Degradation Mechanism against Hydrogenation Cycles in Mg2

Article pubs.acs.org/JPCC

Degradation Mechanism against Hydrogenation Cycles in Mg2−xPrxNi4 (x = 0.6 and 1.0)

K. Sakaki,*,† N. Terashita,‡ H. Kim,† E. H. Majzoub,§ A. Machida,∥ T. Watanuki,∥ S. Tsunokake,‡ Y. Nakamura,† and E. Akiba†,⊥ †

National Institute of Advanced Industrial Science and Technology, AIST Central-5, 1-1-1 Higashi, Tsukuba, Ibaraki 305-0035, Japan Japan Metals & Chemicals Company, Ltd., 232 Oguni, Nishiokitama, Yamagata 999-1351, Japan § Center for Nanoscience and Department of Physics and Astronomy, University of Missouri, St. Louis, Missouri 63121, United States ∥ Japan Atomic Energy Agency, 1-1-1 Kouto, Sayo-cho, Sayo-gun, Hyogo 679-5148, Japan ⊥ International Institute for Carbon-Neutral Energy Research (WPI-I2CNER) and Department of Mechanical Engineering, Faculty of Engineering, Kyushu University, 744 Motooka, Nishi-ku, Fukuoka 819-0395, Japan ‡

ABSTRACT: A long cycle life of metal hydrides is paramount for applications. We present an investigation of the degradation mechanism against hydrogenation cycles in Mg2−xPrxNi4 (x = 0.6 and 1.0). Mg1.0Pr1.0Ni4 shows significant degradation and loss of capacity after only a few cycles. In stark contrast, Mg1.4Pr0.6Ni4 did not show any reduction of hydrogen storage capacity until 100 cycles at 313 K. The peak broadening of X-ray diffraction (XRD) patterns and accumulation of lattice strain were observed concomitantly with an increase of hydrogenation cycles only in Mg1.0Pr1.0Ni4. These changes were not observed in Mg1.4Pr0.6Ni4. In pair distribution function (PDF) patterns, r-dependent peak broadening was observed and it became significant with an increasing number of cycles in Mg1.0Pr1.0Ni4, suggesting an increase of the dislocation density. Mg1.4Pr0.6Ni4 showed higher hardness and more pulverization upon hydrogenation than Mg1.0Pr1.0Ni4. These results suggest that the dominant mechanism, the stress induced by hydrogen occupation release, must change from the formation of dislocations to pulverization by increasing the Mg content of the alloy. We discuss how this leads to better cycling properties.

1. INTRODUCTION Hydrogen storage alloys absorb and desorb hydrogen reversibly under ambient conditions. Since they have advantages including high volumetric hydrogen density and fast kinetics, they have been used for nickel-metal-hydride batteries and are one of the most promising materials to supply hydrogen for fuel cell vehicles. However, durability and hydrogen densities in weight must be improved for these practical uses. In order to improve the durability, the relationship between formation of dislocations and vacancies and degradation has been investigated using X-ray diffraction, transmission electron microscopy (TEM), positron annihilation, and other characterization methods in LaNi5-based alloys.1−12 Fractional substitution of a third element, especially Al and Sn, on the Ni site in LaNi5 improved the durable reversibility13,14 and reduced the lattice strain, dislocation density, and vacancy concentration compared with binary LaNi5.4,5,7−12 Therefore, it was concluded that the accumulation of dislocations and the lattice strain degrades the durability. In addition Al substituted LaNi5 showed the finer particle size after hydrogenation and a higher hardness than binary LaNi5.11 These results suggest one possible mechanism that leads to better durability: the substitution of Al in LaNi5 increases the hardness causing the formation energy of dislocations. This shifts the dominant mechanism of releasing stress, accumulated upon hydrogenation, from the formation of © 2014 American Chemical Society

dislocations to pulverization. As a result, the introduction of dislocations is prevented, leading to good durability. On the other hand, to overcome the low hydrogen weight density, Mg containing hydrogen-absorbing alloys have been investigated because Mg is a light metal element.15−19 Chotard et al. found that Mg1.0La1.0Ni4 absorbed hydrogen up to ∼1.0 H/M (H/M: ratio of the number of hydrogen and metal atoms) with two distinct plateaus on the pressure−composition (P−C) isotherms, 20 although most reports indicate Mg1.0RE1.0Ni4 absorbed hydrogen only up to ∼0.7 H/M with one plateau.21−26 We found the dependence of the hydrogenation properties and the crystal structure of the hydride on the Mg/RE ratio in Mg2−xRExNi4. The compound with x = 1.0 forms two different hydride phases, an orthorhombic Mg1.0RE1.0Ni4H∼4 and a C15b Mg1.0RE1.0Ni4H−6; those with x < 1.0 have only C15b Mg2‑xRExNi4H∼3.6, and those with x > 1.0 transform to amorphous hydride with or without formation of orthorhombic hydride.27−30 This indicates that the metal lattice of Mg2−xRExNi4 with x < 1.0 expanded isotropically while that of Mg2‑xRExNi4 with x ≥ 1.0 expanded anisotropically. Therefore, the dislocations and the lattice strain Received: January 24, 2014 Revised: March 5, 2014 Published: March 7, 2014 6697

dx.doi.org/10.1021/jp500855d | J. Phys. Chem. C 2014, 118, 6697−6705

The Journal of Physical Chemistry C

Article

Fourier transformation of the powder diffraction data according to eq 1:

introduced by hydrogenation, and the cyclic properties, may be influenced by the Mg/RE ratio because of different behavior of the lattice expansion. In this work, the cyclic properties, lattice strain generated upon hydrogenation, the hardness, and particle size distribution after hydrogenation cycles were evaluated in Mg2−xPrxNi4 to clarify the degradation mechanism. To identify the origin of the lattice strain and dislocations, the atomic pair distribution function (PDF) technique, which gives the probability of finding atom pairs separated by distance r, was used.31 Our previous PDF experiments revealed that the dislocation density increased with hydrogenation cycles in Ti−V solid solution alloys and that the change in r-dependent peak broadening in the PDF data had close correlation with the cycling properties.32 The mechanism of the degradation will be discussed from the viewpoint of the mechanical properties.

G (r ) =

2 π

∫Q

Q max

Q [S(Q ) − 1]sin(Qr ) dQ

min

(1)

where Q is the magnitude of the momentum transfer, Qmax and Qmin are maximum and minimum values of the magnitude of the momentum transfer for Fourier transformation, and S(Q) is the total scattering structure function.31 Because of the unfavorable signal-to-noise ratio at the high-Q regions, Q[S(Q)−1] was truncated above 20 Å−1 of Qmax before the transformation. The data processing program PDFgetX234 was used to obtain the synchrotron X-ray PDFs. Real space modeling was carried out using the PDFgui program.35 The degree of the r-dependent peak broadening was evaluated using a parameter Qbroad which is used for expressing the PDF peak width as follows:36

2. EXPERIMENTAL PROCEDURE 2.1. Sample Preparation. This work used the same alloy ingots that were investigated in ref 27. Ingots of Mg, Pr, and Ni, each with purity higher than 99.9%, were used as the starting materials. Mg1.0Pr1.0Ni4 and Mg1.4Pr0.6Ni4 were prepared by high-frequency induction melting using an alumina crucible and were cast into a water-cooled board mold under a He atmosphere. The as-cast alloys were annealed at 1323 K for 10 h under an Ar atmosphere. Chemical analysis showed that the compositions of the annealed alloys were Mg1.4Pr0.6Ni4.1 and Mg1.0Pr1.0Ni4.0. The annealed alloys were crushed into powders and put into a stainless steel vessel for pressure−composition (P−C) isotherms. The vessel was evacuated by rotary pump at 423 K for more than 2 h. P−C isotherms of Mg1.0Pr1.0Ni4 and Mg1.4Pr0.6Ni4 were measured three times at 313 K, followed by cyclic measurements at 313 K. For the cycle test, 3 MPa of hydrogen gas was applied to the samples. When the absorption reaction was completed, the hydrogen pressure was around 2.7 MPa, which is higher than the absorption plateau pressure. The dehydrogenation was carried out by evacuating the vessel using a rotary pump. The reaction time for hydrogenation and dehydrogenation was 10 min. The P−C isotherms were measured at 313 K after cyclic measurements. Mg1.0Pr1.0Ni4 and Mg1.4Pr0.6Ni4 used for synchrotron X-ray experiments were repeatedly hydrogenated at 298 K up to the first, fifth, 10th, and 25th cycle. 2.2. Synchrotron X-ray Diffraction Experiment. After hydrogenation cycles, the samples were taken out of the vessel and crushed into powders with particle diameters less than 20 μm. These fine powders were loaded into a kapton capillary. The diameter of the kapton capillary was 1.0033 mm, and its wall thickness was 0.0508 mm. Synchrotron X-ray total scattering data of Mg2−xPrxNi4 (x = 0.6 and 1.0) after the hydrogenation cycles were collected at room temperature at the Japan Atomic Energy Agency (JAEA) beam time of BL22XU at SPring-8.33 During the data collection, the samples were spun. The energy and the size of the monochromatic X-ray were 69.87 keV (wavelength, 0.17748 Å) and 0.4 × 0.4 mm2, respectively. The two-dimensional (2D) imaging plate with 400 mm × 400 mm (R-AXIS V, Rigaku) was used. 2.3. Data Processing and Analysis for Synchrotron Xray Pair Distribution Function. The signal from an empty kapton capillary was subtracted from the raw data, and various other corrections were made.31 The PDF is obtained by a sine

σij = σij′ 1 −

δ1 δ − 22 + Q broad 2rij 2 rij rij

(2)

where σ and σ′ are the peak width with and without corrections, respectively, and δ1 and δ2 are parameters to correct for the effects of correlated motion. 2.4. Peak Profile Analysis. Structural refinement in reciprocal space was carried out by Rietveld analysis using the RIETAN-2000 program.37−39 Standard Si powder from NIST was used for calibration of the detector angle. A pseudo-Voigt function is used for expressing peak profiles, where full width at half-maximum (fwhm) of the Gaussian part and the Lorentzian part, HKG and HKL, are expressed as follows:40,41 Gaussian part: HK G = [8 ln 2(U tan 2 θK + V tan θK + W + P sec 2 θK )]1/2 (3)

Lorentzian part: HK L = (X + Xe cos ϕK )sec θK + (Y + Ye cos ϕK )tan θK (4)

θK is the Bragg angle. V and W in the Gaussian part are parameters depending on the instruments. X and Xe in the Lorentzian part are parameters for crystallite size effect. U and Y are isotropic strain parameters, and Ye is an anisotropic strain parameter. ϕK is the angle between the scattering vector and an anisotropic peak broadening axis. The crystallite size and lattice strain were evaluated by the following equations.40,41 crystallite size:

180Kλ πX lattice strain from Gaussian part: π [8 ln 2(U − Ui)]1/2 × (100/%) 180 lattice strain from Lorentzian part: π (Y − Yi ) × (100/%) 180 K is the Sherrer constant equal to 0.9, λ is the wavelength of the X-ray, and Ui and Yi are instrument contributions. 2.5. Evaluation of Mechanical Property and Particle Distribution. The mechanical properties were measured using 6698



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where E0 is the total energy with equilibrium lattice constants, V0 is the equilibrium volume, B0 is bulk modulus, and B′ is the pressure derivative of the bulk modulus.

a dynamic ultramicrohardness tester (DUH-211, Shimadzu). The advantage of this hardness system is that both Martens hardness (HM) and elastic properties can be evaluated simultaneously from the loading and unloading curves. The indentation load was 98 mN. HM is defined as the maximum applied load, Pmax, divided by contact area A in this system: HM = Pmax /A=Pmax /(26.43h2)

3. RESULTS 3.1. Durable Reversibility of Mg 1.0 Pr 1.0 Ni 4 and Mg1.4Pr0.6Ni4. Figure 1 shows the change of hydrogen storage

(5)

where h is the penetration depth. The particle size distribution was measured before hydrogenation and after the eighth hydrogenation cycle using a laser diffraction light scattering method (MT3300EX II, Nikkiso Co. Ltd.). The powders were distributed into water with a small amount of EMULGEN A-60 (Kao Chemicals). Before measurements, the powders were well mixed using an ultrasonic bath. 2.6. Evaluation of Elastic Properties. In order to evaluate the elastic properties of Mg2−xPrxNi4 and the hydride, firstprinciples calculations were performed using the Vienna ab initio simulation package (VASP)42,43 with the generalized gradient approximation (GGA) proposed by Perdew et al.44 For Pr the special GGA potentials in which the f electrons are kept frozen in the core were used. This is a standard model for the treatment of localized f electrons.44 Potentials based on the all-electron projector augmented wave (PAW) method were used.45,46 Full geometry optimizations, including atomic positions, cell volume, and cell shape, were all allowed to change until the maximum force dropped below 0.001 eV/Å. The self-consistent field (SCF) convergence criterion was set at 10−7 eV. The plane wave cutoff energy was chosen to be 425 eV. The k-point spacing was 0.1 Å−1. The elastic coefficients of Mg2−xPrxNi4 were obtained from a least-squares fit to stress tensors computed for a set of strained crystals from the first-principles calculations. For this calculation, three different strains (0.005, 0.01, and 0.015) were used to get more accurate results for the fitting procedure.47 From the obtained elastic coefficients, Young’ modulus, E, bulk modulus, B, and shear modulus, G, were determined using the Voigt averaging scheme. In the Voigt average, the expressions for the cubic system are given by the following equations;

Figure 1. Change of hydrogen storage capacity against cycles in Mg2−xPrxNi4.

capacity against hydrogenation cycles in Mg2−xPrxNi4. In Mg1.0Pr1.0Ni4, the hydrogen content dramatically decreased from 0.64 to 0.57 H/M during the first 15 cycles. In the subsequent cycles up to the 45th cycle, the hydrogen content was still reduced to 0.52 H/M but the reduction rate became smaller. In further cycles up to 100 cycles, the reduction of hydrogen content became insignificant with increasing hydrogenation cycles. P−C isotherms measured before and after cyclic measurement also showed significant reduction of hydrogen capacity. This means that this reduction did not come from slow kinetics; the hydrogen storage capacity itself was decreased. On the other hand, negligible reduction in the hydrogen content was observed during the hydrogenation cycles in Mg1.4Pr0.6Ni4. These results indicate that the Mg/Pr ratio affects not only the hydrogenation property and crystal structure of the hydride27−30 but also the durable reversibility. 3.2. Structural Degradation of Mg1.4Pr0.6Ni4 against Hydrogenation Cycles. Figure 2a shows the change of XRD patterns against the hydrogenation cycles of Mg1.4Pr0.6Ni4. The crystal structure was C15b with space group F4̅3m. No secondary phases were formed during the hydrogenation cycles. The lattice constant, the atomic position of Ni at the 16e site, and the thermal factor for all elements did not change during hydrogenation cycles. The lattice strain and crystallite size were evaluated by Rietveld refinement. The crystallite size did not change during hydrogenation cycles and was around 100 nm. The accumulation of the lattice strain was not observed as shown in Figure 2b. Figure 3a shows the change of PDF patterns against hydrogenation cycles of Mg1.4Pr0.6Ni4. All PDF patterns clearly show sharp peaks up to 100 Å. There was no change in PDF patterns with the number of hydrogenation cycles, and no significant peak broadening was observed over the whole r range. Even though peak broadening was invisible, the degree of broadening was evaluated using a parameter Qbroad which is defined by eq 2. As shown in Figure 3b, Qbroad did not increase during hydrogenation cycles.

B = (C11 + 2C12)/3 G = (C11 − C12 + 3C44)/5 E=

9GB 3B + G

(6)

However, the number of independent elastic constants increases from three to nine if the crystal structure changes from a cubic structure to an orthorhombic structure. Therefore, only the bulk modulus was evaluated for the hydrides of Mg2−xPrxNi4. Using Birch−Murnaghan’s equations of state (EOS), the equilibrium lattice constants, the total energies, and the bulk modulus of Mg2−xPrxNi4, Mg2−xPrxNi4H4, and Mg2−xPrxNi4H7 were evaluated from the fitting of energy− lattice constant curves as follows:48 E = E0 +

⎡⎛ V ⎞2/3 ⎤2 ⎡⎛ V ⎞2/3 ⎤3 9 9 V0B0⎢⎜ 0 ⎟ − 1⎥ + V0B0(B′ − 4)⎢⎜ 0 ⎟ − 1⎥ ⎢⎣⎝ V ⎠ ⎥⎦ ⎢⎣⎝ V ⎠ ⎥⎦ 8 16

(7) 6699



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Figure 2. Change of (a) X-ray diffraction patterns and (b) lattice strain against hydrogenation cycles in Mg1.4Pr0.6Ni4.

Figure 3. Change of (a) PDF patterns and (b) broadening parameter, Qbroad, against hydrogenation cycles in Mg1.4Pr0.6Ni4.

Figure 4. Change of (a) X-ray diffraction patterns against hydrogenation cycles in Mg1.0Pr1.0Ni4 and (b) comparison of normalized XRD patterns.

3.3. Structural Degradation of Mg1.0Pr1.0Ni4 against Hydrogenation Cycles. Figure 4a shows the change of XRD patterns for different hydrogenation cycles of Mg1.0Pr1.0Ni4. The crystal structure was C15b with space group F4̅3m. No crystalline secondary phases were formed during the hydrogenation cycles. Mg1.0Pr1.0Ni4, hydrogenated up to 100 cycles at 313 K, had almost the same lattice constant as that before

hydrogenation. It suggests that the residual hydrogen content in crystalline Mg1.0Pr1.0Ni4 did not increase with hydrogenation cycles. The atomic position of Ni at the 16e site and the thermal factor for all elements did not change during hydrogenation cycles. However, the diffraction peaks got broader with increasing hydrogenation cycles. In addition, the peak-tobackground ratio clearly became worse after 25 cycles, as shown 6700



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r region although it was not observed clearly in the lower r region below 6 Å. This suggests that the origin of the damping of the PDF patterns results from r-dependent peak broadening. Similar results were obtained in our previous work on V−Ti systems.32 In V−Ti systems, comparison between PDF patterns from experiments and calculation revealed that the accumulation of dislocations induced such r-dependent peak broadening and the broadening became significant with increasing dislocation density.32 This r-dependent peak broadening was well described with a parameter Qbroad in refinement of the PDF data (see section 2.3 for the details of Qbroad). In addition, Qbroad had good correlation with the degradation of hydrogenation capacity in V−Ti systems.32 Therefore, Qbroad was evaluated by fitting the PDF data of cycled Mg1.0Pr1.0Ni4. Qbroad increased dramatically during the first 10 cycles and then increased gradually in the subsequent 15 cycles as shown in Figure 6b. This suggests that the dislocation density increased during the first 25 cycles in Mg1.0Pr1.0Ni4. In further cycles, significant change in Qbroad was not observed suggesting that the dislocation density did not increase. However, the development of weak amorphous-like intensities was observed in the range of 2−4 Å of a difference curve obtained from PDF refinements. Therefore, these PDF results suggest that degradation in the early stage of cycles came from the accumulation of dislocations and that in the later stage it came from the partial amorphization upon hydrogenation cycles. 3.4. Mechanical Property and Particle Size Distribution. Figure 7 shows HM and the elastic modulus of Mg2−xPrxNi4 measured using a dynamic ultramicrohardness tester. HM is monotonically increasing with increasing Mg content. In particular, Mg1.4Pr0.6Ni4 had a significantly higher HM value than others. This suggests that Mg1.4Pr0.6Ni4 has a higher formation energy for dislocations. The elastic modulus also increased with increasing Mg content. This means that Mg1.4Pr0.6Ni4 showed stronger resistance for deformation in the elastic region than Mg1.0Pr1.0Ni4. Parameters relating to the elastic property of Mg 2−x Pr x Ni 4 and their hydrides, Mg8−xPrxNi16Hy (x = 3, 4, 5; y = 4, 7), were evaluated using first-principles calculations. The results are shown in Table 1. The lattice constants on Mg8−xPrxNi16 obtained by two different calculation methods were almost the same. These calculated values well agreed with the previous experimental values.27 In alloy phase, most of the calculated elastic parameters increased with increasing Mg content. This suggests that Mg-rich Mg8−xPrxNi16 shows stronger resistance for

in Figure 4b. This suggests that the crystalline Mg1.0Pr1.0Ni4 partially transformed to amorphous upon hydrogenation cycles. Such changes in XRD patterns were not observed in Mg1.4Pr0.6Ni4. The crystallite size and the lattice strain were evaluated by Rietveld refinement. The crystallite size remained around 50 nm. During the first 10 cycles, the lattice strain dramatically increased with increasing number of cycles as shown in Figure 5. This behavior is completely different from

Figure 5. Change of lattice strain against hydrogenation cycles in Mg1.0Pr1.0Ni4.

that shown in Mg1.4Pr0.6Ni4. In the subsequent cycles up to the 25th cycle the lattice strain was still increasing, but the rate of increase became smaller. Although the hydrogen capacity continued to decrease up to 100 cycles, significant increase of the lattice strain was not observed after the 25th cycle. XRD patterns showed that the fraction of the amorphous part increased after the 25th cycle. These results suggest that the accumulation of lattice strain was a dominant factor for degradation in the early stage of cycles and the partial amorphization contributed to degradation in the later stage of cycles. Figure 6a shows the change of PDF patterns against hydrogenation cycles in Mg1.0Pr1.0Ni4. Sharp peaks were clearly observed up to 100 Å before hydrogenation. However, a significant loss of structural correlation was observed with increasing r with the number of hydrogenation cycles. In addition, the PDF peak broadening became larger in the higher

Figure 6. Change of (a) PDF patterns and (b) broadening parameter, Qbroad, against hydrogenation cycles in Mg1.0Pr1.0Ni4. 6701



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Figure 7. Martens hardness and indentation elastic modulus of Mg2−xPrxNi4. Closed and open circle marks show Martens hardness and indentation elastic modulus, respectively.

Figure 8. Particle size distribution of Mg2−xPrxNi4. Circles and squares show before and after cyclic measurements, respectively.

deformation. This tendency is similar to that obtained from the above experiment. For the hydrides, the elastic modulus decreased with increasing hydrogen content. This result suggests that the alloy phase has stronger resistance for deformation than their hydride. The particle size distribution before hydrogenation and after eight hydrogenation cycles was shown in Figure 8. The mean particle sizes of Mg1.4Pr0.6Ni4 and Mg1.0Pr1.0Ni4 before hydrogenation were 350 and 340 μm, respectively, and there were no significant differences. The mean particle sizes after the hydrogenation cycles were 37.8 and 112 μm, respectively. Mg1.4Pr0.6Ni4 showed much smaller size with narrower size distribution than Mg1.0Pr1.0Ni4. No significant change in the particle size distribution was observed during further cycles in both alloys. Although Mg1.4Pr0.6Ni4H∼3.6 (13.3%) has a relatively smaller volume expansion than Mg1.0Pr1.0Ni4H∼4 (15.4%),29 Mg1.4Pr0.6Ni4 was significantly pulverized upon hydrogenation. Our results strongly suggest that the difference in pulverization is related to the difference in hardness or brittleness between the two alloys.

while Mg1.0Pr1.0Ni4 lost around 23% of the initial capacity during 100 cycles. Here we discuss the origin of degradation of the hydrogen storage capacity. In general several possible reasons for degradation include formation of structural disorder or inhomogeneity such as dislocations and phase separation. Amorphization is also one of the causes for degradation because hydrogen in an amorphous phase occupies deep trapping sites from where hydrogen is not released under ambient conditions. In fact Pr-rich Mg2−xPrxNi4 became amorphous upon hydrogenation, and most of the absorbed hydrogen atoms were not released at room temperature.27 Mg1.0Pr1.0Ni4 after 100 cycles at 313 K had almost the same lattice constant (7.1010 Å) as that before hydrogenation (7.1048 Å) even though a 23% reduction of capacity was observed. Therefore, the increase of residual hydrogen in the crystalline part after desorption is negligible for degradation in Mg1.0Pr1.0Ni4. Considering that hydrogen storage capacity decreased with three different rates (see Figure 1), the dominant factor of degradation probably varies from the initial cycles to further cycles. Referring to the changes in the PDF broadening parameter and XRD, Qbroad and the lattice strain significantly increased suggesting the formation of dislocations, but the amorphous-like background had not developed yet in the first 10 cycles. From the 10th to 25th cycle, the increasing

4. DISCUSSION 4.1. Origin of Degradation for Hydrogen Storage Capacity. The effect of Mg/Pr ratio on the degradation for hydrogen storage capacity is clear from Figure 1. Mg1.4Pr0.6Ni4 showed no significant reduction of hydrogen storage capacity,

Table 1. Calculated Lattice Constant and Elastic Properties of Mg2−xPrxNi4 and Their Hydridesa composition Mg1.75Pr0.25Ni4 Mg1.25Pr0.75Ni4 MgPrNi4 Mg0.75Pr1.25Ni4 Mg0.25Pr1.75Ni4 E0 (eV/f.u.) V0 (Å3) B′ B0 (GPa) a (Å) b (Å) c (Å)

a (Å) 6.9062 7.0578 7.1267 7.1870 7.3094 Mg5Pr3Ni16 −29.475 88.079 4.5025 113.13 7.0628 7.0628 7.0628

C11 (GPa)

C12 (GPa)

176.36 91.95 166.19 87.79 166.24 87.42 152.36 87.27 140.83 87.69 Mg5Pr3Ni16H28 −54.251 112.89 4.3739 104.99 7.6719 7.6719 7.6719

C44 (GPa)

bulk modulus (GPa)

shear modulus (GPa) 55.6 49.1 47.8 42.4 36.1

Young’s modulus (GPa)

64.51 55.76 53.37 48.93 42.41 MgPrNi4

120.1 113.9 113.7 109.0 105.4 MgPrNi4H4

MgPrNi4H7

Mg3Pr5Ni16

144.5 128.9 125.7 112.5 97.1 Mg3Pr5Ni16H28

−30.475 90.697 4.4815 111.74 7.1321 7.1321 7.1321

−44.916 103.94 4.3515 108.22 5.1004 5.5060 7.4025

−55.358 115.08 4.2817 107.22 7.7212 7.7212 7.7212

−31.248 93.003 4.5165 108.06 7.1920 7.1920 7.1920

−56.203 118.96 4.2413 104.44 7.8071 7.8071 7.8071

a

E0 is the total energy with equilibrium lattice constants, V0 is the equilibrium volume, B0 is bulk modulus, and B′ is the pressure derivative of bulk modulus. f.u. = formula unit. 6702



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tion.4,7,11,13 All of these effects are quite similar to those obtained in Mg-rich Mg2−xPrxNi4. In LaNi5-based alloys, therefore, increase of formation energies of the dislocations is likely to lead to better cyclic properties as well as Mg2−xPrxNi4. V−Ti−Cr solid solution alloys showed partly different tendency from that observed in Mg2−xPrxNi4 and LaNi5-based alloys. Accumulation of the lattice strains degrades cyclic properties,32,51−53 similarly to the above alloys. However, V− Ti−Cr alloys with low vanadium content showing poorer cyclic properties had finer particle size and larger Vickers hardness. This difference can be attributed to the difference between brittle alloys and ductile alloys. Inui et al. suggest that dislocations were set in motion to release stress generated upon hydrogenation in TiFe because of ductility.2 V−Ti−Cr alloys with a higher vanadium content show more ductile and better cyclic properties and larger particle size than those with a lower vanadium content. It is speculated that in ductile alloys such as TiFe and V−Ti−Cr higher ductility will promote the motion of dislocations and result in preventing their accumulation and pulverization. Because of ductility, therefore, V−Ti−Cr alloys showed different relationships among dislocations, hardness, and pulverization from brittle compounds such as Mg2−xPrxNi4 and LaNi5-based alloys. In common, the important factor to improve the cyclic property is to prevent accumulation of dislocations during cycling. At least in brittle intermetallic compounds such as Mg2−xPrxNi4 and LaNi5-based alloys we believe that controlling the formation energy of the dislocations is one of the most important factors to improve cyclic properties.

rate of Qbroad and the lattice strain became smaller and the amorphous-like background started to be developed. In the further cycles, Qbroad and the lattice strain did not increase. However, the development of weak amorphous-like intensities was observed in the diffraction patterns and in a difference curve obtained from PDF refinements. These results suggest that the formation of dislocations contributed to the degradation in the first 10 cycles and the amorphization mainly contributed above 40 cycles. Between these periods, the main origin of degradation probably changed from the formation of dislocations to amorphization. The possible reason for the formation of the amorphous phase in Mg1.0Pr1.0Ni4 is simply explained by the empirical rule proposed by Aoki et al.49,50 This rule indicates that hydrogeninduced amorphization occurs when the ratio of the atomic radii, RA/RB, is above 1.37 in AB2 compounds with C15 structure. Mg1.0Pr1.0Ni4 did not clearly show amorphization in the early stage of the hydrogenation cycles even though Mg1.0Pr1.0Ni4 has RA/RB = 1.372, where RA is the weighted average of RMg and RPr. However, considering that Mg0.8Pr1.2Ni4 whose composition is close to Mg1.0Pr1.0Ni4 showed the hydrogen-induced amorphization upon the first hydrogenation, it is likely that a part of Mg1.0Pr1.0Ni4 gradually became amorphous during the repeated hydrogenation cycles as the increasing background observed in Figure.4b. In conclusion, the origins of the degradation in Mg1.0Pr1.0Ni4 in the early stage and the later stages were the formation of dislocations and the formation of the amorphous phase, respectively. 4.2. Relationship between Cyclic Property and Material Property. As described above, Mg1.4Pr0.6Ni4 showed excellent cyclic properties while the hydrogen storage capacity of Mg1.0Pr1.0Ni4 was reduced by around 23% over 100 cycles. The synchrotron X-ray total scattering experiments showed that the lattice strain and the dislocations introduced upon hydrogenation were accumulated in Mg1.0Pr1.0Ni4 while those in Mg1.4Pr0.6Ni4 were not. Mg1.4Pr0.6Ni4 had the higher hardness and was significantly pulverized upon hydrogenation although Mg1.4Pr0.6Ni4H∼3.6 has relatively smaller volume expansion (13.3%) than Mg1.0Pr1.0Ni4H∼4 (15.4%).29 In addition, the calculated elastic modulus of Mg-rich Mg2−xPrxNi4 was larger than that of Mg1.0Pr1.0Ni4 suggesting that Mg-rich Mg2−xPrxNi4 showed stronger resistance for deformation than Mg1.0Pr1.0Ni4. These results can be explained by differences in the formation energy of the dislocations. If the formation energies of dislocations increase, the pulverization will become energetically more favorable than the introduction of the dislocations for reducing the stress generated upon hydrogenation. This would promote pulverization and prevent formation of the dislocations. Therefore, increase of the formation energy of dislocations would change the dominant mechanism of releasing the stress from formation of dislocations to pulverization and it leads to better cyclic properties. The fact that Mg1.4Pr0.6Ni4 has less dislocation, higher hardness, smaller particle size, larger elastic constant, and better durability than Mg1.0Pr1.0Ni4 suggests that the substitution of Mg for Pr in Mg2−xPrxNi4 probably increased the formation energies of the dislocation. The relationship between cyclic properties and material properties has also been investigated in LaNi5-based alloys and V-based alloys. In LaNi5-based alloys, Al substitution for part of Ni in LaNi5 improved the cyclic properties, reduced the lattice strain, dislocation density, and vacancy concentration, increased hardness, and enhanced pulverization upon hydrogena-

5. CONCLUSION The degradation mechanisms upon cyclic hydrogenation have been investigated in Mg2−xPrxNi4 (x = 0.6, 1.0). The reduction of hydrogen storage capacity was likely caused by the accumulation of dislocations introduced during hydrogenation in the early stages of the cycling measurements and by the partial amorphization in the later stages. Mg substitution for Pr in Mg2−xPrxNi4 probably increased the formation energy of the dislocations, which changed the dominant mechanism for releasing the stress induced by hydrogen occupation from the formation of dislocations to pulverization. Mg substitution also decreased the value of RA/RB and suppressed the formation of an amorphous phase. These two factors improved the cyclic property of Mg2−xPrxNi4 with a Mg-rich composition.

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AUTHOR INFORMATION

Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS A part of this work was supported by New Energy and Industrial Technology Development Organization (NEDO) under its “Advanced Fundamental Research Project on Hydrogen Storage Materials” and “Development of technologies for hydrogen production, delivery and storage system”. The synchrotron X-ray experiments were performed under the Shared Use Program of JAEA Facilities (Proposal No. 2011BE09) at JAEA beamline BL22XU in SPring-8 (Proposal No. 2011B3784). 6703



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Degradation Mechanism against Hydrogenation Cycles in Mg2

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