Origin of Degradation in the Reversible Hydrogen Storage Capacity of

Article pubs.acs.org/JPCC

Origin of Degradation in the Reversible Hydrogen Storage Capacity of V1−xTix Alloys from the Atomic Pair Distribution Function Analysis Hyunjeong Kim,*,† Kouji Sakaki,† Hiroshi Ogawa,† Yumiko Nakamura,† Jin Nakamura,‡ Etsuo Akiba,§ Akihiko Machida,∥ Tetsu Watanuki,∥ and Thomas Proffen⊥ †

National Institute of Advanced Industrial Science and Technology, Tsukuba, Ibaraki 305-8565, Japan Japan Metals & Chemicals Co. Ltd., Nishiokitama, Yamagata 999-1351, Japan § International Institute for Carbon-Neutral Energy Research, Kyushu University, Nishi-ku, Fukuoka 819-0395, Japan ∥ Japan Atomic Energy Agency, Sayo, Hyogo 679-5148, Japan ⊥ Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831-6475, United States ‡

S Supporting Information *

ABSTRACT: Reduction in reversible hydrogen storage capacity with increasing hydrogenation and dehydrogenation cycle number is observed in numerous hydrogen storage materials, but the mechanism behind this unfavorable change has not been elucidated yet. In this study, we have investigated the development of structural defects or disorders in V1−xTixH2, x = 0, 0.2, and 0.5, during the first 15 hydrogen absorption and desorption cycles using the atomic pair distribution function (PDF) analysis of synchrotron X-ray total scattering data to find out the possible structural origin of the poor cyclic stability of V1−xTix alloys. While pure vanadium shows no significant change in the PDF, alloy samples subject to several hydrogenation and dehydrogenation cycles display fast decaying of the PDF profile due to a progressive increase in the PDF peak width with increasing r. This r-dependent PDF peak broadening effect becomes stronger with cycle number. Molecular dynamics (MD) simulations demonstrated that dislocation defects explain characteristic features in our experimental PDFs very well and suggested that a large number of dislocations are formed during hydrogen cycling. We found there is a close relation between the reduced amount of the reversible hydrogen content of V0.8Ti0.2 and the amount of generated dislocations. On the basis of the PDF analysis results, a possible mechanism behind degradation in the reversible hydrogen storage capacity of V1−xTix is discussed.

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INTRODUCTION The ability to reversibly absorb and desorb a large amount of hydrogen at ambient conditions makes vanadium (V) very attractive for energy storage application in fuel cell vehicles and large-scale stationary energy storage systems.1 With ∼0.1 MPa of applied hydrogen gas pressure, this body-centered cubic (bcc) metal first absorbs H/M (the number ratio of hydrogen to metal atoms) ∼0.9 of hydrogen at room temperature changing the structure type to body-centered tetragonal (bct).2 This phase is called a monohydride phase. A further increase in hydrogen gas pressure leads to the formation of a dihydride phase (VH2) with a face-centered cubic (fcc) structure (Figure 1). Although V absorbs a total of ∼4 mass % of hydrogen, because of its very low hydrogen desorption pressure of the monohydride phase, little more than half of the absorbed hydrogen can come out via the transition from the fcc dihydride phase to the bct monohydride phase at ambient conditions. This leaves ∼2.1 mass % of hydrogen for a practical use. One drawback of V is its high cost, and therefore alloying with other inexpensive elements is favorable. However, it is well-known that alloying V with other elements often leads to unfavorable changes in hydrogen storage properties such as poor cyclic stability;3 that is to say, the reversible hydrogen © 2013 American Chemical Society

storage capacity gradually decreases as the hydrogen absorption and desorption process is repeated. Numerous studies have reported that V-based bcc alloys with poor cyclic stability show significant broadening of diffraction peaks by hydrogen cycling.3,4 This is also true for various intermetallic compounds and is usually ascribed to structural defects and lattice strains developed during hydrogen cycling.5,6 Furthermore, the structure of a monohydride phase seems to be an important factor to determine the cyclic stability of V-based bcc alloys. As mentioned earlier, pure V forms a bct monohydride, but when it is alloyed with other elements like Ti, the structure of a monohydride phase often changes to bcc. Such alloys showing a bcc monohydride usually exhibit worse cyclic stability and more pronounced cycling-induced broadening of diffraction peaks than those showing a bct monohydride phase.3,4 Although excellent cyclic stability is one of the prerequisites for practical application, the mechanism behind degradation in the reversible hydrogen storage capacity of V-based bcc alloys during hydrogen cycling has not been fully elucidated yet. Received: November 18, 2013 Published: November 22, 2013 26543

dx.doi.org/10.1021/jp408766r | J. Phys. Chem. C 2013, 117, 26543−26550

The Journal of Physical Chemistry C

Article

Table 1. List of Samples for This Study sample

# of cycles

notation

VH2 VH2 V0.8Ti0.2H2 V0.8Ti0.2H2 V0.8Ti0.2H2 V0.8Ti0.2H2 V0.5Ti0.5H2 V0.5Ti0.5H2

1 10 1 5 10 15 1 5

VH2-1 VH2-10 V0.8Ti0.2H2-1 V0.8Ti0.2H2-5 V0.8Ti0.2H2-10 V0.8Ti0.2H2-15 V0.5Ti0.5H2-1 V0.5Ti0.5H2-5

PCT temperature (K) 303 303 413 413 413 413 573 573

K K K K K K K K

hydrogen by each sample was recorded with increasing applied hydrogen gas pressure up to 5 MPa and that of desorbed hydrogen with decreasing applied hydrogen gas pressure down to 0.01 MPa. The PCT measurement was stopped at 5 MPa of the target cycle, and the dihydride sample was deactivated for synchrotron X-ray total scattering measurements. Synchrotron X-ray Experiment. Synchrotron X-ray total scattering experiments were conducted at the Japan Atomic Energy Agency (JAEA) beamline of BL22XU13 at SPring-8. Powder samples of VH2, V0.8Ti0.2H2, and V0.5Ti0.5H2 were loaded in kapton capillaries with a diameter of 1.0 mm. Data were collected at room temperature using the rapid acquisition pair distribution function (RA-PDF) technique14 with a large image plate (IP) detector, R-AXISV, manufactured by Rigaku.13 The IP detector was mounted orthogonal to the incident X-ray beam of 70.16 keV (λ = 0.1768 Å). The sample-to-detector distance was 300 mm. A series of frames were collected for a data set to achieve good statistics. Data Processing and Modeling. Series of data were combined. The signal from an empty container (a kapton capillary) was subtracted from the raw data, and various other corrections were made.7 The X-ray PDFs were obtained by a sine Fourier transformation of the powder diffraction data according to eq 1

Figure 1. Crystal structure of fcc VH2 and its relations with the structure of bct VH. Dark orange circles and thick black lines represent trace of the bct monohydride structure. In this dihydride phase, hydrogen occupies interstitial sites where it is tetrahedrally coordinated with four vanadium atoms.

In this study, we investigate the average and local structural changes of V1−xTixH2, x = 0, 0.2, and 0.5, during early hydrogen absorption and desorption cycles (less than 15 cycles) using synchrotron X-ray total scattering data. The goal of this study is to identify the origin of the hydrogen cycling-induced diffraction peak broadening. For average and local structural studies, Rietveld and the atomic pair distribution function (PDF) analyses are employed, respectively. The PDF is an interatomic distance distribution which gives the probability of finding atom pairs separated by distance r.7 These days this technique is popularly used for studying the local structure of various types of materials such as amorphous, nanocrystalline, and heavily disordered materials.8−12 Our average and local structural studies on V1−xTixH2 suggest that diffraction peak broadening during hydrogen cycling is due to the formation of a large number of dislocation defects, and a reduction in the reversible hydrogen storage capacity is closely related to such defect formation.

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 and S(Q) is the total scattering structure function.7 Because of the unfavorable signal-to-noise ratio at the high-Q regions, Q[S(Q) − 1] was truncated at Qmax = 20 Å−1 before the transformation. The program PDFgetX215 was used for obtaining the X-ray PDFs. For local structural studies the PDFgui16 program was used for real space modeling and calculation. In the PDFgui program, the PDF peak width is given as eq 2

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EXPERIMENT Sample Preparation. V0.8Ti0.2 and V0.5Ti0.5 alloys were synthesized by arc melting of a 20 g portion of V (purity >99.9%) and Ti (purity >99.9%) metal chunks on a watercooled copper crucible under argon atmosphere. As-synthesized button-shaped alloys were annealed at 1673 K under vacuum for 48 h followed by ice−water quenching. The morphology and composition of samples were analyzed using scanning electron microscope (SEM) and energy-dispersive X-ray spectrometer (EDX), respectively. The composition of each sample was close to the target, and no extra phase was found. Cycle Test. Around 2 g of each sample was cut into small pieces, sealed in a stainless steel container, and heated at 573− 1073 K, depending on the composition, for 1 h under vacuum for activation. The pressure−composition isotherm (PCT) measurement was carried out using a Sieverts-type apparatus at a temperature given in Table 1. The amount of absorbed

σij = σij′ 1 −

δ1 δ − 22 + Q broad 2rij 2 rij rij

(2)

where the first term is related to the anisotropic atomic displacement parameters and the second and third terms describe the effects of the correlated atomic motion that sharpens the first PDF peak. The last term, Qbroad2rij2, expresses broadening due to the limited Q resolution of the diffractometer. This is the only term that gives rise to the rdependent PDF peak broadening in the calculated PDF. A detailed explanation for these parameters can be found in ref 17. The limited Q resolution of the diffractometer also causes the exponential damping of PDF peaks via 2

B(r ) = e−(rQ damp) 26544

/2

(3)



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Figure 2. (a) Rietveld fit to X-ray diffraction data of the V0.8Ti0.2H2-10 sample using a fcc structural model. The difference between data and calculation is offset. (b) Change in lattice strain during hydrogen cycling is shown. Strain was obtained from the Rietveld refinements of X-ray diffraction data of VH2 (magenta open triangles), V0.8Ti0.2H2 (cyan close circles), and V0.5Ti0.5H2 (blue open squares).

In general, by measuring relatively well-crystallized materials, such as Ni, the contribution from the instrument (Qdamp and Qbroad) can be readily estimated. Structural modeling in reciprocal space was carried out by the Rietveld method using the RIETAN-FP program.18 The pseudo-Voigt function, a linear combination of Gauss and Lorentz functions, was used to describe the shape of diffraction peaks. The crystallite size and lattice stain were obtained from Rietveld refinement results as described in ref 19.

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RESULTS AND DISCUSSION The synchrotron X-ray diffraction data of all samples were well explained by a fcc structural model. The representative Rietveld fit of V0.8Ti0.2H2-10 is shown in Figure 2(a), and refined structural parameters are given in Table S1 (Supporting Information). Except a negligibly small amount of a monohydride phase in some of the samples, no other extra phase was found. The monohydride phase was probably formed during the deactivation process because there was no consistency in its appearance. With increasing cycle number, the broadening of diffraction peaks was observed. Analysis of the peak shape with a pseudo-Voigt profile function suggests that the broadening is mainly due to the development of large lattice strain. The crystallite size was ∼1000 Å for all the samples, and no considerable cycling-induced crystallite size reduction was found. Figure 2(b) shows lattice strain obtained from the Rietveld analysis. The accumulation of strain over the cycles is clearly displayed. Samples that underwent several hydrogen absorption and desorption cycles show unusually rapid profile damping in their PDFs. In Figure 3, the PDFs of the first (red lines) and fifth or tenth (blue lines) cycles are compared over a wide-r range to illustrate how the profile damping evolves in each sample. Interestingly, depending on the composition, a notable change in the PDF occurs at different times of the cycling period. In the case of VH2 the PDF remains almost unchanged for ten cycles (Figure 3(a)). However, substitution of Ti for 20% of V dramatically alters the outcome. Initially, the PDF peak height of V0.8Ti0.2H2 decays with r in a comparable rate to that of VH21 (red line in Figure 3(b)), but after ten cycles it falls off markedly faster (blue line in Figure 3(b)). A further increase of Ti (or decrease of V) accelerates the decaying (Figure 3(c)); even the PDF of V0.5Ti0.5H2-1 falls off faster than that of V0.8Ti0.2H2-10. On close examination of the X-ray PDFs of V0.8Ti0.2H2, where the evolution of the cycle-induced PDF profile decay was readily seen, we found that the most prominent change

Figure 3. X-ray PDFs of (a) VH2, (b) V0.8Ti0.2H2, and (c) V0.5Ti0.5H2. The PDFs of the first and higher cycle (VH2-10, V0.8Ti0.2H2-10, and V0.5Ti0.5-5) are plotted on top of each other with red and blue lines, respectively.

occurred during the first five cycles, and a rather moderate change was followed in subsequent cycles (Figure 4(a)).

Figure 4. (a) X-ray PDFs of V0.8Ti0.2H2 samples after several hydrogen absorption and desorption cycles. PDF peaks at (b) low- and (c) highr regions are compared.

Furthermore, comparison of low- and high-r regions (Figure 4(b) and (c)) reveals that the PDF peaks get broadened with increasing cycle number, and this effect is larger at higher-r regions. This strong r-dependent peak broadening blurs the features of high-r peaks and leads to the damping of the overall PDF profile. We did not find any extra peak developed in the low-r regions during test cycles. This rules out the possibility of secondary phase formation, i.e., disproportionation. The 26545



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refined model provided much broader PDF peaks in low-r and much sharper PDF peaks in high-r regions than experimental data resulting in a characteristic residual curve shown in Figure 5(a). However, once the Qbroad parameter, the only parameter currently available in the PDFgui program that gives rise to rdependent peak broadening, was refined, the fit was dramatically improved (Figure 5(b)). The Rwp value of 9% is a strong indication that the model successfully reproduces the PDF features over a wide-r range (1.5 ≤ r ≤ 100 Å). This means the Qbroad2rij2 term in eq 2 is acceptable to describe the rdependent peak broadening of our experimental data, and the refined Qbroad value can be tentatively used for a quantitative measure of such broadening. PDF refinement results for other samples are summarized in Table S2 (Supporting Information). As mentioned earlier, the Qbroad parameter originally deals with the broadening due to the limited instrument resolution, and it is normally estimated from standard samples such as Ni. Therefore, the Qbroad values obtained from the PDF refinements above include the contributions from both the sample and the instrument. The sample-related portion can be readily extracted by subtracting the value obtained from the Ni PDF. The Qbroad value after making this correction was plotted against cycle number in Figure 6(a). Figure 6(b) shows the amount of

broadening of PDF peaks reflects the weakening of the structural correlation. Hence, it is most probable that the hydrogen absorption and desorption reaction creates lattice defects or disorders which largely disturb mid-to-long-range structural order. Although the identification of the structural defects or disorders responsible for the r-dependent PDF peak broadening is the most important part of this study, we first estimate the degree of broadening and see whether it has any close relationship to degradation of the reversible hydrogen storage capacity. For this purpose, the PDF refinements using the average fcc structural model were carried out. This fcc model does not include H atoms since their signals in X-ray PDFs are negligible. We first refined a scale factor, a lattice parameter, an isotropic atomic displacement parameter (Uiso) for V and Ti atoms, and a peak sharpening parameter7 until the best fit was obtained. The fitting result for V0.8Ti0.2H2-10 is shown in Figure 5(a), and the refined values of the structural parameters are

Figure 5. Representative fit to the X-ray PDF of V0.8Ti0.2H2-10 using the fcc structural model. (a) A poor fit was obtained when Qbroad was fixed to the value obtained from the Ni data. (b) By refining the Qbroad parameter, the fit was improved significantly. Blue open circles, red lines, and green lines correspond to experimental PDFs, calculated PDFs, and difference curves, respectively. The difference curves are offset for clarity.

Figure 6. (a) Qbroad values obtained from the PDF refinements of V0.8Ti0.2H2 data and (b) the reversible hydrogen storage capacity of V0.8Ti0.2 at 413 K are plotted as a function of cycle number. The reversible hydrogen storage capacity was normalized by the initial value. Red dashed lines are for guidance.

given in Table 2. Apparently, the fit is poor. This is because we did not use any parameter that can handle r-dependent PDF peak broadening during this refinement, and consequently, the

reversibly absorbed hydrogen by V0.8Ti0.2 as a function of cycle number. Interestingly, there is a close correlation between Qbroad and the reversible hydrogen storage capacity. During the first five cycles when the Qbroad value increases quickly, the amount of reversibly absorbed hydrogen is reduced rapidly. In subsequent cycles, changes in both values become more gradual. This leads us to an idea that structural defects or disorders which induce the r-dependent PDF peak broadening probably play a critical role in degradation in the reversible hydrogen storage capacity of V1−xTix. Various structural defects were reported to form in V-based ternary alloys during the hydrogenation and dehydrogenation process. For instance, stacking faults, twin boundaries, and dislocations were observed in the high-resolution transmission electron microscopy (TEM) images of V−Ti−Mn alloys,20 and the formation of vacancies and dislocations was identified in

Table 2. Structural Parameters Obtained from the PDF Refinement of V0.8Ti0.2H2-10 Data Using the fcc (Space Group Fm-3m) Modela a (Å) Uiso (Å2)b Qbroad (Å−1) Rwp (%)d

Qbroad was fixed

Qbroad was refined

4.305(2) 0.0110(5) 0.0077c 47.8

4.305(2) 0.0050(3) 0.060(9) 9.3

The refinement range was 1.5 ≤ r ≤ 100 Å. bThe same Uiso parameter was used for V and Ti atoms. cThe value is from the PDF refinement of Ni data. dRwp indicates the goodness of a fit. a

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Figure 7. (a) Extrinsic stacking fault in the fcc VH2 model is depicted. Only V atoms are shown. (b) Calculated PDFs using the fcc VH2 models with (blue line) and without (red line) stacking faults are compared. (c) Stacking faults do not affect the PDF peaks in low-r regions significantly but (d) do broaden the peaks in high-r regions resulting in PDF profile damping. Some peaks remain sharp like the one indicated by a green arrow.

molecular dynamics (MD). A pair of edge dislocations (Burgers vector of 1/2 ⟨111⟩ on (110) slip planes) was embedded in periodic MD cells filled with bcc V lattice ranging in size from ca. 80 × 80 × 30 Å3 to 210 × 210 × 30 Å3. This allows us to vary dislocation density (ρdislocation) from 3.3 × 1012 cm−2 to 4.6 × 1011 cm−2. The atomic positions were relaxed using the Finnis−Sinclair potential23 under isobaric−isothermal (constant NPT) conditions at 0.1 MPa and 300 K for more than 1 ns. The calculated PDFs using the relaxed dislocation models are shown in Figure 8(a). Similar to stacking faults, dislocations

V0.4Ti0.24Cr0.36 by positron annihilation lifetime measurements.21 If present, all these structural defects could readily disturb and/or possibly reduce the hydrogen occupation sites in the fcc dihydride phase (Figure 1). Hence, we need to examine the effect of each of these structural defects on the PDF. We first consider how stacking faults, disruption in the regular ABCABC... arrangement of {111} planes in the fcc structure, appear in the PDF. For this purpose, the VH2 fcc structure (Fm-3m) was converted into P1 space group, and the unit vectors were redefined in such a way that the {111} planes of the fcc structure became perpendicular to the c-axis of the P1 structure. We made a tetragonal-shaped supercell where an extrinsic stacking fault (Figure 7(a)) was repeated every 54.6 Å along the c-direction when periodic boundary conditions were applied. Note that the stacking fault model was not theoretically relaxed, and H atoms were not included in it because of relatively weak signals of H atoms in X-ray data. In Figure 7(b), the calculated PDFs of the defect-free fcc and stacking fault models are compared. Although the PDF of the stacking fault model decays fast with increasing r, its r-dependent peak broadening is not as pronounced as our experimental PDFs; in contrast to the blurry peaks in our data (Figure 4(c)), the model yields relatively sharp well-defined peaks in high-r regions (Figure 7(d)). More importantly, some of the peaks, such as one indicated by a green arrow in Figure 7(d), remain almost unchanged. Because stacking faults do not disturb the in-plane correlation of a (111) plane, the atomic arrangement of a (111) plane remains the same and the corresponding PDF peaks stay sharp. A more detailed discussion about the effect of stacking faults on the PDF can be found in ref 22. Twin boundaries give similar effects on the PDFs. These two defects are probably present in our samples, but the significant rdependent PDF peak broadening found in our data is not attributed to them because no peak remains sharp in our experimental PDFs. This suggests that there exists another type of defects that can induce much stronger r-dependent peak broadening than stacking faults and twin boundaries. Dislocation is another type of defect found in V-based bcc alloys.20,21 Our main interest is how strain fields around a dislocation core emerge in the PDF. This means it is necessary to have a realistic atomic arrangement of the dislocation core and its vicinity through theoretical calculation. However, the formation process and resultant characteristics of dislocations in fcc VH2 are not well-understood. Therefore, as a first approximation we simulated dislocations in bcc V by classical

Figure 8. (a) Calculated X-ray PDFs using the bcc V models with different dislocation densities. PDF profile decays faster as dislocation density increases. Effect of dislocations, i.e., broadening of PDF peaks, is small in (b) low-r regions but large in (c) high-r regions.

or strain fields around a dislocation core also induce rapid PDF profile decaying. In this case, however, r-dependent PDF peak broadening (Figure 8(b) and (c)) is comparable to the experimental data (Figure 4(b) and (c)), and no peaks remain sharp. In addition, an increase in dislocation density mainly lowers the low- to mid-r peaks, and a similar tendency was observed in samples subject to many cycles (Figure 4(a)). This strongly suggests that dislocations are present in our samples, and the number of dislocations increases with cycle number. After the fifth cycle, the Qbroad value does not change as much as the reversible hydrogen storage capacity does (Figure 6). This is probably because a further increase in the dislocation density does not induce more prominent r-dependent PDF peak broadening as seen in Figure 8(a) (when ρdislocation > 4.6 × 1011 cm−2, change in r-dependent PDF peak broadening is relatively small), and the PDFgui program probably could not 26547



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Figure 9. Calculated PDFs using (a) the VH2 stacking fault model and (b) V dislocation (ρdislocation = 4.6 × 1011 cm−2) model were fit using the defect-free fcc and bcc models, respectively, without refining the Qbroad parameter. Blue open circles, red lines, and green lines represent calculated PDFs using the defect models, fits using the defect-free models, and difference curves, respectively. The difference curves are offset for clarity. (c) and (d) Refinement of the same PDFs as (a) and (b) with Qbroad parameter, respectively. Qbroad parameter did not help explain the effect of stacking faults (c) but reproduced the effect of dislocations on the PDF well (d).

in our experimental PDFs and simulation results of dislocation defects. Strain or stress gradients seem to be manifested in the PDF in different ways depending on their types. Therefore, we might be able to identify types of strain or stress fields present in a sample by analyzing how PDF peak widths increase with r. Although our PDF analysis results strongly indicate that a large number of dislocation defects are formed in V0.8Ti0.2H2 during hydrogen cycling, we cannot rule out the presence of vacancy and planar defects in our samples. It is worth noting that the most probable dislocations in the fcc structure are 1/2 ⟨110⟩ and ⟨001⟩ types.25 1/2 ⟨110⟩ dislocations are energetically more favorable than ⟨001⟩, but this type of perfect dislocation does not occur in metals.25,26 Instead, it dissociates into two partial dislocations (e.g., 1/2 ⟨110⟩ → 1/6 ⟨211⟩ + 1/ 6 ⟨12-1⟩) leaving stacking faults between to reduce the elastic energy. Moreover, vacancy defects reduce the intensities of PDF peaks slightly, and it is difficult to identify them from the PDF in the presence of a large number of dislocations. Therefore, it is more realistic to think that the effects of vacancy and planar defects are overshadowed by those of dislocations in our experimental PDF data. Now we consider possible formation mechanisms of lattice defects. As mentioned earlier, the reversible hydrogen storage capacity that we are interested in is the amount of hydrogen that can be absorbed and desorbed during the transition between a bct or bcc monohydride and an fcc dihydride phase. During the hydrogenation process, a bcc lattice stretches along one of its lattice vector directions transforming into an fcc lattice (Figure 1). There are three possible elongation directions in a bcc lattice. If there are multiple origins of growth of a fcc phase, uncertainty in the growth direction leads to lattice mismatch where fcc phases grown in different directions meet and various types of lattice defects like vacancies, dislocations, stacking faults, and twin boundaries can be formed. On the other hand, the elongation direction of a bct cell is rather unambiguous. There is a strong chance that a bct cell will be stretched along the long axis only to transform into an fcc structure. This leads to a much reduced chance to

catch such a small change. Another possible reason is that other kinds of defects may start to dominate after the fifth cycle. The PDF of V0.5Ti0.5H2-1 decays faster than that of V0.8Ti0.2H2-10. This is probably because V0.5Ti0.5 has a large number of structural defects or disorders from the first cycle or from the alloy phase. To examine our defect models further, we attempted to fit the fcc and bcc models to the calculated PDFs of the stacking fault and dislocation (ρdislocation = 4.6 × 1011 cm−2) models, respectively. As expected, poor fit was obtained in both cases when the Qbroad value was fixed to one from the Ni PDF (Figure 9(a) and (b)). Note the difference curve in Figure 9(b). This curve is very similar in appearance to the one found in V0.8Ti0.2H2-10 with a fixed Qbroad (Figure 5(a)). When Qbroad was refined, a dramatically improved fit was obtained only for the PDF of the dislocation model (Figure 9(d)). This behavior is consistent with our data. On the contrary, the addition of the Qbroad parameter did not yield an improved fit to the stacking fault PDF (Figure 9(c)). This is partly because the Qbroad parameter does not reproduce sporadically appearing sharp peaks arising from atom pairs residing on {111} planes. Moreover, r-dependent PDF peak broadening due to stacking faults seems to occur more slowly than Qbroad2rij2 in eq 2. At any rate, the Qbroad2rij2 term works relatively well to reproduce dislocation-induced changes in the PDF. Bear in mind that to model the dislocation effect properly and to obtain any meaningful information from fitting results we need a more exact mathematical expression for strain fields of dislocations in eq 2. Our MD simulation shows that randomly distributed vacancy defects do not cause the attenuation of the PDF. This is in good agreement with the simulation result demonstrated by Gibson that a Gaussian random atomic displacement does not induce the rapid damping of the PDF profile.24 Interestingly, Gibson also showed that a strain or stress gradient or simply the existence of curved lattice planes progressively increases the width of peaks in the PDF as r increases and consequently attenuates the PDF profile.24 This is exactly what we have seen 26548



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on Hydrogen Storage Materials (HYDRO-STAR). 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.2011A3703 and 2011B3784).

form lattice defects. Tetrahedral sites for hydrogen at or near the defects are distorted. Over many hydrogenation and dehydrogenation cycles they can get more heavily distorted and may no longer be occupied by hydrogen resulting in degradation in the reversible hydrogen storage capacity. This is probably the reason why V-based alloys showing a bcc monohydride have poor cyclic stability and those showing a bct monohydride have excellent cyclic stability.3,4 In the case of our samples, while V shows a bct monohydride phase and excellent cyclic stability, V1−xTix, x = 0.2 and 0.5, shows a bcc monohydride phase and poor cyclic stability. In fact, the importance of lattice defects in the decrease in the reversible hydrogen storage capacity of hydrogen absorbing intermetallic compounds was realized earlier. The cyclic stability of LaNi5, one of renowned hydrogen-absorbing intermetallic compounds, can be dramatically improved by substituting a small amount of Sn on Ni sites.27 This is because Sn somehow effectively suppresses the formation of vacancy and dislocation defects during hydrogen absorption and desorption processes.28 Therefore, to improve the cyclic stability of hydrogen storage alloys, we need to restrain the formation of lattice defects during hydrogen absorption and desorption processes. For Vbased alloys, this can be achieved by developing alloys showing a bct monohydride with a large tetragonal distortion. An investigation on distorted hydrogen occupation sites in cycled V1−xTixH2 samples using the neutron PDFs is underway.

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(1) Schlapbach, L.; Züttel, A. Hydrogen-Storage Materials for Mobile Applications. Nature 2001, 414, 353−358. (2) Manchester, F. D. Phase Diagrams of Binary Hydrogen Alloys; ASM International: Materials Park, OH, 2000. (3) Itoh, H.; Arashima, H.; Kubo, K.; Kabutomori, T.; Ohnishi, K. Improvement of Cyclic Durability of BCC Structured Ti-Cr-V Alloys. J. Alloys Compd. 2005, 404, 417−420. (4) Aoki, M.; Noritake, T.; Ito, A.; Ishikiriyama, M.; Towata, S. Improvement of Cyclic Durability of Ti-Cr-V Alloy by Fe Substitution. Int. J. Hydrogen Energy 2011, 36, 12329−12332. (5) Bowman, R. C., Jr.; Luo, C. H.; Ahn, C. C.; Witham, C. K.; Fultz, B. The Effect of Tin on the Degradation of LaNi5‑ySny Metal Hydrides During Thermal Cycling. J. Alloys Compd. 1995, 217, 185−192. (6) Č erný, R.; Joubert, J.-M.; Latroche, M.; Percheron-Guégan, A.; Yvon, K. Anisotropic Diffraction Peak Broadening and Dislocation Substructure in Hydrogen-Cycled LaNi5 and Substitutional Derivatives. J. Appl. Crystallogr. 2000, 33, 997−1005. (7) Egami, T.; Billinge, S. J. L. Underneath the Bragg Peaks: Structural Analysis of Complex Materials; Pergamon Press Elsevier: Oxford, England, 2003. (8) Proffen, T.; Kim, H. J. Advances in Total Scattering Analysis. J. Mater. Chem. 2009, 19, 5078−5088. (9) Malavasi, L. Total Scattering Investigation of Materials for Clean Energy Applications: The Importance of the Local Structure. Dalton Trans. 2011, 40, 3777−3788. (10) Kim, H.; Nakamura, J.; Shao, H.; Nakamura, Y.; Akiba, E.; Chapman, K.; Chupas, P.; Proffen, T. Local Structural Evolution of Mechanically Alloyed Mg50Co50 Using Atomic Pair Distribution Function Analysis. J. Phys. Chem. C 2011, 115, 7723−7728. (11) Kim, H.; Nakamura, J.; Shao, H.; Nakamura, Y.; Akiba, E.; Chapman, K.; Chupas, P.; Proffen, T. Insight into the Hydrogenation Properties of Mechanically Alloyed Mg50Co50 from the Local Structure. J. Phys. Chem. C 2011, 115, 20335−20341. (12) Kim, H.; Nakamura, J.; Shao, H.; Nakamura, Y.; Akiba, E.; Chapman, K.; Chupas, P.; Proffen, T. Variation in the Ratio of Mg2Co and MgCo2 in Amorphous-like Mechanically Alloyed MgxCo100‑x Using Atomic Pair Distribution Function Analysis. Z. Kristallogr. 2012, 227, 299−303. (13) Watanuki, T.; Machida, A.; Ikeda, T.; Ohmura, A.; Kaneko, H.; Aoki, K.; Sato, T. J.; Tsai, A. P. Development of a Single-Crystal X-ray Diffraction System for Hydrostatic-Pressure and Low-Temperature Structural Measurement and Its Application to the Phase Study of Quasicrystals. Philos. Mag. 2007, 87, 2905−2911. (14) Chupas, P. J.; Qiu, X.; Hanson, J. C.; Lee, P. L.; Grey, C. P.; Billinge, S. J. L. J. Rapid-Acquisition Pair Distribution Function (RAPDF) Analysis. Appl. Crystallogr. 2003, 36, 1342−1347. (15) Qiu, X.; Thompson, J. W.; Billinge, S. J. L. PDFgetX2: A GUI Driven Program to Obtain the Pair Distribution Function from X-ray Powder Diffraction Data. J. Appl. Crystallogr. 2004, 37, 678. (16) Farrow, C. L.; Juhas, P.; Liu, J. W.; Bryndin, D.; Božin, E. S.; Bloch, J.; Proffen, Th.; Billinge, S. J. L. PDFfit2 and PDFgui: Computer Programs for Studying Nanostructure in Crystals. J. Phys.: Condens. Matter 2007, 19, 335219. (17) Farrow, C. L.; Juhás, P.; Liu, J. W.; Bryndin, D.; Božin, E. S.; Bloch, J.; Proffen, Th.; Billinge, S. J. L. PDFgui user guide, 2009. (18) Izumi, F.; Momma, K. Three-Dimensional Visualization in Powder Diffraction. Solid State Phenom. 2007, 130, 15−20. (19) Nakamura, Y.; Akiba, E. In-situ X-ray Diffraction Study on LaNi5 and LaNi4.75Al0.25 in the Initial Activation Process. J. Alloys Compd. 2000, 308, 309−318.

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CONCLUSIONS We have investigated the origin of degradation in the reversible hydrogen storage capacity of V1−xTix alloys by examining local structural changes in V1−xTixH2, x = 0.2 and 0.5, as well as pure V during early hydrogen cycling (less than 15 cycles). Our experimental X-ray PDFs of V1−xTixH2, x = 0.2 and 0.5, show significant r-dependent peak broadening, and this effect becomes stronger with cycle number. The PDF analysis and MD simulation results indicate that dislocations are responsible for such broadening of the PDF peaks, and the number of dislocations increases with cycle number. We found a correlation between reduction in the reversible hydrogen storage capacity of V0.8Ti0.2 and increase in the density of dislocation defects. Our PDF analysis results strongly suggest that dislocations play an important role in reduction in the reversible hydrogen storage capacity of V-based bcc alloys during early hydrogen cycling.

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ASSOCIATED CONTENT

S Supporting Information *

Rietveld and PDF analysis results on synchrotron X-ray data of V1−xTixH2, x = 0, 0.2, and 0.5. This material is available free of charge via the Internet at http://pubs.acs.org.

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REFERENCES

AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS H.K. thanks Nobuhiko Takeichi and Itoko Matsumoto for help with experiments. This work was partly supported by the New Energy and Industrial Technology Development Organization (NEDO) under the Advanced Fundamental Research Project 26549



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(20) Matsuda, J.; Nakamura, Y.; Akiba, E. Microstructure of Ti-V-Mn BCC Alloys Before and After Hydrogen Absorption-Desorption. J. Alloys Compd. 2011, 509, 4352−4356. (21) Kawasuso, A.; Arashima, H.; Maekawa, M.; Itoh, H.; Kabutomori, T. TiCrV Hydrogen Storage Alloy Studied by Positron Annihilation Spectroscopy. J. Alloys Compd. 2009, 486, 278−283. (22) Neder, R. B.; Korsunskiy, V. I. Structure of Nanoparticles from Powder Diffraction Data Using the Pair Distribution Function. J. Phys.: Condens. Matter 2005, 17, S125−S134. (23) Finnis, M. W.; Sinclair, J. E. A Simple Empirical N-Body Potential for Transition-Metals. Phil. Mag. 1984, A50, 45−55. (24) Murray Gibson, J. Understanding the Limits of Pair-Distribution Functions for Nanoscale Correlation Function Measurement. J. Phys.: Condens. Matter 2007, 19, 455217. (25) Hull, D.; Bacon, D. J. Introduction to Dislocations, 5th ed.; Butterworth Heinemann, Oxford, England, 2011. (26) Häkkinen, H.; Mäkinen, S.; Manninen, M. Edge Dislocations in Fcc Metals: Microscopic Calculations of Core Structure and Positron States in Al and Cu. Phys. Rev. B 1990, 41, 12441−12453. (27) Lambert, S. W.; Chandra, D.; Cathey, W. N.; Lynch, F. E.; Bowman, R. C., Jr. Investigation of Hydriding Properties of LaNi4.8Sn0.2, LaNi4.27Sn0.24 and La0.9Gd0.1Ni5 after Thermal Cycling and Aging. J. Alloys Compd. 1992, 187, 113−135. (28) Sakaki, K.; Date, R.; Mizuno, M.; Araki, H.; Nakamura, Y.; Shirai, Y.; Bowman, R. C., Jr.; Akiba, E. Behavior of Vacancy Formation and Recovery During Hydrogenation Cycles in LaNi4.93Sn0.27. J. Alloys Compd. 2009, 477, 205−211.

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Origin of Degradation in the Reversible Hydrogen Storage Capacity of

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