Formation of the Ternary Complex Hydride Mg2FeH6 from Magnesium

Article pubs.acs.org/JPCC

Formation of the Ternary Complex Hydride Mg2FeH6 from Magnesium Hydride (β-MgH2) and Iron: An Electron Microscopy and Energy-Loss Spectroscopy Study Mohsen Danaie,*,† Alexandre Augusto Cesario Asselli,‡,§ Jacques Huot,‡ and Gianluigi A. Botton† †

Department of Materials Science and Engineering, Brockhouse Institute for Materials Research and Canadian Centre for Electron Microscopy, McMaster University, Hamilton, Ontario, Canada ‡ Physics Department and Institut de Recherche sur l’Hydrogène (IRH), Université du Québec à Trois-Rivières, Trois-Rivières, Quèbec, Canada § Programa de Pós-Graduaçaõ em Ciência e Engenharia de Materias, Universidade Federal de São Carlos, São Carlos, SP, Brazil ABSTRACT: We examined the formation of the ternary complex hydride phase Mg2FeH6 during the thermal hydrogen absorption of a ball-milled powder mixture of MgH2 and Fe. Analytical measurements, scanning transmission electron microscopy, and electron energy-loss spectroscopy, with the sample cooled to liquid nitrogen temperature, were utilized to identify the various phases present and to study the features of this phase transformation. The low-loss region of the electron energy-loss spectrum was mainly used to distinguish various constituents. Mg2FeH6 was initially formed during ball milling of MgH2 and Fe, demonstrating a co-continuous structure with MgH2 and Mg. Upon the first desorption, this phase was transformed into elemental Mg and Fe. During the initial stages of the subsequent thermal absorption, MgH2 was first formed with enhanced kinetics because of the presence of Fe. This was followed by the nucleation of Mg2FeH6 between MgH2 and Fe and its growth with a columnar morphology. This morphology was dictated by the diffusion direction of the atomic hydrogen from the catalytic Fe cap. As the Mg2FeH6 columns grew, the capping Fe particle and the MgH2 substrate were consumed. When the system was maintained at high temperature for a long time, these columnar Mg2FeH6 structures coalesced into a semispherical morphology.

1. INTRODUCTION Magnesium-based hydride systems have long been considered as potential candidates for solid-state hydrogen storage applications. Pure magnesium forms a stable hydride, tetragonal β-MgH2, with 7.6 wt % hydrogen content and a standard heat of formation of around −74.4 kJ/mol.1 The abundance of magnesium resources has also led to relatively lower costs, which is a major advantage for industrial applications. The main drawback for a large group of Mgbased hydrides is the excessive stability of these phases. High operating temperature has mostly limited the viability for this group of hydrides for stationary hydrogen storage. Faster cycling operation with enhanced kinetics has been demonstrated with the advent of efficient catalyst phases.2−4 The ternary hydride phase Mg2FeH6 belongs to a category of hydrides known as “complex metal hydrides”, Mmδ+[THn]δ−, where M is an alkaline, alkaline earth, or divalent rare-earth metal and T is a transition metal.5 The reason behind this nomenclature is that an important building block for the crystal structure of this group of hydrides is the metal−hydrogen complex centered by the transition metal, that is, [THn]δ−. For the case of Mg2FeH6, this complex is the octahedral anion [FeH6]4−.6 Although the gravimetric hydrogen capacity of Mg2FeH6 is lower than that of pure magnesium hydride (5.4 versus 7.6 wt %), its volumetric hydrogen content exceeds that of MgH2 (9.1 × 1022 versus 6.5 × 1022 atom/cm3). The experimentally measured heat © 2012 American Chemical Society

of formation values for this complex hydride vary in the literature (for example, −98 ± 3,6 −86 ± 6,7 and −81 ± 288 kJ/mol), but the consensus is that it is more stable than β-MgH2. These properties have generated interest in Mg2FeH6 both for solid-state hydrogen storage and also for heat storage applications.9 In 1984, Didisheim et al.6 reported the first synthesis and crystal structure analysis of Mg2FeH6. They later refined their synthesis technique for higher yield,10 based on high-pressure sintering of the compacted elemental powders (2:1 mixture of Mg to Fe) followed by hydrogen charge−discharge cycling (90 bar H2 pressure, 450 °C). The challenge for the synthesis of Mg2FeH6 stems from the facts that magnesium and iron are immiscible in their elemental state (maximum solid solubility of Fe in Mg is 0.00041 at. % Fe) and the intermetallic phase Mg2Fe is not thermodynamically stable.11 This is unlike the case for Mg2NiH4, which belongs to the same group of complex metal hydrides and has a high-temperature polymorph similar in crystal structure to Mg2FeH6.12 Mg2Ni is a stable intermetallic phase, and Mg2NiH4 can be formed simply by the direct hydrogenation of this phase.13 As an alternative and faster route to the synthesis of Mg2FeH6, hydrogenation of the mechanically milled magnesium and iron Received: August 29, 2012 Revised: November 6, 2012 Published: November 9, 2012 25701

dx.doi.org/10.1021/jp3085843 | J. Phys. Chem. C 2012, 116, 25701−25714

The Journal of Physical Chemistry C

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powder mixture has been proven to be feasible.7 Small amounts of iron can be dissolved in metallic magnesium through mechanical alloying, potentially as clusters within the hexagonal close-packed matrix.14 Huot et al.15 demonstrated that the milling atmosphere is a crucial factor, with hydrogen presence enhancing the final yield of Mg2FeH6 significantly. They later showed that starting with a mixture of MgH2 (rather than Mg) and Fe powder, Mg2FeH6 could be formed directly during mechanical milling.16 Close analysis of the data in the cited study led the authors to conclude that Mg2FeH6 is formed from MgH2, and not from metallic Mg. X-ray diffraction measurements on sequentially interrupted ball-milling batches of magnesium and iron powders under hydrogen atmosphere confirmed this observation.17 Polanski et al.18 further corroborated this finding by an in-situ synchrotron X-ray diffraction study of various starting powder combinations (Mg/Fe or MgH2/Fe). They conclusively showed that, regardless of the use of metallic magnesium or MgH2 powder with iron, after cycling and during absorption, MgH2 forms prior to Mg2FeH6, suggesting the following reaction for the formation of Mg2FeH6 phase 2MgH2 + Fe + H 2(g) → Mg 2FeH6

We used X-ray diffraction (Bruker D8-Focus diffractometer with Cu Kα radiation, λ = 1.5406 Å) for phase identification. XRD measurements were performed under ambient conditions. We also performed Rietveld quantitative analysis on the XRD diffractograms to obtain the abundances of the phases present. The commercial software Topas was used for this analysis (TOPAS V4: General Profile and Structure Analysis Software for Powder Diffraction Data, Bruker AXS, Karlsruhe, Germany, 2008). For microscopy analysis, powder samples were dispersed, with no added solvents, onto ultrathin amorphous carbon grids. This was done after gently crushing the powder using an agate mortar and pestle. Sample preparation was performed in a pure and dry argon glovebox (contents of both O2 and H2O below 0.1 ppm). Specimens were transferred to the microscope room in a sealed plastic bag and were exposed to air for only a short time during sample loading. All microscopy sessions were performed using a cryogenically cooled sample holder (T = 95 K) to increase the stability time of the hydride phase under the electron beam. At room temperature, β-MgH2 quickly transforms to a mixture of Mg and MgO upon exposure to an electron beam.20 For microstructure characterization of the powders, we used a FEI Titan 80-300 (scanning) transmission electron microscope, equipped with an X-ray energy-dispersive spectrometer [Oxford Inca, Si(Li) detector] and a Gatan Image Filter (GIF) electron energy-loss spectrometer. The accelerating voltage was set at 300 kV during our experiments. The energy spread of the primary electron beam, measured as the full-width at half-maximum of the zero-loss peak in vacuum, was around 0.7 eV. An energy dispersion of 0.03 eV per pixel was used for electron energy-loss spectroscopy (EELS). Two general techniques were utilized to fully characterize the microstructure of the samples: (1) bright-field (BF) and darkfield (DF) imaging, along with selected-area diffraction (SAD) pattern analysis, and (2) spectrum imaging in scanning transmission electron microscopy (STEM) mode, that is, STEM/EELS. In the former methodology, the crystal structures of the phases present were locally captured. Using BF/DF imaging, by selecting and tilting various diffracting spots to the optical axis, the distributions of various phases were studied. In the latter technique, the microscope was operated in STEM mode, where a fine electron probe was scanned on the sample, rather than illuminating the whole sample with a parallel beam as in conventional TEM. In this way, the total exposure of the sensitive hydride phase to the electron beam was greatly reduced. A high-angle annular dark field (HAADF) detector was used in this mode to obtain an image of the specimen from the highly scattered electrons. The contrast observed in the final image is mostly due to variation in atomic number (Z), with a higher Z resulting in a brighter intensity. Using the spectrum imaging routine, we acquired electron energy-loss (EEL) spectra at each defined pixel in a selected region of the sample. In some cases, we applied both of these techniques to a common region of the sample. For the EELS measurements, we focused on the low-loss region of the spectra (0−50 eV energy loss), where the plasmon excitation peaks of various phases can be observed. Assuming a free-electron model for the valence electrons of the solid, one can express the volume plasmon excitation energy (Ep) as (page 136 of ref 21):

(1)

This reaction pathway was also supported by quantitative analysis of the hydrogen uptake kinetics during reactive mechanical milling of Mg/Fe (2:1 ratio) powder mixtures.19 In the present study, we focused on the microscopy characterization of this phase transformation. A detailed study of how the phase transition from MgH2 and Fe into Mg2FeH6 proceeds, in terms of both morphology and phase distribution, could potentially lead to a better understanding of this process. Such an understanding could then be utilized to design new synthesis procedures with enhanced yield of the ternary hydride phase.

2. EXPERIMENTAL METHODS We milled a mixture of MgH2 and Fe powders with the molar ratio of 4:1, respectively. As starting materials, magnesium hydride powder (300 mesh, 98.0% purity) and iron powder (−22 mesh, 99.998% purity), both provided by Alfa Aesar, were used. High-energy ball milling of the MgH2/Fe mixture was carried out in a Fritsch Pulverisette 4 planetary mill, using balls and a bowl made of hardened stainless steel. The sample was milled at the rotational speed of 600 rpm for 48 h under an argon atmosphere. The ball-to-powder weight ratio was 40:1. All handling procedures were performed inside an argon-filled glovebox. To study the morphology and characteristics of the complex hydride phase at various stages, we focused on three powder samples for our microscopy measurements: sample 1, as-milled sample; sample 2, powder specimen “quenched” at the early stages of Mg2FeH6 formation; and sample 3, fully hydrogenated sample. For the preparation of samples 2 and 3, the as-milled powder was sealed in the stainless steel sample holder within the glovebox and then transferred to a Sieverts apparatus. In both cases, the powders were first fully desorbed at 673 K under 1 bar initial pressure of H2. For the hydrogen absorption step, the temperature was kept constant (at 673 K), and an initial hydrogen pressure of 40 bar was applied to the samples. For sample 2, we quenched the specimen after 2 min of hydrogen absorption. Quenching was achieved by closing the sample valve to the hydrogen gas source, retracting the heating unit from the sample holder area, and cooling the sample holder to room temperature in a few seconds using a cold water bath. For sample 3, we maintained the powder sample in the absorption step for 16 h.

Ep = (28.82 eV)(zρ /A)0.5 25702

(2)



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where z is the number of valence electrons per molecule, ρ is the density (in g/cm3), and A is the molecular weight. Because the density of the valence electrons undergoes a dramatic change between a metallic phase and the corresponding hydride phase, the plasmon excitation peaks of these phases (metal versus hydride) demonstrate a significant shift in energy. This shift can be utilized in mapping the phases that give rise to them through either energy-filtered TEM (EFTEM) or spectrum imaging in STEM (STEM/EELS). In an earlier study,22 EFTEM was used to map the hydride and metallic phases in a ball-milled MgH2/ Mg system. In this study, we performed spectrum imaging (STEM/EELS) to spatially map Mg, MgH2, and Mg2FeH6 within the microstructure. X-ray energy-dispersive spectrometry (XEDS) elemental maps (in STEM mode) were also occasionally acquired. We used the commercial software JEMS for simulating the selected-area diffraction patterns.

the system toward full hydrogenation capacity. To study the microstructure of the system in transition from the fast to the slow regimes, we quenched a powder sample at the designated point in Figure 1A, denoted as sample 2. The plot in Figure 1B shows the hydrogen uptake kinetics for this sample prior to quenching. The powder sample held at the absorption step for 16 h (i.e., that corresponding to the full plot in Figure 1A) was also chosen for microstructure characterization (designated here as sample 3). XRD results for the three powder samples are presented in Figure 2. Table 1 reports the outcome of the Rietveld refinement

3. RESULTS 3.1. Hydrogen Sorption and XRD Analysis. The hydrogen absorption step of the as-milled powder sample is presented in Figure 1. Figure 1A depicts the hydrogen uptake kinetics over

Figure 1. Hydrogen sorption of the as-milled MgH2/Fe powder sample. (A) Full hydrogenation plot with the sample held in the absorption step for 16 h (sample 3). The stage where sample 2 was quenched during absorption is marked on the plot with an arrow. (B) Hydrogen absorption of powder sample 2 prior to quenching. Note the difference in units for the horizontal (time) axes in panels A and B.

Figure 2. XRD results for: (A) sample 1, as-milled MgH2/Fe; (B) sample 2, fully desorbed and then hydrogenated for 2 min at 673 K and 40 bar H2 pressure; and (C) sample 3, maintained in the hydrogen absorption step for 16 h.

the period of 16 h. Two distinct regimes can be observed during hydrogenation of the as-milled sample, namely, initial fast hydrogen absorption to around 5 wt % H within the first 2−3 min, followed by a far slower step spanning about 15 h that takes

analysis of our diffraction data, listing the quantities of the phases present in each sample. The data in Figure 2A show that Mg2FeH6 formed during the milling process, with the peaks for 25703



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Fe(4a), and H(24e) with a displacement of x = 0.242).6 Faint diffraction peaks from metallic magnesium can also be observed in this SAD pattern. The dark-field image, presented in Figure 3B, was formed using a selection of diffraction rings of the Mg 2FeH6 phase, demonstrating fine and equiaxed grain morphology. The low-loss EEL spectra of this area (referred to as Area 1 in Figure 4A), along with two other areas, are presented in Figure 4A. These spectra were all acquired in TEM mode under identical microscope settings. The corresponding selected-area electron diffraction data for areas 2 and 3 are shown in panels B and C, respectively, of Figure 4, with the simulated ring pattern of the ternary hydride phase also included. As for area 1, in these two areas, Mg2FeH6 is the dominant phase present. The EEL spectra in Figure 4A are normalized with respect to the corresponding integrated zero-loss peak (ZLP) intensity (I0) in each case. Because the ratio of the area under the zero-loss peak (I0) to the area under the whole EEL spectrum (It) is exponentially related to the sample thickness [It/I0 = exp(t/λ), where λ is the total inelastic mean free path; see page 294 of ref 21), it can be quickly recognized that, from area 3 to areas 2 and 1, the sampled regions became gradually thinner. However, independent of thickness effects, two main peaks are present in these spectra: one broad peak at 16.5 eV and another (appearing as a small shoulder to the former) sharp peak at around 10.6 eV. The sharp plasmon peak at 10.6 eV is the well-established volume plasmon peak for the metallic Mg phase.22−24 Also, in the cited studies, it was found that the volume plasmon peak for the β-MgH2 phase is close to 14 eV. The characteristic plasmon peak for MgO is close to 22 eV.25 Given all of these values and the fact that the majority of the signal in our acquired SAD patterns from all three areas was from Mg2FeH6 (Figures 3C and 4B,C), it is safe to conclude that the peak at 16.5 eV corresponds to the plasmon excitation of the ternary hydride phase. With this new information regarding the plasmon peak energy (Ep) of the Mg2FeH6 phase, it is possible to map this phase and also other phases present with known Ep values. For enhanced spatial resolution, this is best performed in STEM mode. Figure 5A presents the HAADF-STEM image of an as-milled particle. A spectrum image was acquired for this particle; that is, a low-loss EEL spectrum was gathered for every single pixel of the scanned HAADF image signal. The graph in Figure 5B presents three spectra from three different locations on this particle. For site 1 (red spectrum), one can observe the plasmons for Mg (11 eV)

Table 1. Relative Quantities of the Phases Present in the Three Powder Samplesa sample 1, as-milled 2, partially absorbed 3, fully absorbed

Mg2FeH6

MgH2

Mg

Fe

MgO

68 (36) 13 (4)

16 (35) 43 (60)

8 (19) 5 (8)

4 (4) 29 (19)

4 (6) 10 (9)

49 (22)

32 (59)

0 (0)

8 (7)

10 (12)

Values are in weight percentages (estimated error of ±1%), with the equivalent (mol %) given in parentheses.

a

this complex hydride prominently present. After complete desorption of this sample, the present hydride phases (β-MgH2 and Mg2FeH6) transformed into the elemental phases Mg and Fe. Figure 2B shows the powder diffraction data for sample 2, which was hydrogenated for only 2 min at 673 K after complete desorption. The data for the completely hydrogenated sample (sample 3) are shown in Figure 2C. As can be seen in Table 1, the highest content of the Mg2FeH6 phase was found in the as-milled state. In sample 2, MgH2 was the dominant hydride phase, with iron mostly present in elemental form. After 16 h of being maintained in the hydrogen absorption step (sample 3), the relative contribution of the complex hydride phase increased significantly. No metallic magnesium was present, which indicates that this sample was close to full hydrogenation. Some residual metallic Fe (8 wt %) was still present that could react with MgH2 following reaction 1 and form the ternary hydride phase. Mg2FeH6 was the most abundant hydride phase in sample 3. Based on the XRD results, we concluded that, during the first fast regime of hydrogenation, mainly MgH2 forms. The formation of the ternary hydride phase, Mg2FeH6, can be associated with the slower second portion of the absorption curve presented in Figure 1A. Hence, sample 2 captures the early stages of Mg2FeH6 formation within the microstructure. 3.2. Electron Microscopy Analysis (TEM and STEM/ EELS). 3.2.1. Sample 1: As-Milled MgH2 and Fe (4:1) Powder Sample. The majority of this powder sample (in weight percentage) was the ternary complex hydride phase. A representative region of this specimen is presented in Figure 3. The selected-area diffraction (SAD) pattern confirms that this area was almost exclusively Mg2FeH6. (See the comparison with the simulated ring pattern for Mg2FeH6 in Figure 3C.) Mg2FeH6 has a cubic crystal structure, with a lattice parameter of 6.437 Å [space group Fm3̅m (No. 225); Wyckoff positions Mg(8c),

Figure 3. Conventional TEM results showing typical microstructure of sample 1, as-milled powder: (A) bright-field, (B) dark-field, (C) selected-area diffraction pattern with the simulated ring pattern of Mg2FeH6. 25704



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Figure 4. (A) Low-loss EEL spectra acquired in TEM mode from three different regions of the as-milled sample (sample 1). Intensities of the spectra were normalized to the area underneath the corresponding zero-loss peaks (I0). Area 1 is the same as the area shown in Figure 3. SAD patterns corresponding to areas 2 and 3 are presented in panels B and C, respectively. In both cases, the simulated ring pattern of Mg2FeH6 is also included.

Figure 5. STEM/EELS results from the as-milled powder sample (sample 1). (A) HAADF image, with the area selected for spectrum imaging marked with the frame. (B) HAADF signal with marked locations and corresponding low-loss EEL spectra (similar colors for the spectra and the corresponding locations). (C) Maps with various selected energy ranges for highlighting Mg (10.1−12.1 eV), MgH2 + Mg2FeH6 (14.3−16.3 eV), and MgO and amorphous C (21.7−23.7 eV).

In Figure 5C, the signal intensities for various energy ranges, corresponding to the given phases, are mapped as energy-filtered images. Because of the proximity of the peaks for MgH2 and Mg2FeH6, one map shows the distribution of these two phases.

and MgH2 (14 eV). For location 2 (green spectrum), only a peak at 16.4 eV is present, which we already identified to be from Mg2FeH6. For location 3 (black spectrum), a broad peak at around 22 eV is present, corresponding to the Ep value for MgO. 25705



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Figure 6. Microstructure of sample 2: (A) bright-field TEM image, (B) SAD pattern, and (C) SAD pattern with Fe (bcc) ring pattern simulation. (D) Same area in HAADF-STEM. XEDS elemental maps of the region in the red frame in panel D are presented for (E) Mg and (F) Fe. Green arrows in panels A and D point to the same location in the sample.

(The energy ranges used for the maps are given in the figure caption.) It can be observed that the distributions of Mg and the two hydride phases show co-continuous microstructures, with both the metallic and hydride phases being evenly distributed across the particle. The map for MgO shows a stronger signal at the edges. Because the thin oxide layer covers the surface, the interaction volume of MgO with the electron beam is the largest on the edge. It can be noted that the amorphous carbon from the support film is also highlighted in this last map. This is because the broad plasmon peak for amorphous carbon (used to support the powder sample on the copper grid) is also between 20 and 34 eV.26

3.2.2. Sample 2: Absorbed for 2 min at 673 K [P(H2) = 40 bar]. This sample was fully desorbed and then quenched during the absorption step at the early stages of Mg2FeH6 formation. From the XRD data and Table 1, only 13 wt % of this sample was the ternary complex hydride phase, with 43 and 29 wt % comprising MgH2 and elemental Fe, respectively. Figure 6A is a TEM bright-field image of this sample. Analysis of the SAD pattern (Figure 6B) shows that a mixture of Fe, β-MgH2, and Mg2FeH6 phases was present (Figure 6C). The same area imaged in STEM mode using the HAADF detector is shown in Figure 6D. The bright intensity of the spherical “globules” dispersed across the larger particle indicates the higher atomic number (Z) 25706



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the other hand, was acquired from within the columnar section. The plasmon peak for this spectrum is located at 18 eV. Because we knew from electron diffraction data that Mg2FeH6 phase was present in this sample (see Table 1 and Figure 6C), our anticipation was to observe the low-loss peak that we identified for this phase in sample 1, namely, the peak at 16.5 eV. On the contrary, in sample 2, the only plasmon peak observed, other than the known peak for the volume plasmon of β-MgH2 at around 14 eV, was found at 18 eV. Also, this plasmon at 18 eV was spatially correlated with the columnar features that could be located across the microstructure of sample 2. The two images presented on the right side of Figure 7C map the intensities corresponding to the MgH2 plasmon and the 18 eV peak. It can be observed that the protruding “column” is highlighted entirely in this latter map (marked with an arrow). We thus interpret this columnar morphology as the initial growth stages of the ternary complex hydride phase, Mg2FeH6. The potential reasons behind the observed discrepancy between the plasmon peak location of this phase in the as-milled state (sample 1) compared to the freshly formed phase (sample 2) will be revisited in the Discussion section. Another example of the columnar growth of the Mg2FeH6 phase is presented in the STEM-HAADF image of Figure 8A. As above, the column was capped with an iron particle. In this specific case, the columns also showed a tapered geometry, with smaller cross section closer to the Fe particle (two cases marked by arrows). The spectrum imaging data acquired from the marked region in panel A are presented in Figure 8B. The HAADF signal is shown at the top left. The low-loss EEL spectrum from the location marked in the HAADF signal is presented in the graph at the bottom, indeed showing that the plasmon peak for the growing Mg2FeH6 phase is located at 18 eV. The energy-filtered intensity image, using this plasmon peak, is presented at the top right image in Figure 8B. As can be observed in this map, the growing columns are highlighted entirely, indicating a consistency in volume plasmon peak location across this phase. To demonstrate the incorporation of iron atoms into the growing Mg2FeH6 columns, we acquired X-ray energydispersive spectra from this area. Figure 8C presents the XEDS data for the two selected locations (marked in Figure 8A with crosshairs). Location 1 was on one of the Fe particles, and location 2 was within one of the growing columns. In the spectrum for data point 1, a strong Fe signal can be observed. This signal is also present, although at a lower intensity, in spectrum 2, that is, in the growing columnar phase. The signal for Mg is present in both locations. For the case of data point 1, this could potentially be caused by the underlying Mg-containing phases. A remark should be made here regarding the Fe plasmon and why it was not used here for mapping this phase. Iron, in its elemental bcc state, gives rise to a broad plasmon peak at around 23.0 eV (data not shown here), in agreement with the literature (see page 420 in ref 21). Because we had overlapping peaks around this same energy value in our system, for example, MgO (Ep = 22 eV) and amorphous C (Ep = 20−34 eV), to avoid potential confusion, we decided not to use the Fe plasmon signal for mapping. Alternatively, the pronounced Z difference between iron and the rest of the microstructure resulted in a distinctive contrast for this phase in the STEM-HAADF signal. Hence, we used this contrast difference to locate the Fe particles within the studied microstructure. XEDS was also used to identify the presence of Fe in the ternary hydride phase, with Figure 8C being an example.

of that phase. Panels E and F of Figure 6 show the XEDS elemental maps of Mg and Fe, respectively. The iron map confirms that the bright areas are indeed Fe-rich. The diffraction pattern (Figure 6C) and the presence of the body-centered cubic (bcc) Fe diffraction ring pattern confirm that the fine particles with the brighter intensity in the HAADF signal are elemental Fe. Because this sample was initially fully desorbed, a high concentration of metallic Fe from the decomposition of the Mg2FeH6 phase was present (see Table 1). Polanski et al.27 reported a similar type of morphology for Fe after complete desorption of the Mg2FeH6 phase. We also point out a columnar feature protruding from the main particle in the center of the view here, marked with green arrows in both the BF (Figure 6A) and HAADF (Figure 6D) images. This type of morphology was encountered in other areas of sample 2 as well. Figure 7A shows the STEM-HAADF signal from another area of this sample. Panel B of Figure 7 is a higher-magnification

Figure 7. Microstructure of sample 2 (A) under STEM-HAADF imaging conditions. (B) Higher-magnification view of the red frame in panel A. (C) Corresponding spectrum imaging data. The graph on left shows two spectra from two locations, designated in the HAADF signal. The two maps on the right show the intensity of the signal in the selected energy ranges for the two phases: MgH2, 13.3−15.5 eV; Mg2FeH6, 17.3−19.5 eV.

image of the marked area in panel A. The arrow in Figure 7B points to another example of these columnar features. The brighter HAADF intensity of the uppermost section of this column indicates that it is capped with Fe. Figure 7C presents STEM/EELS results from this area. Two EEL spectra are plotted in the graph on the left, with corresponding beam positions marked on the HAADF signal. Spectrum 1, acquired from the interior section of the particle, shows a peak at around 14.5 eV corresponding to the plasmon peak for MgH2. Spectrum 2, on 25707



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Figure 8. (A) Columnar growth of the Mg2FeH6 phase imaged using STEM-HAADF. (B) Spectrum imaging data from the marked box in panel A. The HAADF image is shown on top left, and the energy-filtered intensity map for the Mg2FeH6 phase (16.9−19.0 eV) is the top right image. The graph at the bottom shows a plot of the low-loss EEL spectrum from the location marked in the HAADF signal (top left image). XEDS spectra from two locations, marked with + in panel A, are presented in panel C. Copper peaks are from the supporting grid.

Figure 9. Microstructure of sample 2 under STEM-HAADF imaging conditions . Spectrum imaging results from the red frame in (A) is presented in (B) and (C). (B) Two low-loss EEL spectra from two different locations on the specimen (marked on the HAADF signal on the right side). (C) Two maps from two energy ranges corresponding to MgH2 (12.9−15.4 eV) and Mg2FeH6 (17.2−19.8 eV).

Figure 9A shows the STEM-HAADF image of an area in sample 2 with a large population of the columnar features. The low-loss EEL spectra from the columns show the peak at 18.0 eV, which we interpreted above as the plasmon for the newly formed Mg2FeH6. Two examples are given in Figure 9B, with the

spectrum from location 1 showing only the 18.0 eV peak. The spectrum acquired at site 2 shows the MgH2 plasmon peak (at 14.6 eV), with a shoulder peak at around 18 eV. Mapped intensities of these two peaks are presented in Figure 9C. All of the columnar features are highlighted in the map for the 25708



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Figure 10. (A) Microstructure of sample 2 under STEM-HAADF imaging conditions. Spectrum imaging results from the red frame in panel A are presented in panels B and C. (B) Three low-loss EEL spectra from three different locations on the specimen (marked on the map on the right side). (C) HAADF signal and map from the energy ranges corresponding to Mg2FeH6 (17.1−19.7 eV).

Mg2FeH6 plasmon. Also, an area of MgH2 can be located in the map constructed form the 14 eV peak. It should be noted that the Mg2FeH6 columns were mostly capped with Fe (the bright region at the top of the columns). Another example of the initial columnar growth mode of the Mg2FeH6 phase is presented in Figure 10. The STEM-HAADF image in Figure 10A shows that this column was around 110 nm long and capped with iron. The low-loss spectra along this column consistently exhibited the plasmon peak at 18.0 eV, as shown for the three cases depicted in Figure 10B. The intensity map of this peak (Figure 10C) highlights the entire column. No MgH2 phase was present in this location. Even though the columnar features of the Mg2FeH6 phase were easy to locate in sample 2, the majority of this sample is better represented by the results shown in Figure 6. The STEMHAADF signal of a similar area can be found in Figure 11. Note that, even in this area, we could still find the columnar features (marked by an arrow in Figure 11A). Figure 11B is a highermagnification view of the marked box in Figure 11A. From the contrast in the HAADF image, one can observe that Fe particles were dispersed over the larger particle, acting as a substrate. STEM/EELS data from this area are presented in Figure 11C. The low-loss EEL spectra from the matrix show solely the MgH2 volume plasmon peak at around 14 eV. An example spectrum is plotted in the accompanying graph, with the beam location marked on the HAADF image. The intensity map of this peak, on the right, shows that the entire particle, except for the regions with relatively thick obstructing Fe globules that appear dark, is brightly highlighted as single-phase MgH2. 3.2.3. Sample 3: Absorbed for 16 h at 673 K [P(H2) = 40 bar]. This powder sample was maintained in the hydrogen absorption step for a long time. According to the XRD results (see Table 1), Mg2FeH6 was the dominant phase present. Figure 12 presents a representative microstructure of this sample. A selected-area diffraction pattern of this region is shown in Figure 12C, superimposed with the simulated ring pattern for the ternary

Figure 11. Region in sample 2 with no Mg2FeH6 growth (except for the location indicated by the arrow). (A) STEM-HAADF micrograph. (B) Higher-magnification view of the red box shown in panel A. (C) Spectrum imaging data. The graph shows the low-loss EELS of the designated location within the particle. The map corresponding to the MgH2 plasmon (selected energy region of 13.0−15.5 eV) on the right shows that the particle is entirely β-MgH2. 25709



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through nucleation and growth of this phase with a columnar morphology. A schematic representation of the early stages of Mg2FeH6 formation and growth during thermal absorption is presented in Figure 15. Starting with the as-milled Mg2FeH6 and after complete desorption, the existing phases would be metallic Mg with fine metallic Fe particles (not shown in the schematic). With the start of the absorption cycle, because of the catalytic activity of iron for the dissociation of gaseous hydrogen,28 MgH2 would form with enhanced kinetics (first regime in the absorption plot in Figure 1A). At this stage, β-MgH2 would be present with metallic Fe particles (see Figure 15A and also Figures 6 and 11). Mg2FeH6 would then nucleate at the contact point between the magnesium hydride and metallic iron and grow with a columnar morphology (see Figures 15B, 7B, and 8B). As the columns of the ternary hydride phase grow, MgH2 and Fe particles would gradually be consumed and, if present at the correct stoichiometry, would entirely disappear. Because sample 2 in our study was quenched at early stages of Mg2FeH6 growth, we observed only a few columns with their iron caps fully consumed. The tapered geometry observed for some of the columns also indicates the shrinkage of the Fe particle on top as the reaction proceeded (see, for example, Figure 8). In practice, a fraction of the Fe particles lose contact with the MgH2 grains and do not participate in Mg2FeH6 formation (Figure 15C and Table 1). The same holds for the excess MgH2 that is present. In our experiments, we focused only on the initial formation of the Mg2FeH6 phase, and what we have depicted in the schematic in Figure 15C is a projection of the observed trend to the final stages of the reaction. This columnar morphology of the ternary hydride phase is present only immediately after formation. If the powder specimen were maintained at high temperature for a long time (as was the case for sample 3 in the present study), this morphology, because of the thermodynamic drive to reduce the surface energy, would agglomerate into a more equiaxed morphology. The columnar morphology reported herein could possibly be the origin of the so-called “vermicular” structures observed in the scanning electron microscopy of the Mg2FeH6 system.29 This microstructure was first reported by Bogdanovic et al.,30 who presented TEM micrographs showing that the ternary hydride phase was growing directly from Fe particles, with no metallic Mg or MgH2 present. They concluded that magnesium was participating in the reaction as a vapor phase. In contrast, our present study clearly shows that Mg2FeH6 forms as columns between MgH2 and metallic Fe particles. Because the TEM

hydride phase. The majority of the diffracted signal can be correlated with the Mg2FeH6 phase, with faint reflections of Fe and MgO also present. A dark-field image of this area is shown in Figure 12B, with the position of the objective lens aperture (OLA) with respect to the tilted SAD pattern depicted in the inset. Because of the close proximity of the diffraction rings, this DF image includes reflections from all of the phases present. The grain structure, nonetheless, shows an equiaxed morphology across the particle. Figure 13A shows an STEM-HAADF image of an area in sample 3. Figure 13B shows the spectrum imaging results from the marked region in Figure 13A. According to the location of the plasmon peak across the particle, two phases were present, βMgH2 and Mg2FeH6 (with the peak at around 18 eV). The intensity maps of the corresponding phases are presented at the top, with two examples of the low-loss spectra shown in the bottom graph. The majority of this area represents the ternary hydride phase. The distinctive columnar morphology that was observed for the early stages of Mg2FeH6 formation was not prevalent in sample 3. The STEM-HAADF image in Figure 14A presents an area that contains a column-like feature. The STEM/EELS data in Figure 14B, indeed, show that the low-loss EEL spectrum of the column-like section has a peak at around 18 eV (see the map marked as Mg2FeH6 I and plot 2 in the graph at the bottom of panel B). In other areas of this particle, however, we found plasmon peaks located at lower energies. The map marked as Mg2FeH6 II in Figure 14B shows the signal intensity of another energy range. Spectrum 3 in the graph demonstrates a plasmon peak closer to 16.7 eV. A peak at energy levels lower than 18 eV was observed in other areas of sample 3 as well. It appears that the plasmon peak position of the ternary hydride phase, after a long annealing step, shifts toward its value in the as-milled state, namely, 16.5 eV (see section 3.2.1). Because of peak overlap of this shifting signal, it is hard to construct intensity maps of every observed Ep value. The peak around 14.5 eV, corresponding to MgH2, can be found and mapped as shown in Figure 14B. In the Discussion section, we will return to this phenomenon of the shifting plasmon peak for the ternary complex hydride phase.

4. DISCUSSION 4.1. Growth and Morphology of Mg2FeH6 during Hydrogen Absorption. Our microscopy observations and spectroscopy measurements demonstrate that the thermal formation of the ternary complex hydride Mg2FeH6 proceeds

Figure 12. Representative microstructure of sample 3: fully hydrogenated powder sample, as observed by conventional TEM: (A) bright-field, (B) darkfield, with the location of the objective lens aperture (OLA) with respect to the tilted diffraction pattern shown in the corner. (C) SAD with the simulated ring pattern for Mg2FeH6. 25710



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Figure 13. (A) STEM-HAADF image showing an area in sample 3. (B) Spectrum imaging results from this region. The three images at the top are from left to right: HAADF signal, map of MgH2 plasmon (13.6−15.4 eV), and map of Mg2FeH6 plasmon (17.3−19.1 eV). The graph at the bottom presents two low-loss EEL spectra from the two locations on the particle (marked with cross-hairs on the maps above).

Figure 14. (A) STEM-HAADF image of an area in sample 3. (B) Spectrum imaging data from the area marked in panel A. Three examples of the lowloss EEL spectra are shown in the graph at the bottom. The beam location in each case is designated by cross-hair with the same color (also numbered). The three maps on the right side of the HAADF image are for MgH2 (13.9−15.7 eV), Mg2FeH6 plasmon peak I (16.9−18.8 eV) and Mg2FeH6 plasmon peak II (17.3−19.0 eV).

experiments in the study by Bogdanovic et al. were performed at room temperature, there is a possibility that the MgH2 phase was promptly decomposed upon electron exposure (see the related discussion in Experimental Methods section). The ternary

hydride phase, Mg2FeH6, is far more stable under an electron beam even at room temperature. Figure 15D shows a close depiction of a growing Mg 2FeH6 column. As evidenced by our microscopy observations 25711



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and Fe (or, more accurately, their cations, Mg2+ and Fe2+; see Figure 15D). Because this latter diffusion process is slow, hydrogen cycling of the Mg2FeH6 system generally suffers from poor kinetics. In systems based on Mg and Fe that are tailored mainly toward fast sorption kinetics, the formation of the ternary complex hydride would rather to be avoided.32,33 An important question to consider here is the location of the reaction front in the growing Mg2FeH6 column (Figure 15D). In other words, which of the two metallic cations is the diffusing entity? An analogous question for the case of magnesium hydride, that is, the diffusion characteristics of various point defects within the MgH2 structure, was addressed recently by first-principles calculations.34−36 Ab initio calculations of the Mg2FeH6 phase that have appeared in the literature so far have been focused on the electronic, structural, and thermodynamic properties of the ground state of this phase.37−39 Further calculations could provide insights regarding the diffusion dynamics of potential species within the ternary hydride phase. Our experimental results cannot provide a definitive answer to the question raised above. However, considering the crystal structure of Mg2FeH6, we found that Fe is located within an octahedral cage of six covalently bonded H atoms (dFe−H = 1.54 Å), whereas Mg is within a much more open space (dMg−H = 2.27 Å).40 Hence, considering only the structural aspects, magnesium ions seem to be more mobile than iron ions. Once Mg2+ reaches the Mg2FeH6/Fe interface, all of the reacting agents for forming the ternary hydride phase would be present, namely, Mg, Fe, and atomic H. The ternary phase would form under the Fe particle, and the particle would move by the growing phase and form the protruding structure as the Fe is consumed (see the tapered morphology in Figure 8). From this scenario, the following questions arise: Why does Mg2FeH6 not form directly from metallic Mg, and why does MgH2 form first? The initial formation of β-MgH2 might be solely due to better kinetics, compared to the formation of the ternary hydride phase, which requires the diffusion of heavier species. In fact, Riktor et al.41 reported experimental results on the formation of Mg2FeH6 from the elemental phases Mg and Fe. Based on the premise that Mg2FeH6 has a lower equilibrium pressure than MgH2, they chose a H2 pressure during absorption where MgH2 would be thermodynamically unfavorable. In that case, Mg2FeH6 formed directly from Mg and Fe. Also in that study, the authors applied Johnson−Mehl−Avrami (JMA) theory to study the kinetics of the formation of the ternary hydride phase. Their analysis indicated that the growth of Mg2FeH6 is interface-controlled. In comparison, the same analysis for the ball-milled MgH2 powder points to a diffusion-controlled growth of a fixed number of nuclei.41 The columnar morphology that we observed during the growth of the ternary hydride phase could be correlated with the interface-controlled nature of the growth. The controlling interfacial processes could be one or a combination of the following: (1) H2 adsorption and dissociation on the Fe particles and (2) atomic jumps of the Mg2+ cations from the magnesium (or MgH2) reactant phase into the product Mg2FeH6 phase. 4.2. Shift in the Low-Loss Peak for the Ternary Hydride Phase. By performing electron energy-loss spectroscopy on the as-milled powder sample (sample 1), which mainly consisted of the Mg2FeH6 phase, we acquired the characteristic low-loss plasmon peak of this phase (Ep = 16.5 eV). This peak was then used as a “fingerprint” to map the location of this phase in sample 2, which represented the initial stages of the Mg2FeH6 growth. The peak associated with the plasmon excitation of the freshly formed ternary hydride phase showed a shift to higher energies

Figure 15. Schematic representation of the sequences of the initial steps in Mg2FeH6 formation: (A) MgH2 forms with enhanced kinetics due to the catalytic activity of Fe particles. (B) Mg2FeH6 grains nucleate at the contact point of Fe and MgH2 and grow with a columnar morphology. (C) Both MgH2 and Fe are gradually consumed by the growing Mg2FeH6 phase. (D) Close view of the red frame shown in panel B. Refer to text for details.

(Figures 7−10), in the initial stages of growth, the columns are always capped with Fe particles. This points to the twofold role that iron fulfills during this phase transformation, acting both as a catalyst for H2 dissociation and as a reactant for Mg2FeH6 formation. The catalytic role of the capping particle (Fe) and the resulting columnar morphology shows a close resemblance to the case of Si and Ge nanowire growth through chemical vapor deposition, using gold particles as a catalyst (see ref 31 and references cited therein). In that case, Au capping particles form eutectics with Si and Ge and mediate the integration of the incoming flux of atoms into the growing nanowires. Given the observed columnar morphology, as a consequence of directional flow of atomic hydrogen from the catalytic Fe particles, we can confirm that gaseous hydrogen actively participates in Mg2FeH6 formation. One of the proposed pathways for the formation of the ternary hydride phase is the reaction17 3MgH2 + Fe → Mg 2FeH6 + Mg

(3)

If this reaction were the formation route for the ternary complex hydride, the newly formed Mg2FeH6 would have nucleated at the contact point of the two reactants, namely, MgH2 and Fe, and grown into the initial MgH2 particles. There would have been no reason to develop a protruding morphology as we observed here. Also, we did not detect any trace of metallic Mg (Ep = 10 eV) close to the growth sites of the Mg2FeH6 phase (Figures 7−10). We have to point out here that the preceding statement is true only for the thermal hydrogenation of this system. We also observed the formation of Mg2FeH6 during ball milling under an Ar atmosphere. We cannot rule out the possibility that the hydrogen-free reaction, that is, reaction 3, is in operation during milling. The formation of Mg2FeH6 requires the long-distance diffusion of not only atomic H but also the metallic species Mg 25712



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hydrogenation. X-ray diffraction measurements correlated the first stage with the formation of β-MgH2 and the secondary sluggish regime with Mg2FeH6 formation. Three samples were chosen for microscopy and spectroscopy characterization: (1) asmilled powder sample, (2) sample quenched immediately after the first fast stage of hydrogen absorption, and (3) fully hydrogenated sample (kept for 16 h in the absorption step). The as-milled powder consisted mostly of Mg2FeH6 and β-MgH2 phases. The low-loss EELS peak of the ternary hydride phase was located at 16.5 eV. This allowed us to map the various phases present in this sample, using the spectrum imaging (STEM/EELS) technique. The microstructure was shown to be a co-continuous mixture of Mg and the two present hydride phases, Mg2FeH6 and MgH2. Sample 2 captured the early stages of Mg2FeH6 formation during thermal absorption. The microstructure of this sample was mostly larger particles of MgH2, covered with smaller globules of elemental Fe. The ternary hydride phase, Mg2FeH6, was found to be forming with a columnar morphology, emerging from the MgH2 particles and capped with Fe. The low-loss EELS peak of the Mg2FeH6 phase in this columnar morphology was consistently at 18.0 eV, showing a shift compared to our measurements in the as-milled state of this phase. The iron particles in this context performed a twofold role of first acting as catalytically active centers for generating atomic hydrogen and second acting as one of the reactants to ultimately form Mg2FeH6. The microstructure at the fully hydrogenated stage (sample 3) mostly consisted of Mg2FeH6 and MgH2, with an equiaxed grain structure. Longterm annealing during absorption caused the Mg2FeH6 columns to coalesce into semispherical particles to reduce the overall surface area. The plasmon peak position corresponding to the Mg2FeH6 phase at various locations in sample 3 differed. The energy loss values of the peak maximum for this phase in sample 3 were between 18 and 16.5 eV. The observed shift in the volume plasmon energy for the Mg2FeH6 phase is believed to be due to potential modifications of the atomic positions within this structure in the early stages of phase transformation. The columnar growth of the ternary hydride phase provides a complementary access to the kinetics of growth, in addition to kinetics measurements based on the instantaneous hydrogen pressure. This feature could be exploited to reach a better understanding of this phase transformation. On the other hand, the optimum conditions for the initiation of growth are still unknown. The particle size of Fe, the orientation of the Mg or MgH2 grains, and the nature of the interface at Fe/Mg2FeH6 and (Mg or MgH2)/Mg2FeH6 interphase boundaries could each play a role. By conducting experiments with a more controlled geometry, it would be possible to gain access to some of this information.

(Ep = 18.0 eV). This was measured in numerous locations in the columnar Mg2FeH6 phase and proved to be statistically valid. Once the sample was fully hydrogenated (sample 3), the Ep values that we measured on the Mg2FeH6 phase showed a scatter in peak maximum value, that is, Ep was not consistent in various locations and ranged from 18.0 to 16.5 eV. In all our measurements, no significant shift was observed for the volume plasmon peak values of β-MgH2 (∼14.0 eV) or Mg (∼10.7 eV). As pointed out for eq 2, as a first approximation, the energy loss value for the volume plasmon excitation can be directly correlated with the square root of the volume density of the valence electrons within the unit cell. This estimation works quite well for “free-electron” solids (i.e., metals). A change in the valence electron density can then be detected by a shift in the plasmon energy. This phenomenon has been successfully used in metallic alloys to measure the composition, by calibrating the plasmon peak location (see, for example, ref 42). The case for semiconducting materials is far more complicated. For an accurate description of the low-loss peak for semiconducting solids, one needs to take into account other potential factors, such as local field and excitonic effects.43 Mg2FeH6 is a semiconductor with a band gap of 1.74 eV.40 So, as mentioned in the preceding paragraph, the shift in the low-loss peak can be difficult to interpret. Two plausible scenarios could potentially lead to the observed shift in the plasmon peak of Mg2FeH6. First, it is possible that, during the initial stages of growth, the arrangement of the hydrogen atoms is slightly different from the stable arrangement in the as-milled or fully annealed state. Subtle modifications in hydrogen crystallographic location cannot be detected by XRD or electron diffraction. As was shown to be the case for Mg2NiH4,44 an isostructure of Mg2FeH6, different allotropic varieties of this phase can be stable at different temperatures and hydrogen pressures. It is wellknown that, in simpler hydride materials, such as zirconium hydride,45 a shift in the plasmon peak location can be detected by a change in the hydrogen content (also see page 309 in ref 21). We are planning to conduct an in situ neutron diffraction study, monitoring the initial stages of Mg2FeH6 phase formation, to conclusively detect possible modifications of the H arrangement within the crystal structure. The second possible reason behind the shift in the plasmon peak could be a modification of the crystal structure that also involves the heavier elements, that is, Mg or Fe atoms. It is possible that, because of the limitations in the kinetics of the reaction, given that it entails long-range diffusion of Mg cations, a metastable variation of the Mg2FeH6 phase, with an offstoichiometric ratio of Mg to Fe, forms first. This metastable phase could then transform to the equilibrium state after longterm annealing under hydrogen. This would explain the gradual shift of the plasmon peak toward the value originally observed in the as-milled state. Performing a more quantitative diffraction analysis could validate or disprove this hypothesis.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Tel.: +1 905 525 9140 Ext. 24862. Fax: +1 905 521 2773.

5. CONCLUSIONS We have studied the thermal formation of the ternary complex hydride, Mg2FeH6, utilizing (scanning) transmission electron microcopy and electron energy-loss spectroscopy with a cryogenically cooled specimen stage. Mg2FeH6 was formed during the high-energy ball milling of the initial powder mixture of MgH2 and Fe (4:1 molar ratio). After full desorption of the asmilled powder, the hydrogen uptake kinetics (at 673 K and 40 bar H2) demonstrated an initial fast absorption up to 5 wt % within the first few minutes followed by a slow regime of

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank Dr. Vincent Mauchamp for insightful comments on the low-loss EELS results of this work. One of the authors (M.D.) thanks Mr. Andy Duft and Dr. Andreas Korinek for their technical assistance. A.A.C.A. gratefully thanks the Brazilian 25713



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(28) Liang, G.; Huot, J.; Boily, S.; van Neste, A.; Schulz, R. J. Alloys Compd. 1999, 292, 247−252. (29) Polanski, M.; Plocinski, T.; Kunce, I.; Bystrzycki, J. Int. J. Hydrogen Energy 2010, 35 (3), 1257−1266. (30) Bogdanovic, B.; Reiser, A.; Schlichte, K.; Spliethoff, B.; Tesche, B. J. Alloys Compd. 2002, 345, 77−89. (31) Gai, P. L.; Sharma, R.; Moss, F. M. MRS Bull. 2008, 33 (2), 107− 114. (32) Zahiri, B.; Harrower, C. T.; Amirkhiz, B. S.; Mitlin, D. Appl. Phys. Lett. 2009, 95 (10), 103114. (33) Amirkhiz, B. S.; Zahiri, B.; Kalisvaart, W. P.; Mitlin, D. Int. J. Hydrogen Energy 2011, 36 (11), 6711−6722. (34) Du, A. J.; Smith, S. C.; Lu, G. Q. J. Phys. Chem. C 2007, 111 (23), 8360−8365. (35) Park, M. S.; Van de Walle, C. G. Phys. Rev. B 2009, 80, 064102. (36) Tao, S. X.; Kalisvaart, W. P.; Danaie, M.; Mitlin, D.; Notten, P. H. L.; van Santen, R. A.; Jansen, A. P. J. Int. J. Hydrogen Energy 2011, 36 (18), 11802−11809. (37) Halilov, S. V.; Singh, D. J.; Gupta, M.; Gupta, R. Phys. Rev. B 2004, 70, 195117. (38) Zhou, H. L.; Yu, Y.; Zhang, H. F.; Gao, T. Eur. Phys. J. B 2011, 79, 283−288. (39) Zareii, S. M.; Sarhaddi, R. Phys. Scr. 2012, 86, 015701. (40) Orgaz, E.; Gupta, M. J. Phys.: Condens. Matter 1993, 5 (36), 6697− 6718. (41) Riktor, M. D.; Deledda, S.; Herrich, M.; Gutfleisch, O.; Fjellvag, H.; Hauback, B. C. Mater. Sci. Eng. B 2009, 158, 19−25. (42) Duly, D.; Cheynet, M. C.; Brechet, Y. Acta Metall. Mater. 1994, 42 (11), 3843−3854. (43) Olevano, V.; Reining, L. Phys. Rev. Lett. 2001, 86 (26), 5962− 5965. (44) Darriet, B.; Soubeyroux, J. L.; Pezat, M.; Fruchart, D. J. LessCommon Met. 1984, 103 (1), 153−162. (45) Woo, O. T.; Carpenter, G. J. C. Scripta Metall. 1986, 20 (3), 423− 426.

agency CNPq for a doctoral fellowship. This research was supported by the NSERC Hydrogen Canada (H2CAN) Strategic Research Network. The microscopy work was carried out at the Canadian Centre for Electron Microscopy, a facility supported by NSERC and McMaster University.

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Formation of the Ternary Complex Hydride Mg2FeH6 from Magnesium

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