiometric oxides embedded within the crystal lattice, which hindered hydrogen access to the surface. This was further supported by a drastic reduction in the specific surface area to nearly zero. Hence, excessive mechanical processing resulted in structural degradation and a decrease in sorption activity.
Based on the conducted study, the optimal mechanical dispersion parameters for the Mg-Ni-Ce system were established: 350 rpm for 1 h. Under these conditions, the favorable crystal structure formed during SPS is preserved, and a significant increase in the active surface area is achieved, leading to enhanced hydrogen sorption kinetics.

Author Contributions

Methodology, N.M., A.S. and O.O.; Validation, A.M.; Writing—original draft, A.S.; Writing—review & editing, N.M., A.M. and T.T.; Visualization, T.T.; Project administration, N.M. All authors have read and agreed to the published version of the manuscript.

Funding

This work was carried out with the financial support of the Science Committee of the Ministry of Science and Higher Education of the Republic of Kazakhstan within the framework of grant funding of the Ministry of Higher Education of the Republic of Kazakhstan for project AP19574566 on the topic “Development of hydrogen storage materials based on Mg-Ni-Ce”.

Data Availability Statement

The original contributions presented in this study are included in the article. Further inquiries can be directed to the corresponding authors.

Conflicts of Interest

The authors declare no conflict of interest.

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Crystals 15 00743 g001
Figure 1. Working chamber of the spark-plasma sintering installation (a) and the external appearance of the sintered material (b).
Figure 1. Working chamber of the spark-plasma sintering installation (a) and the external appearance of the sintered material (b).
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Figure 2. SEM images of the cross-section of the Mg-Ni-Ce system material after sintering: 1–3—areas of different contrast.
Figure 2. SEM images of the cross-section of the Mg-Ni-Ce system material after sintering: 1–3—areas of different contrast.
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Figure 3. Mapping of sintered material of the Mg-Ni-Ce system at ×7000 magnification.
Figure 3. Mapping of sintered material of the Mg-Ni-Ce system at ×7000 magnification.
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Figure 4. SEM images of the powder material of the Mg-Ni-Ce system after 1 h of dispersion; 1–3—areas of different contrast.
Figure 4. SEM images of the powder material of the Mg-Ni-Ce system after 1 h of dispersion; 1–3—areas of different contrast.
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Figure 5. Mapping of the powder material of the Mg-Ni-Ce system after 1 h of dispersion.
Figure 5. Mapping of the powder material of the Mg-Ni-Ce system after 1 h of dispersion.
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Figure 6. SEM images of the powder material of the Mg-Ni-Ce system after 2-h dispersion.
Figure 6. SEM images of the powder material of the Mg-Ni-Ce system after 2-h dispersion.
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Figure 7. Mapping of the powder material of the Mg-Ni-Ce system after 2 h of dispersion.
Figure 7. Mapping of the powder material of the Mg-Ni-Ce system after 2 h of dispersion.
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Figure 8. Diffraction pattern of sintered material of the Mg-Ni-Ce system.
Figure 8. Diffraction pattern of sintered material of the Mg-Ni-Ce system.
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Figure 9. Diffraction pattern of finely dispersed powder material after dispersion for 1 h.
Figure 9. Diffraction pattern of finely dispersed powder material after dispersion for 1 h.
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Figure 10. Diffraction pattern of finely dispersed powder material after dispersion for 2 h.
Figure 10. Diffraction pattern of finely dispersed powder material after dispersion for 2 h.
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Figure 11. Kinetics of hydrogen sorption up to 1 wt.% in materials: activated powder of pure Mg, sintered Mg-Ni-Ce system and powder material of the Mg-Ni-Ce system after dispersion for 1 h.
Figure 11. Kinetics of hydrogen sorption up to 1 wt.% in materials: activated powder of pure Mg, sintered Mg-Ni-Ce system and powder material of the Mg-Ni-Ce system after dispersion for 1 h.
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Table 1. Crystal lattice parameters for the Mg and Mg2Ni phases.
Table 1. Crystal lattice parameters for the Mg and Mg2Ni phases.
MaterialPhaseLattice Parameter (a, Å)Lattice Parameter (c, Å)Microdeformation (ε)Crystallite Size (D, HM)
Sintered by SPS methodMg3.2159 ± 0.00025.2208 ± 0.000150.00073 ± 0.0002148.28 ± 0.015
Mg2Ni5.2284 ± 0.0001513.3397 ± 0.000170.00262 ± 0.000261.34 ± 0.018
After dispersion
1 hMg3.2152 ± 0.000145.2234 ± 0.000150.00144 ± 0.0001737.77 ± 03013
Mg2Ni5.2368 ± 0.0002113.3311 ± 0.000160.00298 ± 0.0001833.67 ± 0.001
2 hMg3.3571 ± 0.000235.4312 ± 0.000150.02365 ± 0.0002332.24 ± 0.002
Table 2. Specific surface area measured by the BET method.
Table 2. Specific surface area measured by the BET method.
MaterialSpecific Surface Area, m2/g
Sintered on SPS1.22 ± 0.007
After dispersion
1 h6.34 ± 0.01
2 h0 (−0.115)
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Mukhamedova, N.; Miniyazov, A.; Sabyrtayeva, A.; Tulenbergenov, T.; Oken, O. Dispersion of Sintered Mg-Ni-Ce Materials for Efficient Hydrogen Storage. Crystals 2025, 15, 743. https://doi.org/10.3390/cryst15080743

AMA Style

Mukhamedova N, Miniyazov A, Sabyrtayeva A, Tulenbergenov T, Oken O. Dispersion of Sintered Mg-Ni-Ce Materials for Efficient Hydrogen Storage. Crystals. 2025; 15(8):743. https://doi.org/10.3390/cryst15080743

Chicago/Turabian Style

Mukhamedova, Nuriya, Arman Miniyazov, Aisara Sabyrtayeva, Timur Tulenbergenov, and Ospan Oken. 2025. "Dispersion of Sintered Mg-Ni-Ce Materials for Efficient Hydrogen Storage" Crystals 15, no. 8: 743. https://doi.org/10.3390/cryst15080743

APA Style

Mukhamedova, N., Miniyazov, A., Sabyrtayeva, A., Tulenbergenov, T., & Oken, O. (2025). Dispersion of Sintered Mg-Ni-Ce Materials for Efficient Hydrogen Storage. Crystals, 15(8), 743. https://doi.org/10.3390/cryst15080743

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