Destabilization of Mg Hydride by Self-Organized Nanoclusters in the

Positron Doppler broadening depth profiling demonstrates that after hydrogenation, nanometer-sized MgH2 clusters are formed which are coherently embed...
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Destabilization of Mg Hydride by Self-Organized Nanoclusters in the Immiscible Mg−Ti System Kohta Asano,*,†,‡ Ruud J. Westerwaal,‡ Anca Anastasopol,§ Lennard P. A. Mooij,‡,∥ Christiaan Boelsma,‡ Peter Ngene,‡ Herman Schreuders,‡ Stephan W. H. Eijt,§ and Bernard Dam‡ †

Energy Technology Research Institute, National Institute of Advanced Industrial Science and Technology (AIST), AIST Central-5, 1-1-1 Higashi, Tsukuba, Ibaraki 305-8565, Japan ‡ Department of Chemical Engineering, Faculty of Applied Sciences, Delft University of Technology, Julianalaan 136, NL-2628 BL Delft, The Netherlands § Department of Radiation Science and Technology, Faculty of Applied Sciences, Delft University of Technology, Mekelweg 15, NL-2629 JB Delft, The Netherlands ∥ Department of Physics and Astronomy, Uppsala University, Box 516, SE-75120 Uppsala, Sweden ABSTRACT: Mg is an attractive hydrogen storage material not only because of its high gravimetric and volumetric hydrogen capacities but also because of it low material costs. However, the hydride of MgH2 is too stable to release hydrogen under moderate conditions. We demonstrate that the formation of nanometer-sized clusters of Mg reduces the stability of MgH2 by the interface energy effect in the immiscible Mg−Ti system. Ti-rich MgxTi1−x (x < 0.5) thin films deposited by magnetron sputtering have a hexagonal close packed (HCP) structure, which forms a face-centered cubic (FCC) hydride phase upon hydrogenation. Positron Doppler broadening depth profiling demonstrates that after hydrogenation, nanometer-sized MgH2 clusters are formed which are coherently embedded in an FCC TiH2 matrix. The P (pressure)−T (optical transmission) isotherms measured by hydrogenography show that these MgH2 clusters are destabilized. This indicates that the formation of nanometer-sized Mg allows for the development of a lightweight and cheap hydrogen storage material with a lower desorption temperature.

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high.6 This value is around 2−5 times those in hydrides of LaNi57 and V.8 A possible way to destabilize MgH2 and enhance the reaction kinetics is to reduce the size of Mg. It has been calculated that MgH2 is destabilized by decreasing the size on a nanometer scale.9−11 The destabilization is remarkable for a MgH2 cluster with a diameter below 1.3 nm, corresponding to 19 Mg atoms or less.11 On the basis of these calculations, the dehydrogenation temperature for a MgH2 cluster of 0.9 nm would be more than 100 K lower than for bulk MgH2 under a given hydrogen pressure. In addition, the size reduction shortens the diffusion distance of hydrogen in Mg and MgH2 and expands the specific surface area to absorb and desorb hydrogen, which as a consequence enhances the reaction kinetics. Au et al. have reported that Mg nanometer-sized particles supported on high surface area carbon aerogels have the high hydrogen sorption kinetics.12 Recent studies indicate that Mg−Ti composite nanomaterials offer viable destabilization routes via size reduction and are promising for achieving favorable hydrogen storage properties. Anastasopol et al. have prepared Mg−Ti nanometer-sized

INTRODUCTION Mg is one of the typical lightweight metals with a density of 1.74 g cm−3, which is much less than that of many transition and rare earth metals.1 Mg has a hexagonal close-packed (HCP) structure, and that forms the hydride phase MgH2 by hydrogenation. MgH2 has a rutile-type body-centered tetragonal (BCT) structure (α-MgH2) at moderate temperatures and pressures.2,3 The gravimetric hydrogen capacity is 7.6 mass %, which is higher than most metal hydrides. The volumetric hydrogen capacity of 109 g-H2 l−1 corresponds to 2.9 times the gaseous hydrogen concentration under a pressure of 70 MPa and 1.6 times the one of liquid hydrogen at 20 K.4 Mg is one of the most attractive hydrogen storage materials not only because of the high capacity but also the low material costs. However, MgH2 is too stable to desorb hydrogen at moderate temperatures and pressures and thus unsuited for practical applications. The enthalpy for hydride formation is −75 kJ mol−1-H2,2,3 which corresponds to a dehydrogenation temperature of around 550 K under a hydrogen pressure of 0.1 MPa. Furthermore, the slow reaction kinetics for hydrogenation and dehydrogenation is a disadvantage for using it as hydrogen storage material. The diffusion of hydrogen in both Mg and MgH2 has been studied,2,5,6 and the activation energy for hydrogen diffusion in MgH2 of 140 kJ mol−1 is particularly © 2015 American Chemical Society

Received: March 9, 2015 Revised: May 11, 2015 Published: May 15, 2015 12157

DOI: 10.1021/acs.jpcc.5b02275 J. Phys. Chem. C 2015, 119, 12157−12164

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The Journal of Physical Chemistry C

Mg, Ti and hydrogen atoms within the FCC hydride phase could not be clarified. We suggested that nanometer-sized Mg clusters embedded in a Ti matrix are formed. The MgH2 clusters resulting from subsequent hydrogenation could then be destabilized by the interface energy effect. In the present work, we verified this concept by synthesizing Ti-rich MgxTi1−x thin films. The structure is investigated on the nanometer scale using the positron Doppler broadening depth profiling method. The stability of the hydrides is evaluated by an optical technique called “Hydrogenography” which allows to measure the P (pressure)-T (optical transmission) isotherms.23,37,38 From the isotherms we found evidence for the formation of destabilized MgH2, while the structural data confirmed the presence of MgH2 entities within a Ti-rich MgxTi1−x−H matrix.

composites by spark discharge generation, and indeed, a destabilization of MgH2 has been found.13 Thin-film multilayer stacks prepared by magnetron sputtering is a more practical approach to systematically study nanometer-sized Mg.14−16 Baldi et al. have grown Mg thin films capped with various transition metals.16 They showed that the equilibrium hydrogen pressure for hydrogenation of Mg at a given temperature depends type of cap layer. Ni and Pd, which are miscible with Mg, effectively increase the equilibrium pressure. On the other hand, Ti, V, and Nb, which are immiscible with Mg, did not show a significant change in the hydrogenation properties of the Mg. The difference was attributed to the elastic clamping exerted by alloying elements.16,17 The thermodynamic change was found for dehydrogenation of MgH2 also. Mooij et al. have shown that the equilibrium hydrogen pressure for hydrogenation and dehydrogenation increases with reduction of the thickness of the Mg layers from 10 nm down to 1.5 nm in the immiscible Mg−Ti multilayer stack.18 They clearly demonstrated that MgH2 is destabilized by interface energy between Mg and Ti layers. As summarized above, the size reduction of Mg in fact tailors the thermodynamics of the Mg−H2 system using the interface energy effect. In the present work, we further explore the immiscible Mg−Ti system to destabilize the MgH2 phase. Instead of looking at artificial interfaces we now look at the interfaces formed by self-organization during hydrogenation. Jensen et al. have calculated the mixing enthalpies for MgxTi1−x in the range of 0.0156 ≤ x ≤ 0.9844 both for quasirandom and nanocluster models by density functional theory (DFT) calculations.19−21 They find that Ti segregates into nanometer-sized clusters in Mg-rich compositions while nanometer-sized Mg segregates in the Ti-rich ones. This has been verified for Mg-rich MgxTi1−x thin films prepared by magnetron cosputtering.21−25 In the as-deposited state, these alloys have an HCP structure, similar to the pure Mg and Ti metals. The crystal structure of MgxTi1−x hydrides in the range of x ≥ 0.90 resembles that of the rutile-type MgH2 BCT phase while for x ≤ 0.87 a fluorite-type TiH2 face centered cubic (FCC) phase is found.22 These two hydride phases coexist in the range of 0.87 ≤ x ≤ 0.90. Using extended X-ray absorption fine structure (EXAFS) spectroscopy, Baldi et al. have shown that for MgxTi1−x−H in the range of 0.53 ≤ x ≤ 0.90 the local chemical segregation consists of a Mg matrix and nanometersized Ti domains within a large coherent FCC grain.26 Srinivasan et al. have measured nanometer-sized embedded Ti domains in Mg0.65Ti0.35−D from a shift of the 2H magic angle spinning nuclear magnetic resonance (MAS NMR) signal.27−29 Using the positron Doppler broadening depth profiling method, Leegwater et al. have found evidence that TiH 2 domains are embedded in a MgH 2 matrix for Mg0.70Ti0.30−H and Mg0.90Ti0.10−H.30 MgxTi1−x alloys have also been prepared by ball milling of Mg and Ti powders in the range of 0.25 ≤ x ≤ 0.80.31−36 These samples have HCP, FCC, and body-centered cubic (BCC) structures depending on the compositions and the milling conditions. The hydride phase with an FCC structure is in this case found only in the Ti-rich powders32,34,36 in contrast to the FCC and BCT hydride phases found in Mg-rich thin films.22,24 Asano et al. have synthesized Mg0.25Ti0.75H1.62 from a BCC Mg0.25Ti0.75 alloy prepared by ball milling.36 This hydride consists of a single FCC phase as shown by the atomic pair distribution function (PDF) analysis on synchrotron X-ray total scattering data and 1H NMR. However, the local distribution of

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EXPERIMENTAL SECTION

Thin Film Preparation. MgxTi1−x thin films were deposited at room temperature on quartz substrates under an argon pressure (of purity 6 N) of 0.3 Pa in an ultrahigh vacuum (UHV) RF/DC magnetron cosputtering system (base pressure 10−7 Pa) with off-centered sources. The thin films for X-ray diffraction (XRD) measurements had a gradient composition of Mg and Ti, determined by controlling their sputtering rates. A gradient MgxTi1−x (0.13 ≤ x ≤ 0.59) thin film with a thickness of 500 nm was deposited on seven individual 10 × 10 × 0.5 mm substrates. The thin films were covered by uniform 5 nm layer of Pd to prevent oxidization and promote hydrogenation. We also deposited a pure Ti thin film covered by Pd for reference. The thin films for positron Doppler broadening of annihilation radiation depth profiling were a homogeneous Mg0.25Ti0.75 thin film with a thickness of 200 nm covered by uniform 5 nm layer of Pd, which were deposited on rotating 10 × 10 × 0.5 mm substrates. The thin films for the hydrogenography were deposited on 70 × 5 × 0.5 mm substrates. A Ti uniform layer of 3 nm was deposited first by rotating the substrate. After the rotation was stopped, a gradient MgxTi1−x (0.13 ≤ x ≤ 0.59) thin film with a thickness of 60 nm was deposited. Finally, the thin films were covered by 8 nm of Ti and 10 nm of Pd. Structural Analysis. The XRD was measured using a Bruker D8 Advance diffractometer with Co Kα radiation with λ = 1.7890 Å. The diffraction patterns of the thin films before hydrogenation in the unloaded state were measured in air and those of the hydrides under a hydrogen atmosphere. The interplanar distances of the alloy and hydride phases were calculated from the diffraction peak positions. The positron Doppler broadening of annihilation radiation with 511 keV was measured using positrons with a tunable kinetic energy, Epositron, in the range of 0−25 keV. The parameters S (shape) and W (wing) were determined from the measured electron-positron momentum distribution. Momentum windows for S and W were |p| < 3.0 × 10−3 m0c and 8.2 × 10−3 m0c < |p| < 23.4 × 10−3 m0c, respectively. The symbols of m0 and c are the electron/positron rest mass and the light velocity, respectively. The S parameter is a measure for positron annihilation with valence electrons which provide sensitivity to the electronic structure and the presence of vacancies. The W parameter is a measure of annihilation with semicore electrons which provide chemical sensitivity to the positron trapping site. The S and W parameter depth profiles were analyzed using the VEPFIT program package.39 The detailed procedures for the 12158

DOI: 10.1021/acs.jpcc.5b02275 J. Phys. Chem. C 2015, 119, 12157−12164

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The Journal of Physical Chemistry C positron Doppler broadening depth profiling method were described in the previous report.40,41 Hydrogenation and Dehydrogenation Properties. Hydrogenation properties of the gradient thin film on 70 × 5 × 0.5 mm substrate were measured with hydrogenography.23,37,38 First, the thin film was precycled three times of hydrogen absorption and desorption at 393 K to obtain a reproducible optical response. Then so-called P−T isotherms were measured at the same temperature. The optical transmission, T, is given by the Lambert−Beer relation ln(T/T0) is proportional to CHt, where T0 is the initial transmission, CH is the hydrogen content in the film, and t is the film thickness.42

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RESULTS AND DISCUSSION Structure before and after Hydrogenation. A gradient MgxTi1−x (0.13 ≤ x ≤ 0.59) thin film covered by a uniform layer of Pd was deposited by magnetron cosputtering of Mg and Ti, as described in detail in the Experimental Section. In order to determine the crystalline structure of the gradient thin films, we first performed XRD measurements (Figure 1a). For x = 0 the diffraction peak at 2θ ≈ 44.8° corresponds to the 002 reflection of HCP Ti. A much smaller broadened peak at 2θ ≈ 47° might be related to either the 101 reflection of HCP Ti or the 111 reflection of FCC Pd. A Ti thin film could have a double growth direction especially at the lower deposition temperature.43 In the 0.16 ≤ x ≤ 0.49 range two clear peaks are observed which shift to the lower 2θ side with increasing Mg content (Figure 1a). We attribute the peaks to the 100 and 101 reflections of HCP MgxTi1−x, respectively. However, for x = 0.56 a sudden change in pattern is observed, which we interpret to be due to the presence of the 002 and 101 reflections. Subsequently, the interplanar distances, d, were calculated from the peak positions and compared to those in the range of 0.53 ≤ x ≤ 0.90 that have been reported by Baldi et al.26 (Figure 1b). For x > 0.5 and x = 0 the d002 shows an almost perfect Vegard’s law44 dependence. This indicates that as-deposited Mg-rich MgxTi1−x (x > 0.5) consists of a single HCP phase that originates from a supersaturated solid solution of randomly dispersed Mg and Ti atoms. In addition, the 101 reflection observed below x = 0.7 agrees well with a Vegard’s law dependence. Baldi et al. have suggested that this reflection is due to a satellite diffraction peak origination from the coexistence of Ti and Mg domains with coherent boundaries.26 However, the diffraction peak observed at 2θ ≈ 44.8° for x = 0.56 shifts to the higher 2θ side with decreasing Mg content (Figure 1a). Therefore, we suggest that this corresponds to the 101 reflection. For x < 0.5, the disappearance of the 002 reflection (Figure 1a) shows that the Ti-rich MgxTi1−x (x < 0.5) has a different crystal orientation than the Mg-rich samples. The d101 and d100 values reduce with decreasing Mg content and have smaller values than predicted by the Vegard’s law dependence (Figure 1b). Especially for x < 0.3, the d101 and d100 values are almost the same as for pure Ti, i.e., 0.2251 and 0.2563 nm,45 respectively. This suggests that Ti and Mg domains are locally segregated. Upon hydrogenation, for x = 0 the diffraction peak at 2θ ≈ 39.7° corresponds to the 111 reflection of the TiH2 FCC phase (Figure 2a). A much smaller peak at 2θ ≈ 45° seems to be the 111 reflection of FCC Pd−H. The 111 diffraction peak shifts to the lower 2θ side with increasing Mg content suggesting the formation of an FCC MgxTi1−x−H phase. The d values calculated from the peak position are shown in Figure 2b, together with the data on bulk hydrides of Mg0.25Ti0.75H1.62,36

Figure 1. (a) X-ray diffraction patterns of as-deposited MgxTi1−x (0.16 ≤ x ≤ 0.56) and Ti thin films capped with a Pd layer and (b) interplanar distances d (filled symbols). The d in the range of 0.53 ≤ x ≤ 0.90 (open symbols) is from ref 26. The dotted lines indicate a Vegard’s law dependence of d100, d002, and d101 calculated from HCP Mg and Ti and d111 of FCC Pd reported by ref 45. In this figure, all blue dotted lines refer to bulk samples.

MgH2,45 TiH2,46 and TiH1.56.47 The d111 obtained by the present work and on comparable thin films from Borsa et al.22 and Baldi et al.26 show a linear relationship in the range of 0 ≤ x ≤ 0.81. For x = 0, the d111 value is evaluated to be 0.2633 nm, and this is larger than d111 = 0.2575 nm of bulk TiH2.46 This demonstrates that the out-of-plane expansion by hydrogenation of the thin films is larger than that of the corresponding bulk, resulting in the larger d111 for the present MgxTi1−x−H thin films than Mg0.25Ti0.75H1.62 synthesized from ball milled powder36 (Figure 2b). Indeed, it has been reported that the out-of plane d111 for the FCC Pd−H thin film is also larger than that for bulk.48 In order to investigate the structure on a nanometer scale we performed positron Doppler broadening depth profiling of asdeposited and hydrogenated Mg0.25Ti0.75 thin films (Figure 3). The parameters S and W extracted from the Dopplerbroadened 511 keV annihilation peak reflect the contributions of the valence and semicore electron orbitals at the positron annihilation site, respectively, thus providing information on the local electronic structure, composition, and phase of the 12159

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Figure 3. Representative positron depth profiles of the S and W parameters (symbols) for Mg0.25Ti0.75(−H) thin films capped with a Pd layer in the unloaded state before and after hydrogenation, together with the corresponding fit curves obtained by VEPFIT analysis (full lines).

Figure 2. (a) X-ray diffraction patterns of MgxTi1−x−H (0.16 ≤ x ≤ 0.56) and Ti−H thin films capped with a Pd layer in a hydrogen atmosphere and (b) interplanar distances d (filled red circles). The d in the range of 0.53 ≤ x ≤ 0.90 is from refs 22 and 26 (open red symbols). The d111 of FCC Mg0.25Ti0.75H1.62 is from ref 36 (open blue square). The d110 of BCT MgH2 is from ref 45 (open blue rhombus). The d111 of FCC TiH2 is from ref 46, and that of FCC TiH1.56 from ref 47 (open blue rhombuses). All blue symbols refer to bulk samples.

material. The fit curves for the positron depth profiles of the S and W parameters were obtained by VEPFIT,39 based on the presence of two homogeneous layers (Pd and Mg0.25Ti0.75(−H)) on top of the substrate. The S and W parameters obtained by positrons with Epositron = 1−5 keV are primarily due to the contribution of positron annihilation in the 200 nm Mg0.25Ti0.75(−H) layer rather than the Pd cap layer and the quartz substrate.40 Clearly, the S and W parameters of the Mg0.25Ti0.75 layer decreased and increased by hydrogenation, respectively. The corresponding S-W diagrams for the depthprofiles of both the Mg 0.25 Ti 0.75 (−H) thin films are characterized by linear relations, which indicates that the Mg0.25Ti0.75(−H) thin films consist of a homogeneous layer.40 In Figure 4, the characteristic S−W points of the metal and hydride Mg0.25Ti0.75(−H) layers as deduced from VEPFIT analysis are plotted together with previously reported data on MgxTi1−x(−H) (x = 0.70, 0.90),30,40 Mg(−H),49 and Ti(−H).30,40 For as-deposited Mg0.25Ti0.75, the S−W point is much closer to Mg than to Ti in spite of the Ti-rich composition. This clearly points to the presence of Mg domains

Figure 4. Characteristic S-W points for Mg0.25Ti0.75(−H) (filled circles). Data of MgxTi1−x(−H) (x = 0.70, 0.90) and Ti(−H) is from ref 30, 40. and that of Mg(−H) from ref 49. (open symbols). In this figure all blue symbols indicate metallic thin films before hydrogenation, all red symbols indicate hydrogenated thin films and a black open square refers to bulk Ti.

which are chemically segregated from Ti domains. Namely, the positron affinity of Mg is 2.1 eV lower than that of Ti;30 12160

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Figure 5. (a) Hydrogenography image of the MgxTi1−x−H2 (0.13 ≤ x ≤ 0.59) system at 393 K where the transition of Mg to MgH2 as a function of Ti content is evidenced by an change in transmission and (b) selected P−T isotherms of the MgxTi1−x−H2 (x = 0.16, 0.29, 0.36, 0.43, 0.49, 0.53, 0.59) systems. The dotted lines indicate the hydrogenation and dehydrogenation pressures for 60 nm of Mg.

and Pd formed the hydride and hydrogen solid solution phases, respectively. Those phases do not change in the present pressure range of 1−105 Pa at 393 K.51 Therefore, the optical change in brightness between black and white is due to the hydrogenation and dehydrogenation mainly of the Mg domains, corresponding to the transformation between the metallic Mg and transparent MgH2 phases (Figure 5a). The optical change is increased with Mg content and a hysteresis is observed: the equilibrium hydrogen pressure for hydrogenation is higher than that for dehydrogenation. The P−T isotherms of the MgxTi1−x−H2 (x = 0.16, 0.29, 0.36, 0.43, 0.49, 0.53, 0.59) systems are constructed from the hydrogenography image (Figure 5b). The absorption and desorption isotherms showed sloping plateau regions which are attributed to the coexistence of the metal and hydride phases of the Mg domains in the MgxTi1−x−H thin film. For x = 0.53, which consists of almost equal parts of Mg and Ti, the equilibrium pressure of the plateau region, Peq, is clearly the lowest in the absorption process. The plateau region exhibits a larger slope in the desorption process. The maximum pressure in the desorption plateau region for x = 0.53 is Peq ≈ 200 Pa, which is slightly lower than Peq ≈ 300 Pa for x = 0.49 and 0.59. Gremaud et al. have reported for the Mg-rich compositions that the plateau region shifts to the lower pressure side and shortens with decreasing Mg content in the range of 0.61 ≤ x ≤ 0.85.25 This is due to the distribution of the local chemical composition of Mg and Ti in FCC MgxTi1−x−H, resulting in the various tetrahedral coordination MgjTi4−j (0 ≤ j ≤ 4) of hydrogen. Hydrogen occupying the tetrahedral sites is stabilized by the Ti coordination. The Peq for x = 0.53 is slightly lower than that for x = 0.59 because of the thermodynamic effect of the tetrahedral sites on hydrogen. On the other hand, for the Ti-rich compositions the Peq in the absorption and desorption processes increase with increasing Ti content from x = 0.53 to 0.43. The Peq remains unchanged below x = 0.36. This indicates that the hydride of Mg domains is destabilized with the Ti content in the Ti-rich compositions, which is not in accordance with the thermodynamic effect deduced for the Mgrich compositions. The hydrogenography results suggest that the MgH2 nanodomains which we deduced to be the present

therefore, most positrons are selectively injected in the segregated Mg domains, where they remain trapped and annihilate. On the other hand, the XRD elucidates that the d101 and d100 for x < 0.3 is almost same as those of pure Ti, and any reflections of pure Mg are not observed (Figure 1a,b). These results suggest that Mg forms nanometer-sized domains embedded in a Ti matrix and that the interface between them hardly has an elastic connection, while larger Mg domains with increasing Mg content increase the d values. While the S−W point of Mg0.25Ti0.75−H after hydrogenation is still much closer to Mg(−H) rather than Ti(−H), it shifts toward the S−W point of TiH230,40 on hydrogenation. This is in a striking contrast to the S−W points of MgxTi1−x−H (x = 0.70, 0.90),30,40 which are all close to MgH249 (Figure 4). This disagreement in the S−W points between Mg0.25Ti0.75−H and MgH2 indicates that, at high Ti fraction, d-electrons of TiH2 provide a small but clearly visible contribution, leading to a higher W value. This could be caused by the presence of Ti near the MgH2/TiH2 interfaces which may contribute to the positron annihilation at sufficiently small sizes of the MgH2 clusters embedded in the TiH2 matrix. A similar phenomenon has been reported as the affinity-induced quantum-dot-like positron state which can detect nanometer-sized particles with the size of ∼1 nm embedded in the host materials.50 Alternatively, the observed (small) TiH2 d-electron contribution could also be caused by some positron annihilation in the TiH2 itself, as the TiH2 is expected to contain Ti mono vacancies acting as positron trapping sites for positrons initially implanted in the TiH2 matrix.41 It should be noted that the coherent embedding of nanometer-sized MgH2 domains in TiH2 leads to interfaces with hardly any vacancy-related defects which would act as positron trapping sites.30 Hence, Vegard’s law dependence shown in Figure 2b is due to the elastic interaction of the TiH2 matrix with nanometer-sized MgH2 domains. Hydrogenography. We obtained the hydrogenography image of the MgxTi1−x−H2 (0.13 ≤ x ≤ 0.59) system at 393 K (Figure 5a). Before the actual measurement a hydrogen absorption and desorption cycle at the same temperature were repeated three times to obtain a reproducible response. Ti 12161

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experimental values for Δγ for the hydrogen absorption and desorption processes of the Mg−Ti multilayer stack, respectively: Δγabs = 0.44 J m−2 and Δγdes = 0.17 J m−2.18 The P−T isotherms determined in our studies showed sloping plateau regions (Figure 5b). For x = 0.43, the maximum pressures in the absorption and desorption plateau regions, which corresponds to hydrogenation and dehydrogenation of the smallest Mg domains, are Pnano,abs ≈ 18000 Pa and Pnano,des ≈ 600 Pa at T = 393 K. The Pthick is assumed to be Pthick,abs = 600 Pa and Pthick,des = 100 Pa for 60 nm of Mg. The r is estimated in the absorption and desorption processes, respectively, rabs = 1.66 nm and rdes = 1.21 nm, on the assumption that the Δγ values for spherical Mg domains are equivalent to those for planar Mg layers. However, spherical Mg domains might have smaller Δγ values than planar Mg layers if the present Ti-rich MgxTi1−x−H consists of a Ti matrix and spherical Mg domains within a coherent FCC grain, as shown in Mg-rich MgxTi1−x−H.26 According to eq 6, to achieve the observed destabilization of MgH2 the r reduces with decreasing Δγ. Hence, we conclude that r of Mg domains is ∼1 nm which is larger than the minimum radius of rmin = 0.22 nm as calculated by the principle of the positron potential well. These estimations suggest that the MgH2 domains consist of clusters of on the order of ∼10 to 20 Mg atoms. It is argued that the clusters have a size distribution because sloping plateau regions were observed in the P-T isotherms (Figure 5b). Smaller MgH2 clusters are more effectively destabilized by the interface energy effect, resulting in a higher Peq, as Wagemans et al. have calculated that MgH2 would be strongly destabilized for a cluster of ∼19 Mg atoms or less.11 Equations 6 suggests that the increase in Peq is negligible above r ≈ 25 nm on the assumption that MgH2 and TiH2 domains form coherent interfaces. The Peq is increased with increasing Ti content but remains unchanged below x ≈ 0.4 (Figure 5b). The volumetric fraction of Mg, VMg, is calculated from the atomic weight and density of Mg and Ti on the assumption that Mg and Ti domains are segregated in MgxTi1−x and a volume of their boundaries is ignored (Figure 6). The VMg is increased with increasing Mg

from the XRD and positron Doppler broadening depth profiling are destabilized by an interface energy effect. Size Estimation of Mg Clusters. The positron Doppler broadening depth profiling indicates the formation of nanometer-sized MgH2 domains in Ti-rich MgxTi1−x−H. However, this technique is unsuitable to estimate the exact size of them. Nevertheless, based on the positron state confined in a spherical potential well, we can calculate the minimum size of the nanometer-sized domain in which a positron may annihilate. On the assumption that MgH2 forms a nanometer-sized spherical domain with the radius of r embedded in a TiH2 matrix, the minimum radius rmin required to have at least one positron energy level confined in the embedded MgH2 domain is expressed by eq 152 rmin =

h 4 2m+ΔE+

(1)

where h is Planck’s constant and m+ is the effective mass of positron. The positron affinities for MgH2 and TiH2 are not known. Therefore, we base our estimate on the positron affinity difference, ΔE+, between Mg and Ti, which is 2.1 eV based on ab initio calculations,49 leading to a minimum radius of rmin = 0.22 nm. Further, we may obtain estimated values for the size of the formed MgH2 clusters from the Peq determined in the hydrogenography measurements. Mooij et al. have reported that the Peq of Mg layers with various thicknesses in the Mg−Ti multilayer stack is affected by the change in interface energy, Δγ, between Mg and TiH2 upon hydrogenation, given by eq 2.18

⎛P ⎞ A ln⎜ nano ⎟ = Δγ ⎝ Pthick ⎠ RT

(2)

Here, the symbols of Pnano and Pthick are the Peq of Mg for a nanometer-sized thin film and a thick film at least above 50 nm, respectively. The symbol R is the gas constant and T the temperature. Now we assume that MgH2 forms nanometersized spherical domains with the radius of r embedded in a TiH2 matrix. The Pnano in eq 2 corresponds to the Peq of a nanometer-sized Mg domain. The interface area of A is assumed to be constant because the TiH2 matrix is not transformed during the hydrogen absorption and desorption of Mg domains. A ≈ A Mg | TiH2≈A MgH | TiH2

(3)

2

The interface energy difference is defined as Δγ = γMgH

2 | TiH 2

− γMg | TiH

(4)

2

The interface area per mole is proportional to the molar volume V of Mg with the radius of r. A 4πr 2 3 = 4 3 = V r π r 3

∴A=

3V r

Figure 6. Calculated volumetric fraction of Mg VMg in MgxTi1−x and schematic drawing of an {111} plane in FCC MgxTi1−x−H. The yellow, orange, and blue atoms indicate Mg, Ti, and H, respectively.

(5)

Equations 2 and 5 lead to an expression for r by the Pnano. r=

content but does not follow a linear relationship because of the difference in the density between Mg and Ti. The VMg reaches to around 50% at x ≈ 0.4. This indicates that Mg is no longer present in the form of embedded nanometer-sized clusters above x ≈ 0.4 and that the VMg of 50% is the upper limit to destabilize the hydride of MgH2. Finally, in the Mg-rich

3V Δγ

( )

RT ln

Pnano Pthick

(6) −5

−1

The value of V is 1.40 × 10 m mol as calculated from the atomic weight and density of Mg. Mooij et al. reported 3

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The Journal of Physical Chemistry C

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compositions Mg is a matrix and Ti forms nanometer-sized domains.26 According to the present work for Ti-rich MgxTi1−x−H thin films, Mg forms nanometer-sized clusters embedded in a TiH2 matrix, which destabilizes MgH2 by the interface energy effect. In contrast, for Mg0.25Ti0.75H1.62 powder segregation into Mg and Ti domains has not been found in studies where X-ray PDF and 1H NMR were applied.36 The FCC MgxTi1−x−H thin film was synthesized by hydrogenation of the HCP MgxTi1−x thin film and the FCC Mg0.25Ti0.75H1.62 powder by hydrogenation of the BCC Mg0.25Ti0.75 powder. These results suggest that Mg and Ti atoms rearrange by the first cycle of hydrogenation and Mg forms nanometer-sized clusters during the structural change. The clusters are stable during hydrogenation and dehydrogenation cycling at least over 10 cycles because Mg and Ti are immiscible under the equilibrium state. A structural study using neutron PDF and 2H MAS NMR for the deuterides is in progress to clarify the difference in the local structure between Ti-rich FCC MgxTi1−x−H thin films and powders. Self-organization of Mg nanometer-sized clusters by hydrogenation in the immiscible Mg−Ti system provides an attractive phenomenon to destabilize MgH2. Furthermore, our results of self-organization can most likely be extended to other systems, e.g., which have been reported as Mg−AlTi53 and Mg−Fe54 multilayer stacks with higher interface energy differences Δγ than Mg−Ti one to obtain more destabilized MgH2 for hydrogen storage.

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CONCLUSION In the present work, we successfully synthesized nanometersized MgH2 clusters in Ti-rich MgxTi1−x−H. Our results suggest that the smallest clusters consist of on the order of ∼10 to 20 Mg atoms and are coherently embedded in a TiH2 matrix. This structure synthesized during hydrogenation shows excellent cycling stability and leads to destabilized MgH2 for a Mg content below x ≈ 0.4. At x ≈ 0.4, the volumetric fraction of Mg is estimated to be around 50%, which is the upper limit for the destabilization of MgH2 using the interface energy effect. We experimentally demonstrated that the formation of Mg nanometer-sized clusters embedded in a matrix of immiscible metals is a powerful way to destabilize MgH2, resulting in a lower working temperature for hydrogen storage applications.

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AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS We thank Dr. Kouji Sakaki of the National Institute of Advanced Industrial Science and Technology (AIST) for his advice on the positron annihilation study. This work was supported by the FY 2013 Researcher Exchange Program between the Japan Society for the Promotion of Science (JSPS) and the Nederlandse Organisatie voor Wetenschappelijk Onderzoek (NWO).

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REFERENCES

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DOI: 10.1021/acs.jpcc.5b02275 J. Phys. Chem. C 2015, 119, 12157−12164

Article

The Journal of Physical Chemistry C

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DOI: 10.1021/acs.jpcc.5b02275 J. Phys. Chem. C 2015, 119, 12157−12164