# Crystallography-dependent MgH2 desorption mechanisms revealed by in-situ SEM with Ti3AlC2 MAX phase catalysts

Adrien Heinzelmeier, Fermin Cuevas, Thierry Grosdidier, Antoine Guitton

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Adrien Heinzelmeier, Fermin Cuevas, Thierry Grosdidier, Antoine Guitton. Crystallography-dependent MgH2 desorption mechanisms revealed by in-situ SEM with Ti3AlC2 MAX phase catalysts. 2026. hal-05512033

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# Crystallography-dependent MgH₂ desorption mechanisms revealed by *in-situ* SEM with Ti₃AlC₂ MAX phase catalysts

Adrien Heinzelmeierª, Fermin Cuevasᵇ,ᵇ*, Thierry Grosdidierª, Antoine Guittonª,\*

ªUniversité de Lorraine, CNRS, Arts et Métiers Institute of Technology, LEM3, F-57000 Metz, France,

ᵇUniv Paris-Est Creteil CNRS ICMPE (UMR 7182) 2 Rue Henri Dunant F-94320 Thiais France

## Abstract

In this work, *in-situ* scanning electron microscopy combined with crystallographic analysis is used to directly visualise, for the first time, the destabilization of MgH₂, under a localised electron beam in a Mg-based composite containing a Ti₃AlC₂ MAX phase. H desorption is shown to initiate and propagate along nanometric linear defects aligned with the {110} planes of the rutile MgH₂ lattice, the planes with the highest Mg atomic density. This orientation-selective transformation to Mg, accompanied by a volume shrinkage, releases residual stress fields by producing cracks within Mg and at MgH₂/Mg interfaces. The presence of Ti₃AlC₂ particles further localises or promotes the transformation at the interface between Ti₃AlC₂ and Mg. Such interphases act as favorable sites, analogous to grain boundaries, where the MgH₂→Mg + H₂ transformation preferentially occurs. These findings provide direct experimental evidence of crystallography-dependent desorption channels and point toward microstructural features that may be exploited to accelerate H release kinetics in Mg-based hydrides.

**Keywords:** Hydrogen storage, Desorption, Mg-based composite, MAX phases, Microstructure, Crystallographic defects, Electron beam

## 1. Introduction

Magnesium (Mg)-based alloys are attractive solid-state media for hydrogen (H) storage, combining high capacity and safety through reversible hydride formation [1]. Mg can store up to 7.6% H by weight [2] via sequential steps — H₂ physisorption, dissociation, chemisorption, bulk diffusion, and hydride nucleation — which reverse during de-storage [3]. However, sluggish kinetics often result in absorption/desorption times of several hours and require temperatures above ~200°C. Catalysts and microstructural engineering are therefore critical for enabling fast, low-temperature cycling [4–13].

MAX phases, with general formula M$_{n+1}$AX$_n$ (n = 1–3), consist of a transition metal (M), an A-group element (A), and either carbon and/or nitrogen (X). They are hexagonal, nanolayered solids with a P6$_{3}$/mmc space group [14, 15], comprising n layers of M$_{6}$X octahedra alternating with a single A layer. These structures exhibit marked crystalline anisotropy, with c/a ratios above 3 [14]. Importantly, owing to their distinctive crystallographic structure, MAX phases combine metallic and ceramic properties, which give rise to a range of intriguing deformation, oxidation, diffusion, and thermal mechanisms, making them promising for numerous applications [15–33].

Many transition metals catalyze Mg hydriding. For example, Rizo-Acosta [34] reported that adding a few at.% Ti yields the highest reversible H storage capacity over twenty 15 min absorption cycles, outperforming Sc, V, Y, Zr, and Nb. The active TiH₂ phase promotes H₂ dissociation, enhances H diffusion by ~10⁵, and accelerates MgH₂ nucleation at the TiH₂/Mg interface.

The catalytic effect of MAX phases on H sorption in Mg has been demonstrated [11]. For example, the temperature at which half of the H is desorbed decreases from 310°C for pure MgH₂ to 250°C for Mg + 7 wt.% Ti₃AlC₂, surpassing the effect of Ti, Al, or C alone. Lakhnik et al. [13] further showed that increasing Ti₃AlC₂ content from 0 to 5 wt.% and 7 wt.% reduces desorption temperatures from 277°C to 246°C and 236°C, respectively, while increasing total H desorbed from 3.1 wt.% to 4.1 wt.% and 5.6 wt.%. Johnson-Mehl-Avrami-Kolmogorov (JMAK) analysis [35] indicated a kinetic transition from a contracting-volume model (pure Mg) to a nucleation-and-growth model (composites), with MAX particles acting as nucleation sites for

*Corresponding author: Antoine Guitton, \* Email: antoine.guitton@univ-lorraine.fr

**Current address: Representative Office of CNRS in China, 60 Tianze lu, Liangmaqiao, 100600 Beijing China
Mg.

Among the 342 MAX phases reported to date [36], only a few have been studied as catalysts for H sorption in Mg, including Ti₃AlC₂ [11, 13, 37], Ti₃AlCN [38], TiVAlC [12], and Ti₂VAlC₂ [39]. Interstitial H incorporation has also been investigated in Ti₃AlC₂ and Ti₃SiC₂ [40–45].

While the mechanism of H sorption has received some attention, the desorption of H from MAX–MgH₂ mixtures has not yet been reported. In addition, the MgH₂ → Mg + H₂ transformation also remains poorly addressed.

The objective of this work is to report on the initi-
ation of MgH₂ desorption at grain boundaries (GBs) and
along linear cracks propagating on specific crystallographic
planes. These mechanisms were observed *in situ* during
electron backscatter diffraction (EBSD), where exposure
to the energy electron beam of the scanning electron mi-
croscope (SEM) locally destabilized the hydrides. This
study further examines the catalytic role of MAX phases
in the hydride desorption process.

## 2. Materials and methods

2.1. *Synthesis and characterization of the starting Mg-Ti<sub>3</sub>AlC<sub>2</sub> composite*

Ti₃AlC₂ powders were synthesized from titanium (Ti, 99.5% purity), aluminium (Al, 99.5% purity), and titanium carbide (TiC, 99.5% purity) powders, mixed in a molar ratio of 3Ti:1.1Al:1.9C, by pressureless sintering at 1500 °C. Mg-based composites were then fabricated by hot pressing pure Mg powders (Alfa Aesar, Ward Hill, MA) with 15 vol.% Ti₃AlC₂ phase particles at 700 °C under 30 MPa for 1 h.

Microstructural characterization was performed using
a Zeiss Supra 40 FEG-SEM (Oberkochen, Germany), oper-
ating at a residual pressure of approximately 1 × 10⁻⁴ Pa,
and equipped with energy-dispersive X-ray spectroscopy
(EDX) and electron backscatter diffraction (EBSD) sys-
tems, both from Bruker Nano GmbH (Berlin, Germany).
The EBSD acquisition conditions were as follows: accelera-
tion voltage 20 kV, aperture 120 µm, and working distance
15 mm. EBSD data were processed with the ATEX soft-
ware developed at LEM3 laboratory (Metz, France) [46].
For optimal imaging, surfaces were prepared by mechan-
ical polishing without water (due to the high reactivity
of Mg), followed by ionic polishing in a Gatan PECS II

(Pleasanton, USA) for 24 h at 6 keV with a 0° incidence
angle.

Figure 1 shows a backscattered electron (BSE) and secondary electron (SE) micrograph of the starting microstructure of the Mg – (Ti₃AlC₂)₁₅% composite. In BSE imaging (Z contrast), the Mg matrix appears in dark grey, while the Ti₃AlC₂ phase particles appear in bright contrast and delineate pure Mg domains. The average Mg domain size, measured by the mean linear intercept (MLI) method, is 85 µm. Due to uniaxial pressing during synthesis, these domains are anisotropic: the MLI reaches 127 µm in the transverse direction and 44 µm in the compression direction. This arrangement forms a characteristic necklace-type distribution, as previously reported for Mg-Ti₂AlC composites and other metal–ceramic systems [23, 47]. EDX maps of Ti and Al (not shown in Figure 1) confirm that most particles are close to the Ti₃AlC₂ stoichiometry. Only minor chemical heterogeneities are detected, with about 5.5% Ti-rich particles and 0.5% Al-rich particles, likely resulting from synthesis inhomogeneities.

2.2. Hydridation

Hydriding was carried out under controlled pres-
sure–temperature cycling in a Sieverts-type apparatus,
with H content calculated from real-time pressure mon-
itoring and corrected for non-ideal gas behavior using the
Hemmes model [48]. For the Mg – (Ti₃AlC₂)₁₅% com-
posite (maximum theoretical capacity 5.3 wt%), activation
was performed under 18 bar H₂ by fracturing the sample
into 1 mm pieces in an argon-filled glove box (Jacomex,
< 4 ppm O₂, < 0.2 ppm H₂O), heating to 400 °C, then
holding at 300 °C for 16 h. Three consecutive cycles were
carried out under these conditions, resulting in a final H
content of 2.9 wt%.

3. Results

3.1. Electron beam–induced damage in the hydride

Figures 2(a,b) show BSE micrographs of the same re-
gion, taken before and after the first EBSD acquisition,
whose EBSD phase maps are shown in Figure 2(c) and
(d). Red pixels in Figure 2(c) and (d) correspond to the
Mg phase, while blue pixels indicate the MgH₂ phase. Be-
tween the EBSD acquisition in Figure 2(c) and the micro-
graph in Figure 2(b), the sample was stored under vacuum
<figure><img src="image_2.png" /><figcaption>Figure 1: Characterization of the Mg – (Ti₃AlC₂)₁₅% specimen. (left) SEM micrographs of the microstructure; (right) Higher-magnification SEM view of the Ti₃AlC₂ necklace.</figcaption></figure>

<figure><img src="image_3.png" /></figure>

for several days, and subsequently cleaned from surface oxides by ionic polishing (1 h 30 min, 5 keV).

The brighter contrast of Mg compared to MgH₂ in BSE micrographs arises from its higher mean atomic number ($Z_{Mg} = 12$ vs. $Z_{MgH_2} \approx 4.7$), which increases the BSE yield.

In Figure 2(b), bright contrasts, highlighted by white arrows appear within MgH₂ domains. They are barely visible in Figure 2(a). The thickness of these bright edges reaches up to 2 µm. The details of these features are examined in greater detail in Figure 3. In addition, dark linear contrasts reveal the opening of cracks within the hydrides, also visible in Figure 2(b). Interestingly, these cracks tend to be parallel to each other within the same grain.

After the first EBSD map (Figure 2(c), a second EBSD acquisition was subsequently done in the same region, and the corresponding phase map is shown in Figure 2(d). All Mg grains indexed in Figure 2(c) are also readily indexed in Figure 2(d). Conversely, the MgH₂ phase (blue pixels) shows a sharp decrease in indexing rate, dropping from 36.9% in Figure 2(c) to only 7.8% in Figure 2(d), with most hydride regions remaining non-indexed. Notably, a portion of a hydride grain in Figure 2(d) is more indexed as Mg, as highlighted by the white dashed circle in Figure 2(d).

### 3.2. Crystallographic defects induced by desorption

Figure 3 focuses on two hydride grains from Figure 2 (red and green dotted squares corresponding to

Figures 3(c) and 3(b), respectively), to illustrate the microstructure evolution by comparing BSE micrographs acquired before and after the first EBSD acquisition. In the first grain, the hydrides in Figure 3(a) are undamaged. After the first EBSD acquisition (Figure 3(b)), linear defects (white arrows) appear, extending across the grain from one GB to the other. These defects follow two preferential orientations: 62° ±5° and 173° ±5° from the horizontal, giving the appearance of parallel arrays.

In the second grain (Figure 3(c)), similar linear features are already present prior to EBSD acquisition, oriented at 131° ±2° and 47° ±2° from the horizontal. Together, these two sets form nearly square patterns within the grain, hereafter referred to as "windows", following the nomenclature introduced by Lyu [49]. After EBSD acquisition, additional lines appear (labelled "New lines"), remaining parallel to the initial ones and preserving the same two orientations. A linear crack develops along one of the pre-existing defects (labelled "Linear crack" in Figure 3(d)).

Beyond these defects, bright high-contrast edges are observed to form and thicken along GBs. In Figure 3(a), one GB is naturally visible due to a fine (~200 nm thick) bright edge, while other GBs — identified from EBSD data — are reproduced as dashed blue lines in Figure 3(a). After EBSD acquisition (Figure 3(b)), bright edges appear at all GBs. The initially visible edge transforms into a ~2 µm thick cracked phase (labelled accordingly), with numerous transverse cracks. A similar cracked phase is also present in the second grain (Figure 3(d)). Notably, comparable
<figure><img src="image_4.png" /><figcaption>Figure 2: Hydride degradation during EBSD analysis. (a,b) BSE micrographs acquired before and after EBSD, revealing bright-contrast features at GBs and linear cracks. The dotted white square in (a) indicates the region where the EBSD maps were acquired. The red and green dotted squares will be examined in more detail in Figure 3. (c,d) Corresponding EBSD phase maps after the first and second acquisitions, respectively, highlighting the progressive transformation from $MgH_2$ to Mg.</figcaption></figure>

contrast and microcracks are visible along many of the linear defects described above.

The orientation of these defects has been then compared to the crystallographic orientations of the $MgH_2$ phase to assess their relation to specific lattice planes. Figure 4(a) shows a portion of the EBSD map containing both hydride grains of Figure 3, highlighted in colour, while the rest of the map is in grayscale. The trace of a crystallographic plane is defined by its intersection with the observed cross-section. All plane families were examined to identify the best match between defect orientation and crystallographic traces. The highest correlation ($r = 0.999$; mean error =1.25°) is obtained for the {110} family. The traces of the {110} planes are drawn as blue

lines in Figure 4(a) and transposed onto the BSE micrographs in Figure 4(b) and Figure 4(c). The perfect alignment confirms that the defects are contained within the (110) and (\bar{1}10) planes of the respective hydride grains.

### 3.3. Desorption-induced cracked phase

Figure 5 and Figure 6 respectively show two SEM micrographs and two EDX maps of the cracked phase previously identified in Figure 3(b). Both the Mg and $MgH_2$ phases in Figure 5 can be identified by their BSE contrast. The SE micrograph in Figure 5(b) reveals well-defined interfaces with adjacent $MgH_2$ and Mg grains. Cracks are present within Mg phase that formed at the former $MgH_2$ GBs.
<figure><img src="image_5.png" /><figcaption>Figure 3: Microstructure evolution in hydrides after the first EBSD acquisition. (a,b) BSE micrographs of the same region (red dotted square in Figure 2(a)), taken before and after the EBSD acquisition. (c,d) BSE micrographs of another region (green dotted square in Figure 2(a)), acquired before and after the EBSD acquisition, showing similar damage features.</figcaption></figure>

The EDX map in Figure 6(a) shows oxygen concentration variations in shades of green, clearly indicating the presence of oxygen within the cracked phase. Light green lines also trace the linear defects previously observed in Figure 3(b). In the second grain shown in Figure 5(b), oxygen enrichment similarly follows the linear defects.

To further investigate the damage induced in the hydrides after SEM observations, an attempt was made to reproduce these effects under a localised electron beam scan. A hydride-rich region in the $Mg-(Ti_3AlC_2)_{15\%}$ composite was exposed to a stationary electron beam (20 kV acceleration voltage, 120 µm aperture, 15 mm working distance) — identical to the EBSD acquisition parameters — in order to maximise the accumulated electron dose on the hydrides. The exposed area is indicated by a dashed blue square in Figure 7. Starting from the initial state in Figure 7(a), the exposure times for Figure 7(b), (d), (e), (f)

and (c) were 2 min, 4 min, 7 min, 13 min and 15 min, respectively. Micrographs were taken at two different magnifications.

After only 2 min of exposure (Figure 7(b)), a circular region of bright contrast is visible, extending beyond the exposed square. By 15 min (Figure 7(c)), this bright region — labelled “circular bleaching” — extends approximately 5 µm beyond the scanned area in all directions. Figures 7(d–f) show higher-magnification BSE micrographs acquired inside the exposed area. After 4 min (Figure 7(d)), fine cracks appear, as well as deeper intergranular cracks (labelled accordingly) following the bright edges marking hydride along GBs reported in Figure 3(a). At longer exposure times (Figures 7(e,f)), both fine and intergranular cracks propagate further.

These observations demonstrate that the vacuum conditions and the dose delivered by the electron beam are
<figure><img src="image_6.png" /><figcaption>Figure 4: Crystallographic {110} plane traces aligning with linear defects in hydrides. (a) EBSD map showing {110} plane traces in blue. (b,c) BSE micrographs in which the {110} plane traces, reproduced from the EBSD analysis, are superimposed to highlight their alignment with the observed defects.</figcaption></figure>

<figure><img src="image_7.png" /><figcaption>Figure 5: Cracked phase identification. (a) BSE and (b) SE high-magnification micrographs showing the cracks within the Mg phase present along the former $MgH_2$ GBs. The micrographs were taken after EBSD acquisition.</figcaption></figure>

sufficient to destabilise the hydride. In addition, the presence of oxygen within the cracks may have oxidized the Mg phase. The implications of this phenomenon for understanding $MgH_2$ desorption mechanisms will be discussed in the following section.

## 4. Discussion

In the case of the destabilization of $MgH_2$ in the SEM, the combined effects of SEM vacuum and electron beam, at room temperature and near-zero bar pressure (SEM residual pressure: ~$1 \times 10^{-4}$ Pa), create thermodynamic conditions favorable to H desorption. It is therefore unsurprising that desorption occurs after some time in the SEM. Owing to the moderate desorption kinetics under these conditions, hydrides remain relatively stable long enough to allow their characterization before damage becomes excessive.

From the damage patterns left by EBSD scanning (Fig-

<figure><img src="image_8.png" /><figcaption>Figure 6: Cracked phase identification. EDX elemental maps revealing oxygen enrichment within the cracked phase and along linear defects.</figcaption></figure>

ure 3) and by prolonged exposure to a stationary electron beam (Figure 7), it is evident that the electron beam plays a decisive role in destabilizing the hydrides. Similar to temperature, the beam energy acts as a kinetic factor that helps to overcome the activation energy barrier for H desorption.

Regarding the so-called desorption "windows," Lyu [49] reports in a review on the nucleation and growth behavior of $MgH_2$ the appearance of window-like structures in desorbing hydrides, which serve as H diffusion channels. These features were first observed by Evard *et al.* [50] using optical microscopy. No further detailed characterization of these structures could be found in the literature. In the present work, these desorption windows exhibit patterns identical to the linear defects observed in Figure 3. Because the Mg phase was identified along these linear defects, their role as desorption channels is here confirmed. Furthermore, these linear defects were indexed in the {110} planes of the $MgH_2$ lattice (Figure 4). The {110} planes of the rutile $MgH_2$ structure have the highest density of Mg atoms. The $MgH_2 \rightarrow Mg$ phase transformation involves a volume contraction of approximately 30%, and it is likely that the high Mg atomic density on these planes facilitates this transformation. As a result, the phase transformation propagates preferentially along a given {110} plane, introducing a linear defect into the hydride lattice.

The newly formed Mg phase exhibits fast H diffusion and thus acts as a channel. In addition, the growth of the Mg phase and the associated shrinkage generate tensile residual stress fields, eventually relaxing by cracking, which further assists in the evacuation of $H_2$.

Concerning the mechanisms enabling H transport during $MgH_2$ desorption, the results indicate that the transformation proceeds preferentially along {110} planes be-
<figure><img src="image_9.png" /><figcaption>Figure 7: *In-situ* SEM observation of MgH₂ destabilization under localized electron beam scanning (20 kV, 120 µm aperture). (a–c) BSE micrographs acquired after 0, 2, and 15 min of exposure, respectively. (d–f) BSE micrographs of the exposed region during scanning after 4, 7, and 13 min of exposure, respectively.</figcaption></figure>

cause of their high Mg atomic density. Linear defects along these planes form desorption channels, combining (i) the fast H diffusion through the metallic Mg phase and (ii) the stress-assisted cracking of the hydride matrix. Together, these mechanisms facilitate H transport and evacuation during desorption.

Finally, based on these observations, a coherent scenario can be proposed. All processes occur at interfaces: the interface between Mg and Ti₃AlC₂ acts as a preferential site for the forward phase transformation (Mg → MgH₂), while the GBs of MgH₂ serve as nucleation sites for the reverse, crystallographically controlled transformation (MgH₂ → Mg).

Additionally, Mg oxidation may contribute to this scenario. Oxidation can occur at GBs or at interfaces between MAX phases and Mg due to unavoidable oxygen contamination. Such oxidation could promote the phase transformation toward MgH₂ by locally altering thermodynamic or kinetic conditions. Notably, Mg oxidation has been shown to proceed relatively rapidly even under vacuum conditions [51]. At low temperature and low O₂ partial pressure, oxide layers remain thin. However, during hydrogenation, hydride growth may generate stresses sufficient to disrupt the oxide layer, thereby facilitating H

transport and subsequent hydride formation [51].

## 5. Conclusion

In conclusion, an in-depth microstructural and crystallographic analysis, complemented by *in-situ* desorption experiments under a localised electron beam in the SEM, has enabled direct visualisation and characterisation of MgH₂ destabilisation in Mg – (Ti₃AlC₂)₁₅%.

Linear defects — previously reported but not characterized in [49] — within the MgH₂ phase, forming windows, were characterised and found to be aligned along {110} planes of the MgH₂ lattice. As these features coincide with the transformation sites MgH₂ → Mg, they are confirmed as desorption channels. This behaviour is promoted by the volume contraction accompanying the phase transformation, which generates residual stress fields, and by the high Mg atomic density on the {110} planes of the rutile structure.

In the present case, and within our proposed scenario, it is reasonable to suggest that the interface between Ti₃AlC₂ and Mg may localize the initiation of the transformation to MgH₂ and/or slightly destabilize the hydride, thereby enabling the capture of the desorption process un-
der the electron beam. This is consistent with the decrease
in desorption temperature reported by Lakhnik *et al.* [13]
and attributed to MAX particles acting as nucleation sites
for Mg, in agreement with JMAK analysis [35].

Overall, our results establish a direct link between
the formation of crystallographically aligned desorption
channels and hydrogen release in MgH₂, reveal the role
of atomic density in facilitating the transformation, and
demonstrate how Ti₃AlC₂ particles modulate these path-
ways. These insights provide a structural basis for engi-
neering faster H-release kinetics in Mg-based hydrides.

Acknowledgments

This project received financial support from the CNRS
through the MITI interdisciplinary program, and from the
French State through the "Investment for the Future" pro-
gram operated by the French National Research Agency
(ANR), under reference ANR-11-LABX-0008-01 (LabEx
DAMAS). The authors gratefully acknowledge the use of
the MicroMat experimental facilities at LEM3 (Univer-
sité de Lorraine – CNRS – ENSAM) and thank Prof.
Wenbo Yu (School of Mechanical and Electronic Control
Engineering, Beijing Jiaotong Univerisity, Beijing, 100044,
China) for providing the synthesized samples.

Author contributions

A. Heinzelmeier: Conceptualization; Data curation; Formal analysis; Investigation; Methodology; Validation; Visualization; Roles/Writing - original draft.

F. Cuevas: Resources; Validation; and Writing - review
& editing.

T. Grosdidier: Conceptualization; Formal analysis In-
vestigation; Methodology; Validation; Visualization; and
Writing - review & editing.

A. Guitton: Conceptualization; Formal analysis; Fund-
ing acquisition; Investigation; Methodology; Validation;
Visualization; Roles/Writing - original draft; and Writing
- review & editing.

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