Classifications

• • H—ELECTRICITY
  • H01—ELECTRIC ELEMENTS
  • H01M—PROCESSES OR MEANS, e.g. BATTERIES, FOR THE DIRECT CONVERSION OF CHEMICAL ENERGY INTO ELECTRICAL ENERGY
  • H01M4/00—Electrodes
  • H01M4/02—Electrodes composed of, or comprising, active material
  • H01M4/36—Selection of substances as active materials, active masses, active liquids
  • H01M4/38—Selection of substances as active materials, active masses, active liquids of elements or alloys
  • H01M4/383—Hydrogen absorbing alloys
• • C—CHEMISTRY; METALLURGY
  • C01—INORGANIC CHEMISTRY
  • C01B—NON-METALLIC ELEMENTS; COMPOUNDS THEREOF; METALLOIDS OR COMPOUNDS THEREOF NOT COVERED BY SUBCLASS C01C
  • C01B3/00—Hydrogen; Gaseous mixtures containing hydrogen; Separation of hydrogen from mixtures containing it; Purification of hydrogen
  • C01B3/0005—Reversible uptake of hydrogen by an appropriate medium, i.e. based on physical or chemical sorption phenomena or on reversible chemical reactions, e.g. for hydrogen storage purposes ; Reversible gettering of hydrogen; Reversible uptake of hydrogen by electrodes
  • C01B3/001—Reversible uptake of hydrogen by an appropriate medium, i.e. based on physical or chemical sorption phenomena or on reversible chemical reactions, e.g. for hydrogen storage purposes ; Reversible gettering of hydrogen; Reversible uptake of hydrogen by electrodes characterised by the uptaking medium; Treatment thereof
  • C01B3/0031—Intermetallic compounds; Metal alloys; Treatment thereof
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C16/00—Alloys based on zirconium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C19/00—Alloys based on nickel or cobalt
  • C22C19/03—Alloys based on nickel or cobalt based on nickel
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C30/00—Alloys containing less than 50% by weight of each constituent
• • H—ELECTRICITY
  • H01—ELECTRIC ELEMENTS
  • H01M—PROCESSES OR MEANS, e.g. BATTERIES, FOR THE DIRECT CONVERSION OF CHEMICAL ENERGY INTO ELECTRICAL ENERGY
  • H01M4/00—Electrodes
  • H01M4/02—Electrodes composed of, or comprising, active material
  • H01M4/24—Electrodes for alkaline accumulators
  • H01M4/242—Hydrogen storage electrodes
• • H—ELECTRICITY
  • H01—ELECTRIC ELEMENTS
  • H01M—PROCESSES OR MEANS, e.g. BATTERIES, FOR THE DIRECT CONVERSION OF CHEMICAL ENERGY INTO ELECTRICAL ENERGY
  • H01M10/00—Secondary cells; Manufacture thereof
  • H01M10/34—Gastight accumulators
  • H01M10/345—Gastight metal hydride accumulators
• • H—ELECTRICITY
  • H01—ELECTRIC ELEMENTS
  • H01M—PROCESSES OR MEANS, e.g. BATTERIES, FOR THE DIRECT CONVERSION OF CHEMICAL ENERGY INTO ELECTRICAL ENERGY
  • H01M4/00—Electrodes
  • H01M4/02—Electrodes composed of, or comprising, active material
  • H01M2004/026—Electrodes composed of, or comprising, active material characterised by the polarity
  • H01M2004/027—Negative electrodes
• • H—ELECTRICITY
  • H01—ELECTRIC ELEMENTS
  • H01M—PROCESSES OR MEANS, e.g. BATTERIES, FOR THE DIRECT CONVERSION OF CHEMICAL ENERGY INTO ELECTRICAL ENERGY
  • H01M2220/00—Batteries for particular applications
  • H01M2220/20—Batteries in motive systems, e.g. vehicle, ship, plane
• • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
  • Y02—TECHNOLOGIES OR APPLICATIONS FOR MITIGATION OR ADAPTATION AGAINST CLIMATE CHANGE
  • Y02E—REDUCTION OF GREENHOUSE GAS [GHG] EMISSIONS, RELATED TO ENERGY GENERATION, TRANSMISSION OR DISTRIBUTION
  • Y02E60/00—Enabling technologies; Technologies with a potential or indirect contribution to GHG emissions mitigation
  • Y02E60/10—Energy storage using batteries
• • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
  • Y02—TECHNOLOGIES OR APPLICATIONS FOR MITIGATION OR ADAPTATION AGAINST CLIMATE CHANGE
  • Y02E—REDUCTION OF GREENHOUSE GAS [GHG] EMISSIONS, RELATED TO ENERGY GENERATION, TRANSMISSION OR DISTRIBUTION
  • Y02E60/00—Enabling technologies; Technologies with a potential or indirect contribution to GHG emissions mitigation
  • Y02E60/30—Hydrogen technology
  • Y02E60/32—Hydrogen storage

Definitions

• the present invention relates generally to -metal hydride batteries and more specifically to the negative electrodes thereof. Most specifically, this invention relates to a hydrogen storage material for use in the negative electrodes of a Ni-metal hydride battery.
• the alloys have electrochemical capacities which are higher than predicted by their gaseous capacities at 2 MPa of pressure.
• the hydrogen storage alloy may be selected from alloys of the group consisting of A 2 B, AB, AB 2 , AB 3 , A2B7, AB5 and AB9.
• Vanadium has been regarded as a hydride forming element in the development of multi-phase disordered AB 2 MH alloys.
• the contribution of V to the hydrogen storage properties of AB2 MH alloys was reported previously and can be summarized as follows. Vanadium increases the maximum hydrogen storage capacity of the alloy, but the reversible hydrogen storage capacity decreases due to the increase in hydrogen-metal bond strength.
• V was chosen to be the first modifying element, and the results were very promising: the full electrochemical capacity increased from 204 mAh/g in Tii.5Zr5.5Mjo to 359 mAh/g in Tii. Zr .5V2.5Nt7.s.
• the present invention is a hydrogen storage alloy which has a higher electrochemical hydrogen storage capacity than that predicted by the alloy's gaseous hydrogen storage capacity at 2 MPa.
• the hydrogen storage alloy may have an electrochemical hydrogen storage capacity 5 to 15 times higher than that predicted by the maximum gaseous phase hydrogen storage capacity thereof.
• the hydrogen storage alloy may be selected from alloys of the group consisting of A 2 B, AB, AB2, AB3, A 2 B 7 , AB5 and AB9.
• the hydrogen storage alloy may be elected from the group consisting of: a) Zr(VxNi4.5-x); wherein 0 ⁇ x ⁇ 0.5; and b) Zr(V x Ni3.s-x); wherein 0 ⁇ x ⁇ 0.9.
• x may be: 0.1 ⁇ x ⁇ 0.5; 0.1 ⁇ x ⁇ 0.3; 0.3 ⁇ x ⁇ 0.5; 0.2 ⁇ x ⁇ 0.4. Also, x may be any of 0.1 ; 0.2; 0.3; 0.4; or 0.5.
• the hydrogen storage alloy may further include one or more elements selected from the group consisting Mn, Al, Co, and Sn in an amount sufficient enough to enhance one or both of the discharge capacity and the surface exchange current density versus the base alloy.
• the hydrogen storage alloy When the hydrogen storage alloy has the formula: Zr(V x Ni4.5-x), it may have one or more properties such as: 1) a bulk proton diffusion coefficient greater than 4 x ] 0 "10 cm 2 s "! ; 2) a high rate dischargeability of at least 75%; 3) an open circuit voltage of at least 1 .25 volts; and an exchange current of at least 24 mA g : i .
• the present invention further includes a negative electrode for a Ni-metal hydride battery formed using the inventive alloys and a Ni-metal hydride battery formed using said electrode.
• Fig. 1 is a plot of the XRD patterns using Cu-K as the radiation source for alloys YC#1 to YC#6;
• Fig. 2 plots the unit cell volume of the m-Zr2Ni7 phase as a function of V -content in the alloy
• Fig. 3 plots the phase abundances as functions of V-content in the alloy
• Figs. 4a-4f are SEM back-scattering electron images for alloys YC#1 (a), YC#2 (b), YC#3 (c), YC#4 (d), YC#5 (e), and YC#6 i f), respectively;
• Figs. 5a-5b plot the PCT isotherms measured at 30°C for alloys YC#1 - YC#3 (5a) and YC#4 - YC#6 (5b);
• Fig. 6a plots the half-ceil discharge capacities of the six alloys measured at 4 mA g-1 versus cycle number during the first 13 cycles;
• Fig. 6b plots the high-rate dischargeabilities of the si alloys versus cycle number during the first 13 cycles;
• Fig. 7 plots the open circuit voltage vs. pressure at the mid-point of PCX desorption isotherm measured at 30°C from two series of prior art off-stoichiometric MH alloys (AB2 and AB5);
• J Fig. 8 plots the full discharge capacities at the 10th cycle (open symbol) and open circuit voltage (solid symbol) as functions of V-content in the alloy for the six alloys YC#1 -YC#6;
• Fig. 10a is a plot of the X D patterns using Cu-K as the radiation source for alloys YC#7 to YC#1 1 ;
• Fig. 10b is a plot of the XRD patterns using Cu-K as the radiation source for alloys YC#12 to YC#16;
• Fig. 1 1 is photomicrograph of sample YC#12, and is exemplar ⁇ ' of the photomicrographs of al l of the samples YC#7 - YC#16.
• the present inventors have discovered hydrogen storage alloys that have electrochemical hydrogen storage capacities which are higher than predicted by their respective gaseous hydrogen storage capacities at 2 Mpa of pressure.
• the hydrogen storage alloys may have electrochemical hydrogen storage capacities 5 to 15 times higher than that predicted by the maximum gaseous phase hydrogen storage capacity thereof.
• the hydrogen storage alloy may be any alloy selected from alloys of the group consisting of A 2 B, AB, AB 2 , AB3, ⁇ : ⁇ . AB 5 and A B9.
• the inventors believe that the electrochemical discharge capacity is higher than the capacity obtained from gaseous phase measurement due to the synergetic effects of secondary phases present in the present, un-annealed alloys. While not wishing to be bound by theory, the inventors believe that the secondary phases in the present alloys act as catalysts to reduce the hydrogen equilibrium pressure in the electrochemical environment and increase the storage capacity.
• the term "synergetic effect” is used herein to describe the increase in discharge capacity or high rate dischargeability (HRD) of the main phase in the presence of secondary phases.
• HRD high rate dischargeability
• the synergetic effect arises as a result of the multi-phase nature, which provides various properties that together contribute positively to the overall performance.
• the presence of secondary phases offers more catalytic sites in the microstructure for gaseous phase and/or electrochemical hydrogen storage reactions.
• the secondary phases may have too high of a hydrogen equilibrium pressure and they may not absorb any considerable amount of hydrogen; however, they may act as a catalyst for hydrogen storage of the main phase.
• the abundance of the secondary phase is not as important as the interface area affected by the synergetic effect.
• both the interface area and the penetration depth of the synergetic effect are cmcial for maximizing the advantages of the present invention, such as higher storage capacity, higher bulk diffusion, and other electrochemical properties.
• the penetration depth may be estimated by dividing the improvement in various properties by the interface area from scanning electron micrographs.
• the present invention comprises the use of V as a modifying element to improve the electrochemical properties of ZrNi5 alloy.
• V a modifying element to improve the electrochemical properties of ZrNi5 alloy.
• the main phase(s) of the alloy evolves from ZrNis and cubic Zr 2 Ni 7 to monoclinic Zr 2 Ni 7 , ZrNis and ZrNig, and then finally to monoclinic Zr 2 Ni 7 only with increases in V-content.
• the secondary phase(s) evolves from monoclinic Zr 2 Ni 7 and ZrN3 ⁇ 4 to cubic Zr 2 Ni? and VNi 3 and then to VN12.
• PCT results show incomplete hydriding using the current set-up (up to 1.1 MPa), low maximum gaseous phase hydrogen storage capacities ( ⁇ 0.075 vvt.%, 0.05 H/M), and large hysteresis.
• the maximum gaseous phase storage capacity decreases, in general, with the increase in V-content.
• the highest bulk diffusion coefficient obtained is ⁇ ,06 ⁇ 10 ⁇ 0 era 2 s "1 from the base alloy ZrNks, which is more than double of the coefficient for the currently used AB5 alloy (2.55 x 10 "i o cm 2 s ⁇ ! ).
• the discharge capacity ( ⁇ 177 rriAh g -1 ) and the surface exchange current density are lower than the commercially used AB5 alloy, these properties can be further optimized by introducing other modifying elements, such as Mn, Al, and Co.
• Arc melting was performed under a continuous argon flow with a non-consumable tungsten electrode and a water-cooled copper tray. Before each run, a piece of sacrificial titanium underwent a few melting-cooling cycles to reduce the residual oxygen concentration in the system. Each 12g ingot was re-melted and turned over a few times to ensure uniformity in chemical composition.
• each sample was examined by a Varian Liberty 100 inductively-coupled plasma (ICP) system, A Philips X'Pert Pro x-ray diffractometer (XRD) was used to study the microstructure, and a JEOL-JSM6320F scanning electron microscope (SEM) with energy dispersive spectroscopy (EDS) capability was used to study the phase distribution and composition.
• the gaseous phase hydrogen storage characteristics for each sample were measured using a Suzuki-Shokan multi-channel pressure- concentration-temperature (PCT) system.
• PCT Suzuki-Shokan multi-channel pressure- concentration-temperature
• compositions determined by ICP are very close to the design values.
• the ingots were not annealed in order to preserve the secondary phases, which may be beneficial to the electrochemical properties.
• Formulas in the format of Zr( V, Nij4.5 and associated formula weights are also included in Table I .
• Fig. 1 is a plot of the XRD patterns using Cu-K as the radiation source for alloys YC#1 to #6.
• the vertical line is to illustrate the shifting of the ZrNi9 and VNi2 peaks to lower angles.
• Five structures can be identified: a monoclinic Zr 2 Ni 7 (m-Zr 2 Ni 7 ) (reference symbol o), a cubic ⁇ 2 ⁇ 17 (c-Zr 2 Ni 7 ) (reference symbol ⁇ ), a cubic ZrNis (reference symbol V), a cubic ZrNig (reference symbol ⁇ ), and an orihorhombic VNi 2 phase (reference symbol II).
• the second structure a metastable structure of 2Ni 7 , is cubic with lattice constant a ⁇ 6.68A, An orihorhombic ZraNi? phase has been reported previously but was not observed in the current study, Hf 2 Co? is a similar alloy that contains this stable orthorhombic phase,
• the third structure, a ZrNis cubic structure is AuBes-type.
• the fourth structure the ZrNi phase
• the fourth structure does not exist in the Zr-Ni binary phase diagram and has not been reported before.
• the fifth structure an orthorhombic VN12 phase with a MoPt 2 structure, has a diffraction pattern with peaks overlapping with those of a simple cubic structure, such as ZrNis, with the major difference being a splitting of the (1 30) and (002) reflections near 50°.
• VNi reference symbol J
• the unit ceil volume of each phase increases as the V-content in the alloy increases except for the ni-Zr 2 Ni 7 phase in the alloy with very low V-content (YC#2), Considering that Zr is larger than V, and V is larger than Ni, the increase in unit ceil volume indicates that V occupies the B-site and replaces Ni.
• the unit cell volume of m-Zr 2 Ni? is plotted against the average V-content in the alloy in Fig, 2. In the m-ZraNi?
• the phase abundances analyzed by jade 9 software are listed in Table 2.
• Fig. 3 plots the phase abundances as functions of V-content in the alloy.
• the V-free YC#1 is composed of mainly c-Zr 2 Ni 7 (symbol o) and ZrNis (symbol ⁇ ) with m-Zr 2 Ni 7 (symbol ⁇ ) and ZrNi9 (symbol ⁇ ) as the secondary phases.
• the main phase first shifts to and then to ni-Zr 2 Ni 7 only.
• the secondary phase first changes into c-Zr 2 Ni 7 and then to VN12 (symbol ⁇ ).
• the phase abundances of alloys YC#4, 5, and 6 are very similar at about 70% m -Zs-jXi- and 30% VNi 2 .
• the ZrNig secondary phase can be found within the ZrNis main phase in the shape of a liquid droplet (Fig. 4a-4), the ZrNi secondary phase within the c-Zr 2 Ni 7 phase is manifested as fine crystals with well- defined edges (Fig. 4a-5).
• YC#2 three main phases can be found: Zr 2 Ni7 (Fig. 4b- 1), ZrNis (Fig. 4b-2), and ZrNi 9 (Fig. 4b-3). Within the Zr 2 Ni 7 phase, some areas with slightly darker contrast can be identified.
• the majority of the Zr 2 Ni 7 phase with slightly brighter contrast can be designated as the m ⁇ Zr 2 Ni 7 phase, with the darker regions being the c-Zr 2 Ni 7 phase.
• the major secondary phase with darker contrast (Fig. 4b-4) compared to the main phases is located between the ZrNis and ZrNig phases. This phase has a similar Ni-content to the main ZrNiq phase; however, its V-content is higher than the Zr-content. There must be some V occupying the Zr-site in this case; therefore, this phase is designated as the ⁇ - ⁇ phase. A sharp needle-like inclusion was found in the Zr 2 Ni 7 matrix (Fig. 4b-5).
• this inclusion has a very small amount of V and can therefore be assigned as the Zr3 ⁇ 4Nis phase, which does not exist in the Zr-Ni binary phase diagram.
• Another secondary phase the one with the darkest contrast, has a very' small amount of Zr (Fig. 4b ⁇ 6) and is assigned to be the VNi phase according to the stoicliiometry, which has a XRD diffraction pattern very close to that of TiNi3 ⁇ 4s.
• the brightest contrast comes from the main phase, m-Zr 2 Ni 7 (Fig. 4c- 1). The slightly darker region (Fig. 4c-2) and the sharp crystal (Fig.
• the microstructures of the last three alloys are very similar: Zr 2 Ni 7 as the matrix and VNi 2 as the secondary phase with occasional Zr() 2 inclusions.
• the V-content in the Zr 2 Ni? phase increases slightly from 0.7 to 1.1 and then to 1.6 at.% while the V-content in the VNi 2 phase increases from 29.2 to 31.2 and then to 37.2 at.% in alloys YC#4, 5, and 6, respectively.
• the changes in Zr- content in these two phases are very small in the last three alloys.
• the equivalent gaseous phase plateau pressures are listed in Table 5 and range between 0.032 and 1.126 MPa.
• the plateau pressures of the first five alloys in the electrochemical system are lower than the highest pressure employed in the PCT apparatus (1 .1 MPa). Therefore, the electrochemical environment is able to reduce the hydrogen storage plateau pressure and consequently increases the storage capacity.
• the second method of estimating the equivalent gaseous phase equilibrium hydrogen pressure w r as considered due to the fact that most of the disordered MH alloys lack well- defined plateaus in the a-to-b transition in the PCT isotherm. Instead of the Nernst equation, an empirical relationship between the mid-point pressure in the PCT desorption isotherm and OCV (Fig.
• [0036J Fig. 8 plots the full discharge capacities at the 10th cycle (open symbol) and open circuit voltage (solid symbol) as functions of V-content in the alloy for the six alloys YC#1-YC#6.
• OCV increases as the V-content increases except for alloy YC#2.
• the drop in OCV and the boost in discharge capacity in YC#2 may be related to the shrinkage in unit cell volume of the m-Zr 2 Ni.7 phase as shown in Fig. 2, With the increase in the amount of V substituting Ni, the average strength of metal -hydrogen bond increases, and higher discharge capacity is expected and observed.
• OCV which is closely related to the equilibrium hydrogen pressure, is expected to decrease with the increase in metal-hydrogen bond strength, which is not seen in the current study.
• the OCV was altered by the electrochemical environment and is lower than the value expected from the gaseous phase PCX analysis.
• the increase in OCV with the increase in V-content indicates that the charge/discharge characteristics in this multi-phase alloy system are strongly influenced by either the surface modification due to the reaction with KOFI or by the synergetic effect from the catalytic secondary phases as seen in multi-phase AB 2 MH alloy systems.
• the discrepancy between the gaseous phase and electrochemical behaviors is further highlighted when the discharge capacity is plotted against the maximum gaseous phase hydrogen storage capacity.
• Fig. 9 plots the measured electrochemical discharge capacity vs.
• the electrochemical capacity correlates very well to the abundances of several phases, such as both m- and c-Zr 2 Ni 7 phases, the ZrNi.5 phase, and the VNi2 phase.
• the correlation between the electrochemical capacity and the average V-content is the most significant. With the increase in V-content, the average proton affinity of the alloy increases and contributes to a higher electrochemical storage capacity, which is in opposition with the finding from the gaseous phase study.
• Figs. 10a and 10b The XRD patterns of the ten alloys are shown in Figs. 10a and 10b.
• Four structures can be identified: a monoclinic Zr2Ni 7 (m-Zr 2 Ni 7 symbol o). a cubic Zr 2 Ni 7 (c-Zr2Ni 7 symbol ⁇ ), a hexagonal ZrNis phase (symbol V) and a cubic ZrNis phase (symbol ⁇ ).
• the fourth structure a ZrNis cubic structure, is AuBes-type. Its reported lattice constant a varies slightly among different groups, ranging from 6,702 to 6.683 A. Lattice constants and phase abundances obtained from XRD are listed in Table 8.
• microstructures for this series of alloys were studied using SEM, and a back-scattering electron image (BEI) of sample YC#12, which is shown in Fig, 11.
• BEI back-scattering electron image
• This figure is exemplary of the micrographs of all of the samples. Clear phase segregation can be seen from the micrograph. Two phases of ⁇ 2 ⁇ 17 can be identified (spots 1 and 2) with slightly different in contrast and V -content. Without an in-situ electron backscattering diffraction pattern, we cannot assign crystal structures (c- or m-) to these two phases.
• the OCV increased with increasing V-context except for YC#08.
• the addition of V in the alloy is supposed to increase the stability of the hydride by increasing the size of the hydrogen occupation site and decreasing the electronegativity. In this case, however, the equivalent hydrogen pressure increases (less stable hydride) with the increase in the V-content.
• One possible explanation is due to the reduction in synergetic effect from the reduced secondary phase amount as the V-content increased.
• the half-cell HRD of each alloy defined as the ratio of the discharge capacity measured at 50 mA g ; 1 to that measured at 4 mA g ⁇ 1 , at the 10 th cycle are also listed in Table 9, HRD increased as the V-content in the alloy increased.
• the major differences between the secondary phases in AB 2 alloys and Zr Ni 7 MH alloys are the abundance and distribution thereof.
• the secondary phases (mainly Zr 7 Niio and Zi ⁇ Niii) in AB 2 MH alloy are less abundant and more finely distributed, which causes less resistance to hydrogen diffusion in the bulk.
• D values are similar to those obtained from the ZrV Ni4.5-x alloys of example 1 above and are much higher than those measured in other MH alloy systems, such as AB 2 (9.7* 10 ⁇ u cm 2 s “3 ), AB 5 (2.55* iO " 10 cm 2 s “1 ), La-A 2 B 7 (3.08*10 ⁇ 10 cm 2 s _1 ), and Nd-A 2 B 7 (1 .l4x l0 ⁇ 10 cm 2 s ⁇ r ).
• I 0 decreased with increasing V-content.
• These I 0 values are lower other MH alloys such as AB 2 , A 2 B 7 , and ABs MH alloys. Further impro vement in the surface reaction needs to be performed with substitutions that will increase the surface area and/or catalytic properties.

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