Performance Improvement of VâFeâCrâTi Solid State Hydrogen

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Performance Improvement of V-Fe-Cr-Ti Solid State Hydrogen Storage Materials in Impure Hydrogen Gas Ulrich Ulmer, Daria Oertel, Thomas Diemant, Christian Bonatto Minella, Thomas Bergfeldt, Roland Dittmeyer, R. Jürgen Behm, and Maximilian Fichtner ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.7b13541 • Publication Date (Web): 19 Dec 2017 Downloaded from http://pubs.acs.org on December 20, 2017

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ACS Applied Materials & Interfaces

Performance Improvement of V-Fe-Cr-Ti Solid State Hydrogen Storage Materials in Impure Hydrogen Gas Ulrich Ulmera,*, Daria Oertela, Thomas Diemantb, Christian Bonatto Minellaa,c, Thomas Bergfeldtd, Roland Dittmeyere, R. Jürgen Behmb,c and Maximilian Fichtnera,c

a Karlsruhe Institute of Technology (KIT), Institute of Nanotechnology, P.O. Box 3640, D-76021 Karlsruhe, Germany b Ulm University, Institute of Surface Chemistry and Catalysis, D-89069 Ulm, Germany c Helmholtz Institute Ulm (HIU) for Electrochemical Energy Storage, Helmholtzstr. 11, D-89081 Ulm, Germany d Karlsruhe Institute of Technology (KIT), Institute of Applied Materials, P.O. Box 3640, D-76021 Karlsruhe, Germany e Karlsruhe Institute of Technology (KIT), Institute for Micro Process Engineering, P.O. Box 3640, D-76021 Karlsruhe, Germany

Keywords: Hydrogen storage, surface reactions, impure hydrogen, metal hydrides, surface engineering

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Abstract Two approaches of engineering surface structures of V-Ti-based solid solution hydrogen storage alloys are presented, which enable improved tolerance towards gaseous oxygen (O2) impurities in hydrogen (H2) gas. Surface modification is achieved through engineering lanthanum (La)- or nickel (Ni)-rich surface layers with enhanced cyclic stability in a H2/O2 mixture. The formation of a Ni-rich surface layer does not improve the cycling stability in H2/O2 mixtures. Mischmetal (Mm, a mixture of La and Ce) agglomerates are observed within the bulk and surface of the alloy when small amounts of this material are added during arc melting synthesis. These agglomerates provide hydrogen-transparent diffusion pathways into the bulk of the VTi-Cr-Fe hydrogen storage alloy when the remaining oxidized surface is already intransparent for hydrogen. Thus, the cycling stability of the alloy is improved in an O2-containing hydrogen environment as compared to the same alloy without addition of Mm. The obtained surface-engineered storage material still absorbs hydrogen after 20 cycles in a hydrogen-oxygen mixture, while the original material is already deactivated after 4 cycles.

1. Introduction Solid state hydrogen storage systems such as metal hydrides offer advantages over high-pressure or liquid hydrogen storage systems. These include high volumetric hydrogen densities at moderate temperature and pressure conditions and, depending on the storage material, high gravimetric hydrogen density.1,2 Rapid hydriding and dehydriding kinetics are crucial requirements in order to enable hydrogen uptake and release from the storage material at desired rates.3 Storage materials show maximum


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reaction rates if the metal surfaces are clean and free of adsorbed or chemisorbed impurities, such as CO, CO2, hydrocarbons, ammonia, water or oxygen. Then, catalytic sites promoting the chemisorption and dissociation of hydrogen molecules remain easily accessible for hydrogen. Coverage of the surface by impurities hinders access to the catalytic sites either through the formation of a diffusion barrier formed by adsorbed surface species or through the formation of new chemical species, such as metal oxides, which block the catalytic sites necessary to promote hydrogen adsorption and dissociation. This so-called surface “poisoning” effect can have dramatic consequences on the hydriding kinetics and usable capacity.4–12 AB-, AB2- and AB5-type alloys, such as TiFe, ZrV2, LaNi5 as well as V-Ti-based solid solution alloys have been studied with respect to their behavior upon cycling in impure hydrogen.4–8,13–19 Two main conclusions can be drawn from these studies: 1) the tolerance of most metal hydrides towards critical impurities such as oxygen is too low for most technical applications, and 2) LaNi5 and its related compounds exhibit the highest tolerance towards impurities. The improved tolerance of LaNi5 towards impurities is linked to its surface properties. It is known that the chemical composition at the surface of LaNi5 is different from the composition in the bulk material. Lanthanum (La) exhibits a higher concentration at the surface than in the bulk, and the ratio of La to Nickel (Ni) is 1:1 near the surface (it is 1:5 in the bulk). Near the surface, clusters of metallic Ni are embedded into a matrix of La oxide which allow a homolytic splitting of hydrogen molecules into hydrogen atoms. Furthermore, the hydrogen solubility of La oxide is sufficient to allow significant diffusive transport of hydrogen molecules and atoms. The combination of surface La oxide (solubility of H2 and H) and stable Ni clusters (H2 splitting) is believed to be responsible for the high tolerance of LaNi5 towards oxygen.20



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Different methods to protect metal hydrides from passivation by gaseous impurities have been developed. These include the fluorination of the surface of rare earth metal containing hydrides,21–23 or the deposition of Pd, Pt, Cu or combinations of these elements on the surface of H absorbing alloys.15,24–27 The common idea in these approaches is that the alloy surface gets chemically functionalized or coated with elements that exhibit a reduced reactivity toward gaseous impurities. At the same time, the ability to dissociate hydrogen molecules is retained and the diffusion pathways for the hydrogen atoms to the core of the hydrogen absorbing alloy are provided. A disadvantage of the abovementioned surface treatments is that additional steps during the preparation of the hydrogen storage alloys are necessary, and extra costs are generated. Fluorination is only effective in the case of rare earth containing alloys, and the deposition of Pd or Pt noble metals significantly increases the cost of the storage material due to the high cost and low availability of these metals. Hence, a smart, low-cost surface treatment is still needed. In this work, two surface engineering approaches are presented, which aim at improving the durability upon cycling of V-Fe-Cr-Ti-type solid solution alloys in an oxygen-containing hydrogen environment. V-Fe-Cr-Ti solid solution alloys exhibit higher gravimetric and volumetric hydrogen densities than AB-, AB2- and AB5-type hydrogen storage materials, which makes them attractive for stationary and mobile hydrogen storage applications.28–32 In both surface engineering approaches, the surface structure of LaNi5, the hydrogen storage alloy showing the best cycling stability, is partly imitated through the addition of the A-element (La) or the Belement (Ni) of LaNi5. In the first approach, a small amount of nickel (Ni) is introduced into a V-Fe-Cr-Ti-based solid solution. The introduction of a large amount


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of Ni into V-Ti-based solid solution alloys is known to negatively affect the microstructure and hydrogen storage properties because secondary phases are formed.33–35 Selective leaching of V, Fe, Cr and Ti is therefore applied in order to enrich the alloy surface with Ni while maintaining a low Ni concentration in the bulk. The effect of a Ni-rich surface layer on the hydrogenation/dehydrogenation kinetics in the presence of oxygen impurities is investigated. The second approach involves the introduction of a few atomic percent of rare earth metal in the form of „mischmetal“ (Mm), a mixture of lanthanum and cerium, into the solid solution. Rare earth metals do not dissolve in the V-Ti-Fe-Cr solid solution and are present in the form of agglomerates distributed throughout the alloy.28,36 It is investigated whether the rare earth elements (and/or their oxides) could provide diffusion pathways for hydrogen into the alloy when the surface of the V-Fe-Cr-Ti solid solution is rendered passive by oxide formation. Relevant prior work includes, e.g., surface modification of AB5-based hydrogen storage alloys through base leaching, which has been shown to yield Ni-rich clusters in the surface region. This improves the high-rate dischargeability of the corresponding Ni-metal hydride battery.37 Although studies on the surface properties of V-Ti-based solid solution alloys have already been reported,4,38 our work represents the first thorough investigation which encompasses detailed surface engineering methods and explores the consequences of applying these methods for the hydrogen storage properties in a contaminated H2 environment. V40Fe8TixCryNizbased alloys with several different Ti, Cr and Ni contents were used in this work aiming to engineer high-capacity, high-performance hydrogen storage alloys.

2. Experimental



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V, Ti, Fe, Cr, Ni and Mm (composition was 43 % La, 56 % Ce and 1 % other rare earth metals (Pr, Nd)) raw materials were used for alloy formation. The purity of all materials was > 99.9 %. The alloys reported in this study were prepared by arc melting in a water-cooled copper crucible in an arc melter AM by Edmund Bühler. The alloys were melted, turned around and re-melted four times to improve homogeneity. No annealing treatment was performed. Six different materials were prepared

in

this

work:

V40Fe8Ti28Cr24,

V40Fe8Ti28Cr26Ni8,

V40Ni8Ti28Cr24,

V40Fe8Ti32Cr20, V40Fe8Ti32Cr16Ni4, and V40Fe8Ti26Cr26-6%Mm. The majority of Mm was present at the surface of the Mm-modified alloy ingots. At this point of the material preparation, the ingots had a diameter of ca. 3 cm and the surface was covered by Mm, which had segregated from the V-Ti-based phase during the solidification process. Surface Mm was removed by grinding off the surface of the ingots after melting, and the remaining Mm referred to in the results section of this work was then present in the form of agglomerates distributed within the alloy matrix. The V40Fe8Ti32Cr16Ni4 alloy was used for the leaching experiments. This composition was chosen because the resulting hydride is partly stable at ambient temperature and pressure conditions. Reactivation after the leaching procedure could be achieved when a hydrogenated alloy was leached and subjected to thermal desorption subsequently. The particle size of the metal hydride powder decreases during the initial activation cycles. In order to identify the particle size evolution during the activation, particle size distributions were analysed in preparation for the surface modification and kinetic experiments. Results are shown in Figure S1 in the Supporting Information. Based on this analysis, the material was activated during 15 hydrogenation cycles before applying surface modification procedures or performing kinetic experiments.


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In order to quantify the amount of leached material from the solid surface, the filtrate was analysed by inductively coupled plasma optical emission spectrometry (ICPOES). The number of moles of dissolved atoms in the filtrate was calculated based on the ICP-OES results, and an enrichment factor was calculated after leaching for each element according to equation (1):

=

, ,

,

(1)

where , stands for the molar fraction of element in solution (considering only the metal ions) and , stands for the molar fraction of element in the solid. If > 1, the element was enriched in the electrolyte, if = 1, the element was dissolved in the same molar fraction as present in the solid, and if < 1 the element was enriched in the solid. A larger amount of material (4.5 g) was leached with 9 M NaOH solution at 40 °C for 30 min and used to investigate the hydrogen storage properties of the resulting material. The detailed experimental procedures and equipment used for the leaching experiments and more information about the ICP-OES equipment used in this work can be found in section Experimental-S (Figure S2, Tables S1 and S2) in the Supporting Information. XPS measurements were performed in order to obtain quantitative information on the chemical states and elemental concentration depth profiles before and after cycling in impure H2 gas and before and after the leaching experiments. A PHI 5800 MultiTechnique ESCA system (Physical Electronic) was used. The spectra were acquired using monochromatic Al Kα (1486.6 eV) radiation. The photoelectrons were collected at a detection angle of 45°, using pass energies of 93.9 eV and 29.35 eV at the analyzer for survey and detail spectra, respectively. The surface layers were 7 ACS Paragon Plus Environment


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removed by Ar+-ion sputtering (Isp ≈ 1 µA, Usp = 5 kV). Approximately 1 nm was removed per minute of sputtering. All samples were transferred into the XPS apparatus under inert gas conditions. Detailed information about the SEM and XRD analysis is provided in the Experimental-S section in the Supporting Information. Pressure-composition isotherms were recorded with a home-built Sieverts apparatus, which has been described elsewhere.39,40 The total hydrogen capacity of each material was measured after flushing and evacuation of the apparatus at 5.0 MPa H2 and 298 K. A Swagelok reactor made of stainless steel was used for the hydrogen sorption experiments in contaminated H2. An empty reactor was measured as a reference at the same pressure and temperature conditions that were applied for the measurements of the samples. The whole system was evacuated and the valve between the reactor and the rest of the apparatus was closed. 1 MPa of H2 was then filled into the rest of the apparatus. When the valve between the reactor and the rest of the apparatus was opened, the pressure decreased as a result of hydrogen gas flowing into the reactor. In the case of the empty reactor without sample, a pressure drop by about 0.06 MPa was recorded. When the reactor was filled with an absorbing sample, the final pressure was lower than in the case of the empty reactor. The pressure difference between the measurements with and without sample corresponds to the absorption capacity of the respective sample. After the initial activation cycles, the capacity was assumed as 100 %, and the pressure difference was used to normalize the reacted fraction to 100 % (full capacity; reacted fraction = 1). If during a subsequent experiment in impure hydrogen gas the final pressure was higher than during the initial absorption, it was assumed that less hydrogen was absorbed, which corresponds to a reduction of the reacted fraction. Depending on the surface properties of the respective sample, a slower pressure decrease also corresponds to slower absorption kinetics. During


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hydrogen sorption experiments, the reactor was immersed in a thermal oil bath. The temperature during the absorption of pure H2 was measured at the outside of the reactor during the first experiments. Due to the small amount of sample (approx. 0.3 g), no temperature variation was noticed. Unless noted otherwise, both hydrogen absorption and desorption were performed at 303 K. An example of the pressure curves recorded during hydrogen absorption and the resulting kinetic curves is provided in Figure S3 in the Supporting Information.

3. Results 3.1 Microstructure of the as-cast materials XRD patterns of the as-cast materials are presented in Figure S4 in the Supporting Information. All materials exhibited a body-centered cubic structure. Peaks of minor secondary phases were observed, which can be ascribed to the C14 Laves phase. Lattice parameters of the obtained samples are summarized in Table S3. SEM micrographs of the as-cast samples are presented in Figures S5 in the Supporting Information. As an example, EDX elemental distribution maps of V40Fe8Ti28Cr16Ni8 are shown in Figure 1.



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Figure 1. EDX elemental distribution maps of V40Fe8Ti28Cr16Ni8 in the as-cast state.

All samples exhibited a V-Cr-rich bcc main phase and secondary phases. The C14 Laves phase was observed in all alloys (see Table S3 for detailed compositions). When a high fraction of Ni was inserted into the alloys (8 at% in the case of V40Fe8Ti28Cr26Ni8 and V40Ni8Ti28Cr24), a Ti-Ni-containing phase precipitated within the alloys. This phase formed a three-dimensional network around the bcc main phase, which is in line with previous observations.34,35 For the second series of alloys prepared for this work, the Ni content was reduced to 4 at% (V40Fe8Ti32Cr16Ni4). In that case, no TiNi phase was observed, which suggests that the Ni concentration needs to be higher for this phase to precipitate.


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Figure 2 shows the SEM micrograph of V40Fe8Ti26Cr26-6at% Mm in the as-cast state. Back-scattered SEM images of a related material (V40Fe8Ti26Cr26 with 3 mass% La) are shown in Figure S6.

Figure 2. SEM micrograph of V40Fe8Ti26Cr26-6%Mm in the as-cast state.

Mm inclusions of diameters between 0.5 – 2 µm are distributed evenly throughout the alloy. The composition of the Mm inclusions as determined by EDX was 43 at% La, 56 at% Ce and 1 at% other rare earth metals (Pr, Nd). An oxygen concentration of 59 at% of the total content was also detected, which is not included in the above calculation. Rare earth metal inclusions are known to be formed during arc melting of V-Ti-based solid solution alloys, as the solubility of rare earth metals in the V-Ti main phase is low.28,41

3.2 Cycling and surface properties of unmodified materials The passivation and reactivation behavior of a V-Ti-based solid solution alloy was investigated during cycling in H2 containing 250 ppm O2. Figure 3 shows the hydrogen absorption as a function of time of V40Fe8Ti32Cr20 in pure hydrogen, hydrogen containing 250 ppm O2, and the subsequent reactivation in pure H2. The



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absorption pressure was 1 MPa, whereas desorption was done in dynamic vacuum (continuous vacuum pumping). The temperature was 303 K.

Figure 3. Hydrogen absorption curves of V40Fe8Ti32Cr20 at 1 MPa H2, H2 containing 250 ppm O2 (A) and during reactivation in pure H2 after four cycles in O2-contaminated H2 at 1 MPa and 303 K (B). Numbers at the curves indicate cycle number.

A reacted fraction of 100 % (1.0) corresponds to 2.3 mass% H. The absorption behavior during the first cycle in H2 with 250 ppm O2 was similar to that observed in pure H2. A slight reduction in final capacity was observed during the second cycle. A more pronounced reduction to around 50 % of the original capacity was recorded during the third cycle. No hydrogen was absorbed during the fourth cycle. These results indicate an oxygen-induced surface passivation, which becomes more pronounced with an increasing number of reaction cycles. It is likely that this passivation layer inhibits hydrogen dissociation and absorption of the hydrogen 12 ACS Paragon Plus Environment

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atoms. A similar cyclic response in O2-blended H2 has been observed before for TiFe,8 Zr0.9Ti0.1V2 5 or Ti25V35Cr40.13 The storage material could successfully be reactivated by thermal treatment at 150 °C under dynamic vacuum und subsequent cycling in pure H2 at room temperature. About 90 % of the storage capacity was restored after seven cycles. This corresponds to approximately 1.9 mass% H. It should be noted that the initial heating step was necessary to reactivate the surface sites for hydrogen absorption. Such thermal reactivation treatment has been reported for Ti1.13Fe by Hirata et al.

42

and by Ulmer

et al. for V40Fe8Ti26Cr26 deactivated by organic substances.4 Hirata et al. attribute their observations to the fact that hydrogen desorption is never complete at room temperature but some of the hydrogen remains absorbed in each cycle. The formation and growth of surface oxides plays an important role in the accumulation of hydrogen upon cycling. Thermal desorption competes and goes along with the partial reduction of surface metal oxides and thereby reactivates the surface sites to promote dissociative adsorption and diffusion of H into the bulk.42,43 In order to reveal the passivation mechanism of oxygen-contaminated V40Fe8Ti32Cr20, XPS measurements of the surface and subsurface region were performed. As a reference, the Ti 2p, V 2p and O 1s XP spectra of V40Fe8Ti26Cr26 cycled in pure H2 were also measured. The corresponding results are presented in Figures S7 and S8 of the Supporting Information. XPS confirms that metal oxides are formed at the surface during cycling in H2 blended with O2. This is also in good agreement with previous studies on the passivation behavior and surface structure of other intermetallic hydrogen storage materials.6–8,11,12,22,44



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3.3 Surface modifications through selective leaching and rare earth metal inclusion In section 3.1, we present and discuss the introduction and subsequent enrichment of Ni at the surface of the storage material. Fe or Cr were selected as substitution elements for Ni due to their similar atomic radii and comparable effects on the thermodynamics of hydride formation and dissociation.45 Several compositions of Cror Fe-substituted V-Ti-based solid solution alloys were tested with regard to their respective hydrogen storage properties. The hydrogen capacities and the equilibrium pressure of the tested alloys are summarized in Table 1.

Total hydrogen Plateau pressure Alloy capacity [MPa] [mass%] 3.4 ± 0.1 0.03 (303 K) V40Fe8Ti32Cr20 3.3 ± 0.1 0.09 (303 K) V40Fe8Ti32Cr20Ni4 3.3 ± 0.1 0.07 (298 K) V40Fe8Ti28Cr24 1.8 ± 0.1 0.3 (298 K) V40Fe8Ti28Cr16Ni8 1.7 ± 0.1 0.01 – 1.2 (298 K) V40Ni8Ti28Cr24 1.7 ± 0.1 0.01 – 1.2 (298 K) V40Ni8Ti26Cr26 Table 1. Hydrogen capacities and equilibrium plateau pressures of the materials studied in this work.

All alloys prepared with 8 at% Ni exhibited a reduced total hydrogen capacity as compared to the materials without Ni. This observation is in agreement with previous reports.33–35 It is known that a high fraction of TiNi-based phase, which is formed at higher Ni amount added to the alloy, reduces the hydrogen capacity. TiNi has a lower hydrogen capacity than the V-Ti-based body-centered cubic phase.33–35 At 4 at% Ni, on the other hand, a large total capacity of 3.3 mass% was measured, which was close to the material prepared without Ni. Hydrogen capacities in the range of 3.3 – 3.4


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mass% H are in good agreement with previous reports on V-Ti-based alloys with similar compositions.28,29,46,47 The surface treatment presented in the following paragraphs was performed by dispersing the activated storage alloy powders in aqueous solutions at various pH values (Table S5 in Supporting Information summarizes all varied parameters). Trace amounts of gaseous water in H2 are known to passivate the surface of activated hydrogen storage alloys.7,8 Therefore, suitable reactivation procedures had to be developed to enable hydrogen sorption after the dispersion of the storage alloy powders in aqueous solutions. Several attempts to reactivate the surface-treated alloy powders in their dehydrogenated state at elevated temperatures and pressures between vacuum and 10 MPa H2 remained unsuccessful. However, reactivation of the surfacetreated materials was successful if (partly) absorbed hydrides were used instead of the desorbed materials. Thermal desorption of the contained hydrogen effectively reactivated the surface, and subsequently hydrogen could be ab- and desorbed. Based on the high capacities and low equilibrium pressure values well below 0.1 MPa at ambient conditions, V40Fe8Ti32Cr20 and V40Fe8Ti32Cr16Ni4 were chosen as suitable materials for the leaching experiments. Several acids (H2SO4, HCl, HNO3, HCl + HNO3) and bases (NaOH, Na2CO3) were tested upon their ability to selectively leach V, Fe, Cr and Ti while retaining the Ni in its solid state. Unfavorable factors of enrichment were found for acidic leaching (αi ≤ 1 for V, Fe, Cr, Ti; αi ≥ 1 for Ni), however, the desired dissolution of V and Ti could be achieved (αi ≥ 1 for V, Ti; αi ≤ 1 or 0 for Ni) by leaching under alkaline conditions at pH values > 7. NaOH was selected as an appropriate electrolyte, and several leaching parameters (concentration, temperature, leaching time and ratio between solid and electrolyte) were systematically varied to identify the optimum conditions



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(cNaOH = 9 mol/L, T = 40 °C, S/L = 0.05 g/mL, whereby S/L is defined as the ratio between the mass of solid alloy and electrolyte volume; see below for further explanations). In Figure 4 the factors of enrichment are shown as a function of the NaOH concentration.

Figure 4. Factor of enrichment αi as a function of the NaOH concentration in the electrolyte. The experiments were performed at 60 °C, 30 min leaching time, 0.15 g powder in 3 mL electrolyte.

No Ni was detected by ICP-OES in the electrolyte, indicating that it is not dissolved during leaching in NaOH. The factor of enrichment was αNi = 0 for all tested concentrations of NaOH. For V, αV was high at low NaOH concentrations (0.5 M and 2 M), whereas it decreased at intermediate concentrations and increased again at cNaOH = 9 mol/L. For Ti, the highest factors of enrichment are observed for NaOH concentrations of 8 and 9 mol/L. Ti and V are known to be strongly corroded upon contact with gaseous oxygen, and the resulting surface metal oxides considerably hinder hydrogen diffusion into the bulk.44,49–52 Hence, the leaching procedure was optimized primarily with respect to the selective removal of these elements. In Table 16 ACS Paragon Plus Environment

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S5, the variation of the leaching temperature is shown. All remaining parameters were kept constant. Based on the above mentioned investigations, an appropriate leaching condition was selected as cNaOH = 9 mol/L, T = 40 °C, S/L = 0.05 g/mL, where S/L stands for the ratio between solid alloy to liquid electrolyte. No pronounced changes in the microstructure of the NaOH-treated samples were observed by SEM. SEM micrographs and EDX elemental maps of the cycled and NaOH-treated samples are shown in Figures S9 and S10 in the Supporting Information. In order to further elucidate the elemental compositions of the leached materials, XPS measurements of the surface and subsurface region were performed. The XPS concentration depth profiles of the materials cycled in pure H2 and of the NaOH-treated material were determined by removing the surface atomic layers by Ar+ ion sputtering for varying times and measuring XPS spectra after each sputtering step. By this method, the chemical composition could be determined as a function of sputtering time. Results are shown in Figure 5.

Figure 5. Atomic concentrations of the elements in V40Fe8Ti32Cr16Ni4 as a function of the sputtering time. Open symbols and dashed lines stand for the cycled material without contact to air or moisture (no leaching), closed symbols stand for the



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leached material. C and O are not included, and the concentrations were normalized to 100 %. For both cycled and leached samples the concentration of Ti was higher in the surface and subsurface region of the sample (around 40 at%) as compared to the bulk and decreased with sputtering depth. After 150 s of sputtering, the concentration of Ti in the cycled sample stayed fairly constant at 25 – 27 at%. The concentration of V was higher in the surface and subsurface region of the cycled sample than in the case of the NaOH-treated sample. After the fourth sputtering step (1200 s), similar V concentrations were measured in both samples. As expected, the concentration depth profiles of Ni were similar for both samples. For Cr and Fe, the atomic concentrations were higher at the surface and subsurface region as a result of the NaOH treatment. For Cr, this result is in agreement with the ICP-OES results. In the case of Fe, this indicates a preferential dissolution of V over Fe, which results in a higher relative concentration of Fe as compared to before the leaching treatment although αFe was > 1. These results indicate that the elemental components of the alloy show different responses to the leaching treatment. In the case of V, the concentration in the surface and subsurface region could be reduced successfully through the leaching procedure. By contrast, the concentration of Ti was not reduced as compared to the cycled sample. This finding is in contrast to the ICP-OES analysis, since Ti ions were reproducibly detected in the electrolyte (9 M NaOH). This would indicate the dissolution of significant amounts of this element during the leaching process. However, it is likely that Ti atoms diffuse from the bulk of the material towards the surface and are enriched as a result of two combined driving forces: On the one hand, it is known that Ti-Ti bonds and bonds between Ti and other metals are relatively weak which facilitates diffusion of Ti atoms within alloys. On the other hand, the gain 18 ACS Paragon Plus Environment

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in free energy for the formation of TiO2 from Ti and O2 and thus the driving force for Ti

oxidation

is

rather

high

( : − 850 ∙ ; ! " : − 550 ∙

; #$ % : − 700 ∙ ).50 Therefore, it is likely that the combination of Ti diffusion from bulk to surface and the strong affinity of Ti to oxygen are responsible for the higher relative concentration of Ti at the surface in spite of the selective dissolution of this element in the electrolyte. As the Ti at the surface and the other metals are dissolved, fresh metal atoms are exposed to the electrolyte. Dissolved oxygen and water molecules adsorb at the surface, and Ti shows a preferential tendency to diffuse to the surface and react with these oxygen species. This is also in line with previous reports on, e.g., the oxidation of TiNi shape memory alloys, where the formation of a thick TiO2 surface layer has been observed upon exposure of the alloys to air or an oxygen atmosphere at elevated temperatures.53,54 As described in the Introduction, it is known that surface segregation occurs at room temperature in the case of LaNi5 hydrogen storage alloys, where La has a pronounced affinity for oxygen and diffuses to the alloy surface. The free energy of the surface is reduced through segregation and preferential oxidation of La.20 The hydrogen storage properties of V40Fe8Ti32Cr20, V40Fe8Ti32Cr16Ni4, and leached V40Fe8Ti32Cr16Ni4 were also investigated. V-Ti-based hydrogen storage alloys exhibit two distinct hydride formation and dissociation plateau regions. In the lower plateau region up to a hydrogen content of ∼1.2–1.5 mass%, the monohydride is formed: M + H → MH1, where M stands for metal and H stands for hydrogen. This region typically shows a very low equilibrium pressure of

Performance Improvement of VâFeâCrâTi Solid State Hydrogen

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