JP2017528596A - Laves phase related BCC metal hydride alloys and their activation for electrochemical applications - Google Patents

Abstract

Laves phase related BCC metal hydride alloys have traditionally limited electrochemical capacity. New examples of these alloys useful as electrode active materials are provided. A method for activating such an alloy is also provided. The alloy has the formula I: TiwVxCryMz (I) (where w + x + y + z = 1, 0.1 ≦ w ≦ 0.6, 0.1 <x <0.6, 0.01 <y <0.6, M is selected from the group consisting of B, Al, Si, Sn, and one or more transition metals, and achieves a discharge capacity of 350 mAn / g or more over 10 cycles.

Description

Cross-reference to related applications This application is filed on US Patent Application Publication No. 14 / 340,913 filed July 25, 2014 (US Patent Application No: 14 / 340,913) and July 25, 2014. Claims priority based on filed US Patent Application No. 14 / 340,959 (US Patent Application No: 14 / 340,959), the contents of each of which are incorporated herein by reference. It is.

TECHNICAL FIELD The present invention relates to an alloy material and a manufacturing method thereof. In particular, the present invention relates to a metal hydride alloy material capable of absorbing and releasing hydrogen. An activated metal hydride alloy having a Laves phase related body centered cubic (BCC) structure with inherent electrochemical properties (including high capacity) for electrochemical applications is provided.

BACKGROUND Certain metal hydride (MH) alloy materials can absorb and release hydrogen. These materials can be used as hydrogen storage media for fuel cells and metal hydride batteries including nickel / metal hydride (Ni / MH) and metal hydride / air battery systems and / or as electrode materials. However, due to the limited mass energy density (<110 Wh kg ⁻¹ ), current Ni / MH batteries are deprived of their share in the portable electronics and battery-powered electric vehicle market by lightweight Li-ion technology. Thus, next generation Ni / MH batteries are designed to improve two main goals: increasing energy density and reducing costs.

When a potential is applied between the cathode and the MH anode in the MH cell, the negative electrode material (M) is charged due to the electrochemical absorption of hydrogen and the electrochemical generation of hydroxide ions that generate MH. When discharged, the stored hydrogen is released to form water molecules, releasing electrons. The reaction that occurs at the positive electrode of the Ni / MH cell is also reversible. Most Ni / MH cells use a nickel hydroxide positive electrode. The following charge / discharge reactions occur at the nickel hydroxide positive electrode.

  In Ni / MH cells having a nickel hydroxide positive electrode and a hydrogen storage negative electrode, the electrodes are typically separated by a nonwoven, felt, nylon or polypropylene separator. The electrolytic solution is usually an alkaline aqueous electrolytic solution, for example, 20 to 45% by mass of potassium hydroxide.

One particular group of MH materials useful in Ni / MH battery systems is known as the AB _x class of materials with respect to the crystal arrangement occupied by the membrane components. AB _x type materials are disclosed, for example, in US Pat. No. 5,536,591 (US Pat. No. 5,536,591) and US Pat. No. 6,210,498 (US Pat. No. 6,210,498). Examples of such materials include, but are not limited to, modified LaNi ₅ type (AB ₅ ) and Laves phase-based active material (AB ₂ ). These materials reversibly form hydrides to store hydrogen. Such materials utilize a Ti-Zr-Ni based composition, where at least Ti, Zr, Ni is at least one modification from the group of Cr, Mn, Co, V and Al. Exists with the agent. The material is a multiphase material and may include, but is not limited to, one or more Laves phase crystal structures and other non-Laves secondary phases. Because current AB ₅ alloys have a capacity of about 320 mAh g ^-1 and Laves phase AB ₂ has a capacity up to 440 mAh g ^-1 , these are the most promising alloy alternatives and have high rate discharge performance (HRD), having a good balance among cycle life, charge retention, activation, self-discharge, and applicable temperature range.

Rare earth (RE) magnesium based AB ₃ type or A ₂ B ₇ type MH alloys are promising in part to replace the AB ₅ MH alloys currently used as negative electrodes for Ni / MH batteries due to their high capacity Is a good candidate. Most of the RE-Mg-Ni MH alloys are based only on La as a rare earth metal, but several Nd-only A ₂ B ₇ (AB ₃ ) alloys have been recently reported. In these materials, the AB _3.5 stoichiometry is believed to provide the best overall balance between storage capacity, activation, HRD, charge retention, and cycle stability. The pressure-concentration-temperature (PCT) isotherm of the Nd-only A ₂ B ₇ alloy shows a very sharp take-off angle in the α phase [K. Young, et al., Alloys Compd. 2010 506: 831], a relatively high voltage can be maintained in a low charge state. Compared to the commercially available AB ₅ MH alloy, Nd-only A ₂ B ₇ has a higher utilization of the positive electrode during storage at 60 ° C. and does not increase much resistance, but also has a greater capacity degradation during cycling. [K. Young et al., Int. J. Hydrogen Energy, 2012; 37: 9882]. Another problem with the known A ₂ B ₇ alloys is that the HRD is inferior compared to conventional AB ₅ alloy systems due to the low Ni content in the chemical composition of the alloys.

Other AB _x materials include Laves phase-related body-centered cubic (BCC) materials, which, by way of example, are a group of MH alloys having a two-phase microstructure including a BCC phase and a Laves phase, which conventionally exist as C14 It is. These materials are based on the theoretical electrochemical capacity of 938 mAh g ⁻¹ of a Ti—V—Cr alloy with a complete BCC structure, but unfortunately have very poor electrochemical properties. The Laves phase itself with a similar chemical composition is added to the BCC material. These Laves phase related BCC materials exhibit high density phase boundaries and can combine higher BCC hydrogen storage capacity, good hydrogen absorption kinetics and relatively high C14 phase surface catalytic activity. Much work has been done to optimize these materials. Young et al., Int. J. Hydrogen Energy, http://dx.doi.org/10.1016/j.ijhydene.2014.01.134 (press article) describes a family of these materials with a wide range of BCC / C14 ratios. Studies have been described. These results clearly indicate that these materials have many desirable properties, but the electrochemical discharge capacity produced by these materials does not exceed 175 mAh / g.

  Thus, there is a need for improved hydrogen storage materials. As described herein below, the present invention addresses these needs by providing an activated Laves phase-related BCC metal hydride alloy that exhibits greatly improved electrochemical properties for the first time. . These and other advantages of the invention will be apparent from the following drawings, discussion, and description.

SUMMARY OF THE INVENTION The following summary of the invention is provided in order to facilitate an understanding of some of the innovative features unique to the present invention and is not intended to be a complete description. A full appreciation of the various aspects of the invention can be gained by taking the entire specification, claims, drawings, and abstract as a whole.

  As explained, the alloy material not only has an improved capacity over conventional alloys of similar composition, but also has a significantly improved cycle life at high capacity. Some conventional materials are capable of high capacity, but this capacity decreases rapidly in 1-5 cycles. Due to the improved cycle life in Laves phase related BCC metal hydrides, the capacity gradually decreases. The Laves phase related BCC metal hydride alloy provided herein solves the capacity reduction problem and significantly improves capacity over more cycles by adjusting the ratio of Ti to Cr in the system. Demonstrate.

Formula I below:
_{Ti w V x Cr y M z} (I)
(Where, w + x + y + z = 1, 0.1 ≦ w ≦ 0.6, 0.1 ≦ x ≦ 0.6, 0.01 ≦ y ≦ 0.6, and M is B, Al, Si, Sn, and Selected from the group consisting of transition metals)
A Laves phase related BCC metal hydride is provided, wherein the metal hydride alloy has a capacity of greater than 350 mAh / g in cycle 10. Some embodiments have a capacity of 400 mAh / g or greater, optionally 420 mAh / g or greater, in cycle 10. The alloy optionally contains less than 24% C14 phase. In some embodiments, the alloy is primarily a combination of a BCC phase and a Laves phase, wherein the BCC phase is greater than 5% and less than 95%, and the Laves phase is greater than 5% and less than 95%. Of abundance. Optionally, the alloy comprises a BCC phase crystallite size of less than 400 angstroms, optionally less than 200 angstroms. In some embodiments, the B / A ratio is 1.20 to 1.31, optionally 1.20 to 1.30. Optionally, the ratio x / y is 1-3. Some embodiments of any of the above have the following formula II:
Ti _{13.6 + x} Zr _2.1 V ₄₄ Cr _13.2-x M _27.1 (II)
(Wherein x is greater than 0 and not greater than 12 and M is a combination of Mn, Fe, Co, Ni, and Al)
Including the composition. The alloy of formula II optionally has x of 2, 4, 6, 8, 10 or 12. Some embodiments of any of the above have the formula III:
Ti _{0.4 + x / 6} Zr _{0.6-x / 6} Mn _0.44 Ni _1.0 Al _0.02 Co _0.09 (VCr _0.3 Fe _0.063 ) _x (III)
(X is 0.7 to 2.8)
Including the composition.

  The provided alloys and their equivalents represent excellent materials for use in cell or battery system anodes.

Also provided is a method that is useful for generating an increased stable capacity over conventional methods. Thus, the following formula I:
_{Ti w V x Cr y M z} (I)
(Where, w + x + y + z = 1, 0.1 ≦ w ≦ 0.6, 0.1 ≦ x ≦ 0.6, 0.01 ≦ y ≦ 0.6, and M is B, Al, Si, Sn, and Selected from the group consisting of transition metals)
In the method of activating the Laves phase related BCC metal hydride alloy, the method is achieved by subjecting the Laves phase related BCC metal hydride alloy to an atmosphere containing hydrogen at a hydrogenation pressure, and a previous activation step. Cooling the alloy during the process of producing an activated metal hydride alloy having a capacity exceeding 200 mAh / g, a level that was not present. In some embodiments, the cooling step has a maximum activation temperature of 300 ° C. or less. The atmosphere is optionally a hydrogenation pressure of 1.4 megapascals or higher, optionally 6 megapascals or higher. The method produces an activated metal hydroxide alloy having a capacity of optionally 300 mAh / g or higher, optionally 350 mAh / g or higher, optionally 400 mAh / g or higher, optionally 450 mAh / g or higher. In some embodiments, the activated metal hydride alloy has less than 24% C14 phase. In some embodiments, the activated metal hydride alloy is primarily a combination of a BCC phase and a Laves phase, where the BCC phase is greater than 5% and less than 95%, and the Laves phase is greater than 5%. And less than 95%. Optionally, the activated metal hydride alloy comprises a BCC phase crystallite size of less than 400 angstroms. In any of the above or a combination thereof, the Laves phase related BCC metal hydride alloy is optionally represented by the following formula II:
Ti _{0.4 + x / 6} Zr _{0.6-x / 6} Mn _0.44 Ni _1.0 Al _0.02 Co _0.09 (VCr _0.3 Fe _0.063 ) _x (II)
(Wherein x is 0.7 to 2.8)
belongs to. In any of the above or combinations thereof, the Laves phase-related BCC metal hydride alloy is optionally represented by the formula III:
Ti _{0.4 + x / 6} Zr _{0.6-x / 6} Mn _0.44 Ni _1.0 Al _0.02 Co _0.09 (VCr _0.3 Fe _0.063 ) _x (III)
(Wherein x is 0.7 to 2.8)
belongs to.

1A shows the alloy phase distribution seen in the SEM image of the P8 hydrogen storage alloy, FIG. 1B shows the alloy phase distribution seen in the SEM image of the P9 hydrogen storage alloy, and FIG. 1C shows the SEM of the P10 hydrogen storage alloy. FIG. 1D shows the alloy phase distribution seen in the SEM image of the P11 hydrogen storage alloy, and FIG. 1E shows the alloy phase distribution seen in the SEM image of the P12 hydrogen storage alloy. 1F shows the alloy phase distribution seen in the SEM image of the P13 hydrogen storage alloy, and FIG. 1G shows the alloy phase distribution seen in the SEM image of the P14 hydrogen storage alloy. FIG. 2A shows the main constituent elements in the C14 phase and their metal ratio as a function of Ti content in the alloy design, and FIG. 2B shows the main composition in the BCC phase as a function of Ti content in the alloy design. The element and its metal ratio are shown. FIG. 3A shows the microstructure of a hydrogen storage alloy activated to provide improved electrochemical properties and shows two major phases, C14 and BCC, and FIG. 3B shows improved electrochemical The full width at half maximum (FWHM) of the BCC (110) peak from the control and hydrogen storage alloys activated to provide properties, and the crystallites of the alloys activated by the exemplary process described herein Demonstrate size reduction. FIG. 4 shows a schematic diagram of the C14 unit cell of various alloys composed of alternating A2B and B3 layers stacked along the c-axis, with larger A atoms occupying 4f sites and smaller B atoms. Occupies 2a site (on the A2B layer) and 6h site (on the B3 layer). FIG. 5 shows the FWHM of the BCC (110) peak from a control and hydrogen storage alloy activated to provide improved electrochemical properties, activated by the exemplary process described herein. Demonstrates a reduction in the crystallite size of the finished alloy. FIG. 6 shows the gas phase hydrogen storage properties of various alloy materials formed by the exemplary process described herein. FIG. 7A shows the gas phase hydrogen storage characteristics of various alloy materials, and FIG. 7B shows the gas phase hydrogen storage characteristics of various alloy materials. FIG. 8A shows the half-cell discharge capacity measured at 4 mA / g for the first 13 cycles, and FIG. 8B shows the high rate discharge performance (HRD) for the first 13 cycles. FIG. 9 shows the hydrogen storage capacity converted from gas phase hydrogen storage and measured electrochemically as a function of alloy number. FIG. 10 shows the diffusion coefficient and C14 phase crystallite size as a function of alloy number, and this trend shows that hydrogen diffusion is easier in alloys with smaller C14 phase crystallids.

DETAILED DESCRIPTION OF THE INVENTION The following description of specific embodiments is merely exemplary in nature and is in no way intended to limit the scope, use, or use of the invention, which, of course, varies. Can do. The present invention is described with reference to non-limiting definitions and terms contained herein. These definitions and terms are not intended to serve to limit the scope or practice of the invention, but are presented for illustrative and descriptive purposes only. Although these methods or compositions are described as a sequence of individual steps or as using specific materials, these steps or materials are interchangeable and the description of the invention is easily understood by those skilled in the art. It is understood that a plurality of elements or steps provided in as many ways as possible can be included.

  When an element is described as being “on” another element, it is understood that this can exist directly on top of other elements, or there can be intervening elements between them. Let's go. In contrast, when an element is described as “directly above” another element, there are no intervening elements present.

  Although terms such as “first”, “second”, “third”, etc. may be used herein to describe various elements, components, regions, layers, and / or sections, It will be understood that elements, components, regions, layers, and / or sections of this document should not be limited by these terms. These terms are only used to distinguish one element, component, region, layer or section from another element, component, region, layer or section. Accordingly, the “first element”, “component”, “region”, “layer” or “section” described below may be used to describe the second (or other) without departing from the teachings herein. Can be referred to as an element, component, region, layer, or section.

  The terminology used herein is for the purpose of describing particular embodiments only and is not intended to be limiting. As used herein, the singular forms “a”, “an”, and “the” include plural forms including “at least one” unless the content clearly dictates otherwise. Is intended. “Or” means “and / or”. As used herein, the term “and / or” includes any and all combinations of one or more of the associated listed items. As used herein, the terms “comprises” and / or “comprising” or “includes” and / or “including” include the described features, regions, Specifies the presence of an integer, process, operation, element, and / or component, but the presence of one or more other features, regions, integers, processes, operations, elements, components, and / or groups thereof It will be further understood that this does not exclude additions. The term “or combinations thereof” means a combination comprising at least one of the aforementioned elements.

  All terms (including technical and scientific terms) used herein have the same meaning as commonly understood by one of ordinary skill in the art to which this disclosure belongs, unless otherwise defined. Also, terms as defined in commonly used dictionaries should be construed as having a meaning consistent with the meaning in the context of the related art and this disclosure, and are not clearly defined herein. It will be understood that, as far as it is not interpreted in an idealized or excessively formal sense.

  Hydrogen storage alloys with Laves phase related BCC structure have been studied for quite some time to identify how to promote the synergistic effect between the C14 phase and the BCC phase of the system. Prior studies replacing the elements at the A and B sites have been conducted on a number of mixed phase alloys, some of which have been found to increase or decrease the abundance of the C14 phase. Composition improvements continue as a means of improving both the gas storage and electrochemical hydrogen storage properties of these alloys. These efforts have been linked to some success, but often confused the conclusions and achieving capacities exceeding 200 mAh / g remained difficult. The alloys provided herein represent a simple and sophisticated solution to these problems by showing hydrogen storage alloys of Laves phase related BCC structural materials that exhibit excellent electrochemical properties.

Hydrogen storage alloys having Laves phase-related BCC structures that exhibit excellent electrochemical properties are provided, which are surprisingly superior to conventional materials having similar compositions. Formula I below:
_{Ti w V x Cr y M z} (I)
(Where, w + x + y + z = 1, 0.1 ≦ w ≦ 0.6, 0.1 ≦ x ≦ 0.6, 0.01 ≦ y ≦ 0.6, and M is B, Al, Si, Sn, and Selected from the group consisting of one or more transition metals)
A Laves phase related BCC metal hydride alloy having the following composition is provided: The alloy is activated by certain methods to promote the formation of increased BCC phase and limit the AB ₂ phase in the resulting material. The result is an activated metal hydride alloy with improved electrochemical properties including a capacity of greater than 200 mAh / g in cycle 10 and optionally a capacity of 350 mAh / g or more.

Optionally, the following formula II:
Ti _{13.6 + x} Zr _2.1 V ₄₄ Cr _13.2-x M _27.1 (II)
(Wherein x is a value exceeding 0 and 12 or less, and M is a combination of Mn, Fe, Co, Ni, and Al)
A Laves-related BCC metal hydride alloy having the following composition is provided: Optionally, x is 1, 2, 3, 4, 5, 6, 7, 8, 9, 10, 11 or 12, or any value greater than 0 and less than 12 including non-integers. In some embodiments, x is 2, 4, 6, 8, 10, or 12. Optionally, x is 2 or 4.

Optionally, the Laves phase related BCC metal hydride alloy can have the following formula III:
Ti _{0.4 + x / 6} Zr _{0.6-x / 6} Mn _0.44 Ni _1.0 Al _0.02 Co _0.09 (VCr _0.3 Fe _0.063 ) _x (III)
(X is 0.7 to 2.8)
Of the composition.

  In some embodiments, the Laves phase related BCC metal hydride alloy is well over 200 mAh / g, optionally 220 mAh / g or more, 240 mAh / g or more, 260 mAh / g or more, 280 mAh / g or more, 300 mAh / g or more. 310 mAh / g or more, 320 mAh / g or more, 330 mAh / g or more, 340 mAh / g or more, 350 mAh / g or more, 360 mAh / g or more, 370 mAh / g or more, 380 mAh / g or more, 390 mAh / g or more, 400 mAh / g or more 410 mAh / g or more, 420 mAh / g or more, 430 mAh / g or more, 440 mAh / g or more, 450 mAh / g or more. Optionally, the metal hydride alloy comprises a capacity of 200 to 450 mAh / g. Optionally, the activated metal hydride alloy comprises a capacity of 200-450 mAh / g. Optionally, the metal hydride alloy comprises a capacity of 300-450 mAh / g. Optionally, the activated metal hydride alloy comprises a capacity of 300-380 mAh / g. Optionally, the metal hydride alloy comprises a capacity of 350-450 mAh / g. Optionally, the metal hydride alloy comprises a capacity of 400-450 mAh / g. Optionally, the metal hydride alloy comprises a capacity of 400-420 mAh / g. In some aspects, any of the aforementioned capacities are optionally 3, 4, 5, 6, 7, 8, 9, 10, 11, 12, 13, 14, 15 in two or more cycles. , 16, 17, 18, 19, 20, or more cycles. Optionally, the metal hydride alloy has a capacity of 350 mAh / g or more in cycle 10, optionally 400 mAh / g or more in cycle 10, and optionally 420 mAh / g or more in cycle 10.

  Depending on its composition and lack of oxidation, the physical structure of the material promotes the superior electrochemical properties of metal hydride alloys compared to conventional alloy materials with similar chemical compositions. Metal hydride alloys are optionally formed primarily from BCC and Laves phase structures. Without being limited to any particular theory, it is believed that the structural advantage is in the BCC phase, and the Laves phase increases the synergistic effect caused by the presence of two phases. Thus, a metal hydride alloy optionally produced by the methods disclosed herein optionally has a BCC phase in an abundance of greater than 5% and less than 95%, greater than 5% and Having a Laves phase in an abundance of less than 95%, where the combination of BCC phase and Laves phase exceeds 50% of the material structure. Optionally, the BCC phase is 10% to 95%, 20% to 95%, 30% to 95%, 40% to 95%, 50% to 95%, 60% to 95%, 70% to 95%, or 80 % -95%, optionally even in such cases, the Laves phase exceeds 5%. Optionally, the Laves phase is 10% to 95%, 20% to 95%, 30% to 95%, 40% to 95%, 50% to 95%, 60% to 95%, 70% to 95%, or 80% to 95%, optionally even in such cases, the BCC phase exceeds 5%. In some embodiments, the composition comprises less than 30%, optionally less than 25%, optionally less than 24%, optionally less than 20%, optionally less than 15%, optionally less than 14%, optionally less than 13%, Optionally has a C14 Laves phase that is less than 12%.

  Provided is a hydrogen storage alloy having a larger intergranular region that promotes electrochemical properties and a sufficiently small crystallite size BCC phase to allow for a higher synergistic bond between the storage phase and the catalyst phase. The Alloys provided are 400 angstroms or less, optionally 390 angstroms or less, optionally 380 angstroms or less, optionally 370 angstroms or less, optionally 360 angstroms or less, optionally 350 angstroms or less, optionally 340 angstroms, optionally 330 angstroms or less , Optionally 320 angstroms or less, arbitrarily 310 angstroms or less, arbitrarily 300 angstroms or less, arbitrarily 290 angstroms or less, arbitrarily 280 angstroms or less, arbitrarily 270 angstroms or less, arbitrarily 260 angstroms or less, arbitrarily 250 angstroms or less, any 240 angstroms or less, optionally 230 angstroms or less, optionally 220 angstroms or less, optionally 210 Angstrom, arbitrarily 200 angstrom or less, arbitrarily 190 angstrom or less, arbitrarily 180 angstrom or less, arbitrarily 170 angstrom or less, arbitrarily 160 angstrom or less, arbitrarily 150 angstrom or less, arbitrarily 140 angstrom or less, optionally 130 angstrom Hereinafter, the crystallite size of the BCC phase is optionally 120 angstroms or less. Optionally, the crystallite size of the BCC phase is between 200 angstroms and 300 angstroms. Optionally, the crystallite size of the BCC phase is between 120 angstroms and 300 angstroms. Optionally, the crystallite size of the BCC phase is from 120 angstroms to 200 angstroms. Optionally, the crystallite size of the BCC phase is between 120 angstroms and 160 angstroms. Optionally, the crystallite size of the BCC phase is between 140 angstroms and 160 angstroms.

The physical, structural, and electrochemical properties of the hydrogen storage alloy are improved by preventing the formation of excessive Laves phases in the material, increasing the amount of BCC phase structure relative to the material, or a combination thereof. Thus, a method for activating (hydrogenating) a Laves phase related BCC metal hydride alloy of Formula I, Formula II, or Formula III is provided. This method involves exposing a Laves phase related BCC metal hydride alloy to an atmosphere containing hydrogen at a hydrogenation pressure and simultaneously cooling the alloy to have a desired capacity in cycle 10, optionally having a capacity of 200 mAh / g or more. Forming a metal hydride alloy. When a Laves phase related BCC metal hydride alloy is exposed to hydrogen at an elevated pressure, the temperature of the material increases due to the exothermic nature of the hydride formation reaction. It has been found that raising the temperature of the alloy without control is detrimental to the electrochemical properties of the resulting activated alloy. Thus, in some embodiments, the alloy is hydrogenated by a method that includes an active cooling step. Without being limited to a particular theory, controlling the temperature of the material during hydriding promotes the formation of excess AB _two- phase structure in the alloy during activation. Temperature control is achieved by cooling the reaction vessel, including a water jacketed system or bath, or by other methods known in the art. Optionally, the reaction temperature of the alloy does not exceed 300 ° C.

  In some embodiments, the temperature of the alloy is between room temperature and optionally 300 ° C., optionally 295 ° C., optionally 290 ° C., optionally 285 ° C., optionally 280 ° C., optionally 275 ° C., optionally during hydrogenation. 270 ° C, optionally 260 ° C, optionally 250 ° C, optionally 240 ° C, optionally 230 ° C, optionally 220 ° C, optionally 210 ° C, optionally 200 ° C, optionally 190 ° C, optionally 180 ° C, optionally 170 ° C, optionally 160 ° C, optionally 150 ° C, optionally 140 ° C, optionally 130 ° C, optionally 120 ° C, optionally 110 ° C, optionally 100 ° C, optionally 90 ° C, optionally 80 ° C, optionally 70 ° C, optionally 60 ° C, optionally 50 ° C, optionally 40 ° C, optionally 30 ° C. In some embodiments, the alloy is maintained during hydrogenation at a temperature between room temperature and 300 ° C., or any value or range therebetween.

  It has also been found that increasing the hydrogen pressure is useful in promoting the formation of increased BCC phase in the resulting activated hydrogen storage alloy. Thus, some aspects of hydrogenating Laves phase related BCC metal hydride alloys are 1.4 MPa or more, optionally 1.5 MPa or more, optionally 1.8 MPa or more, optionally 2 MPa or more, optionally 3 MPa or more. , Optionally at a hydrogenation pressure of 4 MPa or more, optionally 5 MPa or more, optionally 6 MPa or more.

  In some embodiments, the Laves phase related BCC metal hydride alloy is hydrogenated using both a hydrogenation pressure greater than 1.4 MPa and a temperature control up to 300 ° C. or less. Thus, the alloy is optionally activated at a hydrogenation pressure of 1.4 MPa to 6 MPa or more, and the alloy is 300 ° C., optionally 295 ° C., optionally 290 ° C., optionally 285 ° C., optionally 280 ° C., optionally 275 ° C., optionally 270 ° C., optionally 260 ° C., optionally 250 ° C., optionally 240 ° C., optionally 230 ° C., optionally 220 ° C., optionally 210 ° C., optionally 200 ° C., optionally 190 ° C., optionally 180 ° C, optionally 170 ° C, optionally 160 ° C, optionally 150 ° C, optionally 140 ° C, optionally 130 ° C, optionally 120 ° C, optionally 110 ° C, arbitrarily 100 ° C, optionally 90 ° C, arbitrarily At 80 ° C., optionally 70 ° C., optionally 60 ° C., optionally 50 ° C., optionally 40 ° C., optionally 30 ° C. In any one of the above temperature ranges, the hydrogenation pressure is arbitrarily 6 MPa, optionally 1.4 MPa, optionally 1.5 MPa, optionally 1.6 MPa, optionally 1.7 MPa, optionally 1.8 MPa. 1.9MPa, arbitrarily 2MPa, arbitrarily 2.1MPa, arbitrarily 2.2MPa, arbitrarily 2.3MPa, arbitrarily 2.4MPa, arbitrarily 2.5MPa, arbitrarily 2.6MPa, arbitrarily 2 0.7 MPa, arbitrarily 2.8 MPa, arbitrarily 2.9 MPa, arbitrarily 3 MPa, arbitrarily 3.1 MPa, arbitrarily 3.2 MPa, arbitrarily 3.3 MPa, arbitrarily 3.4 MPa, arbitrarily 3.5 MPa, arbitrarily 3.6 MPa, arbitrarily 3.7 MPa, arbitrarily 3.8 MPa, arbitrarily 3.9 MPa, arbitrarily 4 MPa, arbitrarily 4.1 MPa, arbitrarily 4.2 MPa, arbitrarily 4.3 MPa, arbitrarily 4.4 M a, Arbitrarily 4.5 MPa, Arbitrarily 4.6 MPa, Arbitrarily 4.7 MPa, Arbitrarily 4.8 MPa, Arbitrarily 4.9 MPa, Arbitrarily 5 MPa, Arbitrarily 5.1 MPa, Arbitrarily 5.2 MPa, Arbitrarily Between 5.3 MPa, optionally 5.4 MPa, optionally 5.5 MPa, optionally 5.6 MPa, optionally 5.7 MPa, optionally 5.8 MPa, optionally 5.9 MPa. In some embodiments, the hydrogenation pressure is 6 MPa or more.

  The resulting activated hydrogen storage alloy produced by the provided method has a capacity of nearly twice, often more than twice that of materials of similar composition produced by conventional methods.

  Various aspects of the invention are illustrated by the following non-limiting examples. The examples are for illustrative purposes and are not limiting in the practice of the invention. It will be understood that variations and modifications can be made without departing from the spirit and scope of the invention.

Tests In connection with a series of experiments illustrating the principles of the present invention, a series of metal hydride alloys of Formula I or Formula II were prepared and hydrogenated at various conditions. The raw material was arc melted using a non-consumable tungsten electrode and a water-cooled copper tray with a continuous flow of argon. Prior to formation, the residual oxygen concentration in the system was reduced by performing several melt-cooling cycles on the sacrificial titanium. Investigate ingots that have undergone several re-melting cycles to ensure uniform chemical composition.

The chemical composition of the prepared alloy sample was measured using a Varian Liberty 100 inductively coupled plasma emission spectrometer (ICP-OES) according to principles known in the art. The results of ICP by atomic percentage of ingot before activation are shown in Table 1:

  The as-cast composition is in excellent agreement with the designed composition. Previous series of alloys showed non-uniform Cr content, which was improved by increasing the power during arc melting. The measured B / A ratio of this series of alloys (P8-P14) varies from 2.61 to 5.36, which ranges from Young et al. Int. J. Hydrogen Energy, http: // dx Similar to the conventional series (P1-P8) from .doi.org / 10.1016 / j.ijhydene.2014.01.134 (article in press).

Phase distribution and composition The alloy phase distribution and composition were examined using a JEOL-JSM6320F scanning electron microscope capable of energy dispersive spectroscopy (EDS). The sample was mounted on an epoxy block, polished, rinsed and dried before being placed in the SEM chamber. Backscattered electron images are shown in FIGS. Several areas have been selected for study by EDS, which are indicated numerically in FIGS. 1A-G, respectively. The results of EDS measurement are shown in Table 2, and the compositions of the compositions P8 to P14 are shown in FIGS. 1A to 1G, respectively.

  As the Ti content in the alloy increases, the grain size of both the BCC phase and the C14 phase decreases and then increases. Unreacted metal Zr was found in the first few alloys (FIG. 1A spot 5). Further, since the increase in the produced Ti content increases the abundance of the C14 phase, almost no non-alloy Zr is found in the final sample. The TiNi phase begins to appear in alloys P13 and P14 as shown in FIGS. 1F-3 and 1G-3, and these alloys have the highest Ti content in the study.

  The Zr content, Ti content, V content, Ni content and B / A ratio of the C14 phase are plotted in FIG. 2A. As the total Ti content of the alloy increases, the Ti content of the C14 phase increases linearly and the Zr content decreases; the Ni content decreases and then stabilizes; the V content is almost unchanged And the B / A ratio decreases monotonically from 2.1 to 1.5.

Hydrideation P8-P14 alloys are subjected to various activation conditions by changing the maximum temperature of the alloy during activation by cooling the system, changing the hydrogen pressure, or both. The four activation processes are shown in Table 3.

  Control activated alloys using conventional activation methods and activated alloys prepared as in the methods of Examples 1-3 are subjected to analysis for gas phase hydrogen storage and electrochemical properties and structural configurations. .

XRD Analysis The microstructure of the alloy was studied using a Philips X'Pert Pro X-ray diffractometer. FIG. 3A shows XRD patterns (a to g, respectively) of all seven types of alloys P8 to P14.
It is clear that both C14 and BCC diffraction peaks are seen. An increase in V in the alloy corresponds to a decrease in the intensity of the BCC peak and a shift to a lower angle, with the C14 peak increasing in intensity and shifting in the same direction as the BCC peak.

  The XRD pattern (control and Examples 1-3) of a P8 sample hydrogenated under various conditions is shown in FIG. 3B. Two sets of diffraction peaks are seen, with C14 and BCC indicating the importance of these structures in the overall system.

  The crystallite size of each phase in all alloy samples is obtained from full pattern fitting of XRD data using the Rietveld method and Jade 9 software. The lattice constants of both phases calculated from the XRD pattern are shown in Table 4 and plotted in FIG.

As the amount of Ti increases, both the lattice constants a and c of the C14 phase and the lattice constant a of the BCC phase increase. In the C14 phase, the c / a ratio increases, stabilizes and decreases with increasing Ti content. The anisotropic growth of the C14 unit cell is due to partial substitution of B-site Cr (metal radius 1.423 angstrom of AB ₂ intermetallic alloy) and relatively large A-site Ti (radius 1.614 angstrom). As the V content in the alloy increases, the lattice constant a of the BCC phase increases from 2.9683 angstroms to 3.0165 angstroms.

  Table 4 also shows the crystallite size of each phase of the P8-P14 alloys. As the Ti content in the alloy increases, the abundance of the C14 phase increases from 11.0% by mass to 56.4% by mass and slightly decreases to 52.5% by mass for alloy P14. The BCC phase shows the reverse trend of decreasing from 89.0% to 47.5% from P8 to P14. At the same time, the C14 crystallite size first increases and then decreases, but the BCC crystallite size decreases monotonically.

FIG. 4 shows the FWHM of the BCC peak (110) for sample P8 as varied between the control and the sample hydrogenated by the methods of Examples 1-3. Table 5 shows the calculated crystallite size of the BCC phase.

  Exemplary hydrides each exhibit a BCC phase crystallite size of less than 300 Angstroms.

Table 6 shows the lattice constants a and c of each hydrogenated P8 sample.

Gas Phase Properties The gas phase hydrogen storage properties of the control and Examples 1-3 were measured using a Suzuki Shokan multi-channel pressure-concentration-temperature (PCT) system. Next, PCT isotherms at 30 ° C. and 60 ° C. were measured. The adsorption and desorption isotherms obtained for the control and hydrogenated alloy P8 as in Example 1 are shown in FIG. The P8 alloy hydrogenated as a control shows significant hysteresis, and the hydrogen mass percentage after activation does not exceed 1.1%. The same material activated by a method using the same hydrogen gas pressure, but controlled so that the maximum temperature of the alloy during activation does not exceed 300 ° C. is 1.5% higher than the completion of the second plateau. Shows a hydrogen mass percentage above.

The hysteresis of the PCT isotherm is defined as ln (P _a / P _d ), where P _a and P _d are the adsorption and desorption equilibrium pressures at the midpoint of the desorption isotherm, respectively. Hysteresis can be used to predict the grinding rate of the alloy during the cycle. The greater the hysteresis of the alloy, the faster the grinding rate during the hydrogenation / dehydrogenation cycle. From the hysteresis, a significant improvement in cycle stability is expected by activation according to the methods of Examples 1-3. In particular, the hysteresis is greatly reduced by cooling the ingot during activation so that the maximum temperature does not exceed 300 ° C.

P8-P14 alloys are similarly hydrogenated for PCT testing. All samples are activated in one thermal cycle in the presence of hydrogen at maximum pressure (5.0 MPa). As a result, the PCT characteristics are not greatly changed by the measurement. The resulting adsorption and desorption isotherms measured at 30 ° C. and 60 ° C. are shown in FIGS. 7A and 7B. As a comparison, as shown in FIG. 7A, the sample P8 ^* was activated and measured with the maximum hydrogen pressure set at 1.1 MPa, and all other samples were measured up to 5 MPa. Table 7 summarizes the information obtained from the PCT study.

Equilibrium pressures of desorption curves at 1.3 mass% and 1.5 mass% hydrogen occlusion were used as plateau pressures of alloys P8 to P10 and alloys P11 to P14, respectively. They tend to decrease with increasing Ti content. The lower equilibrium pressure arises from the expanded unit cell of both BCC and C14 phases. The hysteresis of the PCT isotherm is defined as ln (Pa / Pd), where Pa and Pd are the adsorption and desorption equilibrium pressures at the indicated hydrogen storage concentrations, respectively. In general, as the Ti content increases, the hysteresis decreases, which is expected to increase cycle stability by reducing grinding during cycling. As the Ti content increases, the maximum capacity increases due to the larger unit cell size and the ability to accommodate more hydrogen; however, the reversible hydrogen storage capacity (determined from the lower equilibrium pressure) ) Reduced due to increased metal hydrogen bond strength. Using desorption equilibrium pressures at 30 ° C., 60 ° C., and 90 ° C., the change in enthalpy (ΔH) and entropy (ΔS) is expressed by the following equation: ΔG = ΔH−TΔS = RTlnP (2)
(Where R is the ideal gas constant and T is the absolute temperature)
Estimated by These calculation results are shown in Table 7. These values are mainly useful for comparison of these alloys. As the Ti content in the alloy increases, both -ΔH and -ΔS first decrease and then increase. The ΔH value has no correlation with the plateau pressure, which is rarely seen in alloys with a single dominant phase. ΔS indicates how far the MH system is from a perfectly ordered state (degree of failure). The theoretical value of ΔS is the entropy of hydrogen gas, which is close to −135 Jmol− ¹ K− ¹ . It is interesting to see that alloys with more equal C14 and BCC abundances (P13 and P14) do not have the highest ΔS values. Instead, alloys P10 and P11 with high density BCC / C14 grain boundaries have the highest ΔS.

Electrochemical Properties Evaluation The discharge capacity of each alloy was measured in a flooded-cell configuration against a partially precharged Ni (OH) ₂ positive electrode. For in-cell electrochemical studies, each ingot was first ground and then passed through a 200 mesh screen. The sieved powder was then compressed with a 10 ton press onto an expanded nickel metal substrate to form a test electrode (area about 1 cm ² , thickness about 0.2 mm) without using a binder. This improved the measurement of activation behavior. The discharge capacity of the resulting small electrode was measured in a submerged cell configuration using a partially precharged Ni (OH) ₂ paste electrode as the positive electrode and a 6M KOH solution as the electrolyte. The electrode of each sample was charged for 10 hours at a constant current density of 50 mA / g, then discharged at a current density of 50 mA / g, and subsequently withdrawn twice at 12 mA / g and 4 mA / g. The total capacity (4 mA / g) from the first 13 cycles is plotted in FIG. 8A to investigate the activation and cycling behavior of these alloys. The activation of the alloy is facilitated due to the corrosion resistance of the BCC phase compared to the C14 phase and the corrosion resistance of Cr to KOH, but the abundance and Cr content of the BCC phase are reduced. A decrease in corrosion resistance helps activation, but also decreases cycle stability.

Capacity by discharge rate 50MAg ^-1 and 4MAg ^-1 as measured in the fourth cycle and the second cycle shown in Table 8.

As the overall Ti content increases, both capacities increase and then decrease. A maximum capacity of 463 mAh / g was measured for alloy P12. When the gas phase capacity is converted to the theoretical electrochemical capacity and plotted against the alloy number according to the measured electrochemical storage capacity, the 4 mA / g (low rate) curve follows the maximum gas phase capacity, 50 mA / g (high rate). ) The curve follows a reversible gas phase volume (FIG. 9). In this plot, the gas phase capacity is converted to electrochemical capacity by:
1% by mass of H ₂ = 268 mAhg ⁻¹ (3)

  The electrochemical capacity of this series of alloys is between the boundaries set by the maximum gas phase capacity and the reversible gas phase capacity, which are after the sample has been fully activated by the 5 MPa hydrogenation process. Similar to other MH systems.

The half-cell HRD value of each alloy, defined as the ratio of the discharge capacity measured at 50 mA / g to the discharge capacity measured at 4 mA / g, was measured in the fourth stabilized cycle and is shown in Table 8. With the exception of P10, the HRD of all alloys is stable within 4 cycles. As the Ti content in the alloy increases, the HRD decreases rapidly. Increasing the abundance of the C14 phase of the catalyst does not help HRD performance. Comparing the P8 ^* and P8 HRDs shows that the electrochemical HRD of P8 ^* is higher despite the lower hydrogen pressure used during activation and lower reversibility in the gas phase. (FIG. 7A). This observation can be attributed to various phases of the alloy. The two pressure plateaus seen in the PCT curve can be related to two different phases. The first phase having a lower plateau pressure (0.3 wt% to 1.0 wt% in FIG. 7A) is more than the second phase (1.0 wt% to 1.6 wt% in FIG. 7A). Has excellent speed capability. The C14 phase is believed to be more catalytic than the BCC phase. Therefore, the phase with the lower plateau pressure is assigned to the C14 phase and the phase with the higher plateau pressure is assigned to the BCC phase. The increase width of the first plateau is consistent with the increase in the abundance of the C14 phase, and the decrease width of the second plateau is consistent with the decrease in the abundance of the BCC phase obtained by XRD analysis (Table 4).

To investigate why the HRD decreases with increasing Ti content, both bulk diffusion coefficient (D) and surface exchange current (I _o ) were measured electrochemically. Details of both measurements have already been reported by F. Li, K. Young, T. Ouchi, MA Fetcenko, J. Alloys Compd. 471 (2009) 371-7, and these values are shown in Table 8. . The D value and C14 crystallite size are plotted against the alloy number in FIG. The D value tends to increase with increasing Ti content and the C14 crystallite size tends to decrease. This general trend indicates that the smaller the crystallite size of catalyst C14, the more interparticle regions can be formed, which promotes proton transfer and increases the D value.

  The patents, publications and applications mentioned in this specification are indicative of the level of those skilled in the art to which this invention pertains. These patents, publications, and applications are hereby incorporated by reference to the same extent as if each individual patent, publication, or application was specifically and individually incorporated herein by reference. Built in.

  In view of the above, it should be understood that other variations and modifications of the invention may be implemented. The above drawings, discussion, and description are illustrative of certain specific aspects of the invention, but are not intended to be limited in practice. The following claims include all equivalents that define the scope of the invention.

Claims (24)

Formula I below:
_{Ti w V x Cr y M z} (I)
(Where, w + x + y + z = 1, 0.1 ≦ w ≦ 0.6, 0.1 ≦ x ≦ 0.6, 0.01 ≦ y ≦ 0.6, and M is B, Al, Si, Sn, and Selected from the group consisting of transition metals)
Including the composition of
Laves phase related BCC metal hydride alloy having a capacity of greater than 200 milliamp hours per gram at cycle 10.   The alloy of claim 1 having a capacity of 350 milliamp hours or more per gram.   The alloy of claim 1 having a capacity of 400 milliamp hours or more per gram.   The alloy of claim 1 having a capacity of 420 milliamp hours or more per gram.   5. An alloy according to any one of claims 1 to 4, comprising less than 24% C14 phase.   The metal hydride alloy is primarily a combination of a BCC phase and a Laves phase, wherein the BCC phase is greater than 5% and less than 95%, and the Laves phase is greater than 5% and less than 95%. The alloy according to any one of claims 1 to 4, which is an abundance.   5. An alloy according to any one of claims 1 to 4, comprising a BCC phase crystallite size of less than 400 Angstroms.   5. An alloy according to any one of claims 1 to 4, comprising a BCC phase crystallite size of less than 200 Angstroms.   The alloy according to any one of claims 1 to 4, comprising a B / A ratio of 1.20 to 1.31.   The alloy according to any one of claims 1 to 4, wherein x / y is 1 to 3. Formula II below:
Ti _{13.6 + x} Zr _2.1 V ₄₄ Cr _13.2-x M _27.1 (II)
(Wherein x is greater than 0 and not greater than 12 and M is a combination of Mn, Fe, Co, Ni and Al)
The alloy according to any one of claims 1 to 4, comprising a composition of:   The alloy of claim 10, wherein x is 2, 4, 6, 8, 10 or 12.   The alloy according to claim 10, wherein x is 2 or 4. Formula III below:
Ti _{0.4 + x / 6} Zr _{0.6-x / 6} Mn _0.44 Ni _1.0 Al _0.02 Co _0.09 (VCr _0.3 Fe _0.063 ) _x (II)
(Wherein x is 0.7 to 2.8)
The alloy according to any one of claims 1 to 4, comprising a composition of: Formula I
_{Ti w V x Cr y M z} (I)
(Where, w + x + y + z = 1, 0.1 ≦ w ≦ 0.6, 0.1 ≦ x ≦ 0.6, 0.01 ≦ y ≦ 0.6, and M is B, Al, Si, Sn, and Selected from the group consisting of transition metals)
A method for activating a Laves phase-related BCC metal hydride alloy, comprising:
(I) subjecting the Laves phase related BCC metal hydride alloy to an atmosphere containing hydrogen at a hydrogenation pressure; and (ii) producing an activated metal hydride alloy having a capacity of greater than 200 milliamp hours per gram. Said method comprising cooling said alloy between steps.   The method of claim 15, wherein the cooling step has a maximum activation temperature of 300 degrees Celsius or less.   The method of claim 15, wherein the hydrogenation pressure is 1.4 megapascals or greater.   The method of claim 15, wherein the hydrogenation pressure is 6 megapascals or more.   19. A method according to any one of claims 15-18, wherein the activated metal hydride alloy has a capacity of 300 milliamp hours or more per gram.   19. A method according to any one of claims 15 to 18, wherein the activated metal hydride alloy has a capacity of 350 milliamp hours or more per gram.   19. A method according to any one of claims 15-18, wherein the activated metal hydride alloy has less than 24% C14 phase.   The activated metal hydride alloy is primarily a combination of a BCC phase and a Laves phase, the BCC phase is greater than 5% and less than 95%, and the Laves phase is greater than 5% and less than 95%. 19. The method according to any one of claims 15 to 18, wherein the abundance is.   19. A method according to any one of claims 15 to 18, wherein the activated metal hydride alloy comprises a BCC phase crystallite size of less than 400 Angstroms. The Laves phase-related BCC metal hydride alloy has the following formula III:
Ti _{0.4 + x / 6} Zr _{0.6-x / 6} Mn _0.44 Ni _1.0 Al _0.02 Co _0.09 (VCr _0.3 Fe _0.063 ) _x (II)
(Wherein x is 0.7 to 2.8)
19. A method according to any one of claims 15 to 18 wherein

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