U.S. patent application number 09/739919 was filed with the patent office on 2001-06-21 for modified electrochemical hydrogen storage alloy having increased capacity, rate capability and catalytic activity.
Invention is credited to Fetcenko, Michael A., Koch, John, Mays, William, Ovshinsky, Stanford R., Reichman, Benjamin, Young, Kwo.
Application Number | 20010003997 09/739919 |
Document ID | / |
Family ID | 23116904 |
Filed Date | 2001-06-21 |
United States Patent
Application |
20010003997 |
Kind Code |
A1 |
Fetcenko, Michael A. ; et
al. |
June 21, 2001 |
Modified electrochemical hydrogen storage alloy having increased
capacity, rate capability and catalytic activity
Abstract
A modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy which has at least one of the following
characteristics: 1) an increased charge/discharge rate capability
over that the base Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy; 2) a formation cycling requirement which is reduced
to one tenth that of the base Ti--V--Zr--Ni--Mn--Cr electrochemical
hydrogen storage alloy; or 3) an oxide surface layer having a
higher electrochemical hydrogen storage catalytic activity than the
base Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen storage
alloy.
Inventors: |
Fetcenko, Michael A.;
(Rochester Hills, MI) ; Young, Kwo; (Troy, MI)
; Ovshinsky, Stanford R.; (Bloomfield Hills, MI) ;
Reichman, Benjamin; (West Bloomfield, MI) ; Koch,
John; (Brighton, MI) ; Mays, William;
(Livonia, MI) |
Correspondence
Address: |
Dean B. Watson
Energy Conversion Devices, Inc.
1675 W. Maple Rd.
Troy
MI
48084
US
|
Family ID: |
23116904 |
Appl. No.: |
09/739919 |
Filed: |
December 18, 2000 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
|
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09739919 |
Dec 18, 2000 |
|
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09290633 |
Apr 12, 1999 |
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Current U.S.
Class: |
148/442 ;
420/588; 420/900; 428/559 |
Current CPC
Class: |
C22C 16/00 20130101;
C22C 27/06 20130101; C22C 14/00 20130101; C22C 27/025 20130101;
Y10S 420/90 20130101; H01M 4/383 20130101; C01B 3/0031 20130101;
Y02E 60/10 20130101; Y02E 60/32 20130101; Y10T 428/12104
20150115 |
Class at
Publication: |
148/442 ;
420/900; 420/588; 428/559 |
International
Class: |
B22F 007/04; B32B
015/02 |
Claims
We claim:
1. A modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy, said alloy comprising a base alloy and at least one
modifying element, said base alloy consisting essentially of 0.1 to
60% Ti, 0.1 to 40% Zr, 0 to 60% V, 0.1 to 57% Ni, 5 to 22% Mn and 0
to 56% Cr, said modified alloy characterized by an increased
charge/discharge rate capability over that the base
Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen storage alloy.
2. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 1, wherein said increased charge/discharge
rate capability over that of the base Ti--V--Zr--Ni--Mn--Cr
electrochemical hydrogen storage alloy comprises at least a 300%
greater charge/discharge rate capability at a discharge rate of
1000 mA/g.
3. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 1, wherein said at least one modifying
element comprises Al, Co, Sn and optionally Fe.
4. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 3, further comprising 0.1 to 10.0% Al, 0.1
to 10.0% Co, 0 to 3.5% Fe and 0.1 to 3.0% Sn.
5. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 4, comprising, in atomic percentage, Ti
9.0%, Zr 26.2%, V 5.0%, Ni 38.0%, Cr 3.5%, Co 1.5%, Mn 15.6%, Al
0.4%, and Sn 0.8%.
6. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 1, wherein said alloy has a higher exchange
current density than the base Ti--V--Zr--Ni--Mn--Cr electrochemical
hydrogen storage alloy.
7. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 6, wherein said exchange current density is
at least 40 mA/cm.sup.2 higher than that of the base
Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen storage alloy.
8. A modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy, said alloy comprising a base alloy and at least one
modifying element, said base alloy consisting essentially of 0.1 to
60% Ti, 0.1 to 40% Zr, 0 to 60% V, 0.1 to 57% Ni, 5 to 22% Mn and 0
to 56% Cr, said modified alloy characterized by a formation cycling
requirement which is reduced to at least one tenth that of the base
Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen storage alloy.
9. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 8, wherein said at least one modifying
element comprises Al, Co, Sn and optionally Fe.
10. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 9, further comprising 0.1 to 10.0% Al, 0.1
to 10.0% Co, 0 to 3.5% Fe and 0.1 to 3.0% Sn.
11. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 10, comprising, in atomic percentage, Ti
9.0%, Zr 26.2%, V 5.0%, Ni 38.0%, Cr 3.5%, Co 1.5%, Mn 15.6%, Al
0.4%, and Sn 0.8%.
12. A modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy, said alloy comprising a base alloy and at least one
modifying element, said base alloy consisting essentially of 0.1 to
60% Ti, 0.1 to 40% Zr, 0 to 60% V, 0.1 to 57% Ni, 5 to 22% Mn and 0
to 56% Cr, said modified alloy characterized by an oxide surface
layer having a higher electrochemical hydrogen storage catalytic
activity than the base Ti--V--Zr--Ni--Mn--Cr electrochemical
hydrogen storage alloy.
13. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 12, wherein said at least one modifying
element comprises Al, Co, Sn and optionally Fe.
14. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 13, further comprising 0.1 to 10.0% Al, 0.1
to 10.0% Co, 0 to 3.5% Fe and 0.1 to 3.0% Sn.
15. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 14, comprising, in atomic percentage, Ti
9.0%, Zr 26.2%, V 5.0%, Ni 38.0%, Cr 3.5%, Co 1.5%, Mn 15.6%, Al
0.4%, and Sn 0.8%.
16. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 12, wherein said oxide comprises an oxide of
Mn.
17. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 16, wherein said oxide comprises an oxide of
Mn and Ni.
18. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 17, wherein said oxide comprises an oxide of
Mn, Ni, Co, Ti and optionally Fe, if Fe is present in the
alloy.
19. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 16, wherein said oxide comprises an oxide of
Mn, Co and Ti.
20. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 12, wherein said oxide comprises mixed fine
and coarse grained oxides.
21. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 12, wherein said oxide layer contains
metallic catalytic particles of from 10-50 Angstroms in size
dispersed therein.
22. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 21, wherein said metallic catalytic
particles are from 10-40 Angstroms in size.
23. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 22, wherein said metallic catalytic
particles are from 10-30 Angstroms in size.
24. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 23, wherein said metallic catalytic
particles are from 10-20 Angstroms in size.
25. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 24, wherein said metallic catalytic
particles are Ni--Mn--Co--Ti alloy catalytic particles.
26. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 21, wherein said metallic catalytic
particles are finely dispersed in said oxide from 10-20 Angstroms
apart.
27. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 21, wherein said metallic catalytic
particles have an FCC crystalline structure.
28. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 21, wherein said metallic catalytic
particles have a NiCoMnTi and TiZr oxide support structure forming
a super lattice to promote ionic diffusion and reaction.
29. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 12, wherein said alloy has a higher surface
area than the base Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy.
30. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 29, wherein said alloy has a surface area of
at least 0.2 m.sup.2/g higher than the base Ti--V--Zr--Ni--Mn--Cr
electrochemical hydrogen storage alloy.
31. An electrochemical hydrogen storage alloy having an oxide
surface containing metallic catalytic particles, said particles
being from 10-50 Angstroms in size.
32. The electrochemical hydrogen storage alloy of claim 31, wherein
said metallic catalytic particles are from 10-40 Angstroms in
size.
33. The electrochemical hydrogen storage alloy of claim 33, wherein
said metallic catalytic particles are from 10-30 Angstroms in
size.
34. The electrochemical hydrogen storage alloy of claim 33, wherein
said metallic catalytic particles are from 10-20 Angstroms in
size.
35. The electrochemical hydrogen storage alloy of claim 31, wherein
said metallic catalytic particles are Ni--Mn--Co--Ti alloy
catalytic particles.
36. The electrochemical hydrogen storage alloy of claim 31, wherein
said metallic catalytic particles are finely dispersed in said
oxide from 10-20 Angstroms apart.
37. The electrochemical hydrogen storage alloy of claim 31, wherein
said metallic catalytic particles have an FCC crystalline
structure.
38. The electrochemical hydrogen storage alloy of claim 31, wherein
said oxide comprises an oxide of Mn.
39. The electrochemical hydrogen storage alloy of claim 38, wherein
said oxide comprises an oxide of Mn and Ni.
40. The electrochemical hydrogen storage alloy of claim 39, wherein
said oxide comprises an oxide of Mn, Ni, Co, Ti and optionally Fe,
if Fe is present in the alloy.
41. The electrochemical hydrogen storage alloy of claim 38, wherein
said oxide comprises an oxide of Mn, Co and Ti.
42. The electrochemical hydrogen storage alloy of claim 31, wherein
said oxide comprises mixed fine and coarse grained oxides.
43. The electrochemical hydrogen storage alloy of claim 31, wherein
said alloy comprises a base alloy and at least one modifying
element, said base alloy consisting essentially of 0.1 to 60% Ti,
0.1 to 40% Zr, 0 to 60% V, 0.1 to 57% Ni, 5 to 22% Mn and 0 to 56%
Cr, and said at least one modifying element comprises Al, Co, Sn
and optionally Fe.
44. The electrochemical hydrogen storage alloy of claim 43, further
comprising 0.1 to 10.0% Al, 0.1 to 10.0% Co, 0 to 3.5% Fe and 0.1
to 3.0% Sn.
45. The electrochemical hydrogen storage alloy of claim 44,
comprising, in atomic percentage, Ti 9.0%, Zr 26.2%, V 5.0%, Ni
38.0%, Cr 3.5%, Co 1.5%, Mn 15.6%, Al 0.4%, and Sn 0.8%.
46. The electrochemical hydrogen storage alloy of claim 31, wherein
said oxide surface comprises TiZr oxide, FCC nickel particles, and
regions of NiCoMnTi, said TiZr oxide, FCC nickel particles, and
regions of NiCoMnTi interacting to form a super lattice which
promotes ionic diffusion and reaction.
47. A modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy comprising, in atomic percentage:(Base
Alloy).sub.aCo.sub.bFe.sub.c- Al.sub.dSn.sub.e, where said Base
Alloy comprises 0.1 to 60% Ti, 0.1 to 40% Zr, 0 to 60% V, 0.1 to
57% Ni, 5 to 22% Mn and 0 to 56% Cr; b is 0.1 to 10.0%; c is 0 to
3.5%; d is 0.1 to 10.0%; e is 0.1 to 3.0%; and a+b+c+d+e=100%.
48. The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy of claim 47 comprising, in atomic percentage, Ti
9.0%, Zr 26.2%, V 5.0%, Ni 38.0%, Cr 3.5% Co 1.5%, Mn 15.6% Al 0.4%
and Sn 0.8%.
49. In a method of casting an ingot of a modified
Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen storage alloy
comprising, in atomic percentage (Base
Alloy).sub.aCo.sub.bFe.sub.cAl.sub.dSn.sub.e, where said Base Alloy
comprises 0.1 to 60% Ti, 0.1 to 40% Zr, 0 to 60% V, 0.1 to 57% Ni,
5 to 22% Mn and 0 to 56% Cr; b is 0.1 to 10.0%; c is0 to 3.5%; d is
0.1 to 10.0%; e is 0.1 to 3.0%; and a+b+c+d+e=100% the improvement
comprising: limiting the ingot thickness to less than about 5
inches.
50. The method of casting an ingot of a modified
Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen storage alloy of
claim 49, further comprising limiting the ingot thickness to less
than about one inch.
Description
FIELD OF THE INVENTION
[0001] The instant invention involves electrochemical hydrogen
storage alloys and more specifically to modified VTiZrNiCrMn
alloys. Most specifically, the instant invention comprises a
modified VTiZrNiCrMn-based alloy which has at least one of 1) an
increased charge/discharge rate capability over that of the base
VTiZrNiCrMn electrochemical hydrogen storage alloy, 2) a formation
cycling requirement which is reduced to one tenth that of the base
VTiZrNiCrMn electrochemical hydrogen storage alloy, or 3) an oxide
surface layer having a higher electrochemical hydrogen storage
catalytic activity than the base Ti--V--Zr--Ni--Mn--Cr
electrochemical hydrogen storage alloy.
BACKGROUND OF THE INVENTION
[0002] In rechargeable alkaline cells, weight and portability are
important considerations. It is also advantageous for rechargeable
alkaline cells to have long operating lives without the necessity
of periodic maintenance. Rechargeable alkaline cells are used in
numerous consumer devices such as portable computers, video
cameras, and cellular phones. They are often configured into a
sealed power pack that is designed as an integral part of a
specific device. Rechargeable alkaline cells can also be configured
as larger cells that can be used, for example, in industrial,
aerospace, and electric vehicle applications.
[0003] The materials proposed in the prior art for use as hydrogen
storage negative electrode materials for secondary batteries are
materials that have essentially simple crystalline structures. In
simple crystalline materials, limited numbers of catalytic site are
available resulting from accidently occurring, surface
irregularities which interrupt the periodicity of the crystalline
lattice. A few examples of such surface irregularities are
dislocation sites, crystal steps, surface impurities and foreign
absorbates. For more than three decades, virtually every battery
manufacturer in the world pursued such crystalline electrode
materials for electrochemical applications, but none produced a
commercially viable nickel metal hydride battery until after the
publication of U.S. Pat. No. 4,623,597 (the '597 patent) to
Ovshinsky, et al, which disclosed fundamentally new principles of
electrode material design.
[0004] As taught in the '597 patent (the contents of which are
incorporated by reference), a major shortcoming of basing negative
electrode materials on simple ordered crystalline structures is
that irregularities which result in the aforementioned
catalytically active sites occur relatively infrequently. This
results in a relatively low density of catalytic and/or storage
sites and consequently poor stability. Of equal importance is that
the type of catalytically active sites available are of an
accidental nature, relatively few in number and are not designed
into the material as are those of the present invention. Thus, the
efficiency of the material in storing hydrogen and its subsequent
release is substantially less than that which would be possible if
a greater number and variety of sites were available.
[0005] Ovshinsky, et al, fundamental principles overcome the
limitations of the prior art by improving the characteristics of
the negative electrode through the use of disordered materials to
greatly increase the reversible hydrogen storage characteristics
required for efficient and economical battery applications. By
applying the principles of disorder, it has become possible to
obtain a high energy storage, efficiently reversible, and high
electrically efficient battery in which the negative electrode
material resists structural change and poisoning, with improved
resistance to the alkaline environment, good self-discharge
characteristics and long cycle life and deep discharge
capabilities. The resulting disordered negative electrode materials
are formed from lightweight, low cost elements by techniques that
assure formation of primarily non-equilibrium metastable phases
resulting in high energy and power densities at low cost. These
non-equilibrium, metastable phases assure the formation of
localized states where a special degree of disorder, if properly
fabricated, can come from the structural and compositional disorder
of the material.
[0006] The materials described generally in the '597 patent have a
greatly increased density of catalytically active sites providing
for the fast and stable storage and release of hydrogen. This
significantly improved the electrochemical charging/discharging
efficiencies and also showed an increase in hydrogen storage
capacity. Generally, this was accomplished by the bulk storage of
hydrogen atoms at bonding strengths within the range of reversible
electromotive force suitable for use in secondary battery
applications. More specifically, such negative electrode materials
were fabricated by manipulating the local chemical order and hence
the local structural order by the incorporation of selected
modifier elements into the host matrix to create the desired
disorder, type of local order and metal hydrogen bond strengths.
The resulting multicomponent disordered material had a structure
that was amorphous, microcrystalline, multiphase polycrystalline
(but lacking long range compositional order), or a mixture of any
combination of these structures.
[0007] The host matrix of the materials described in the '597
patent were formed from elements capable of storing hydrogen an
thus are considered hydride formers. This host matrix was modified
by incorporating selected modifier elements which could also be
hydride formers. These modifiers enhanced the disorder of the final
material, thus creating a much greater number and spectrum of
catalytically active sites with an increase in the number of
hydrogen storage sites. Multiorbital modifiers (such as transition
elements) provided the greatly increased number of sites due to
various bonding configurations available. Because of more efficient
storage and release of hydrogen and because of the higher density
of the catalytic sites, the hydrogen more readily found a storage
site. Unfortunately, there remained, until U.S. Pat. No. 5,840,440
('440), an insufficient density of new hydrogen storage sites
formed due to disorder to significantly increase the hydrogen
storage capacity of the material.
[0008] The '597 patent describes the use of, inter alia, rapid
quench to form disordered materials having unusual electronic
configurations, which results from varying the three-dimensional
interactions of constituent atoms and their various orbitals. Thus,
it was taught that the compositional, positional and translational
relationships of the constituent atoms were not limited by
crystalline symmetry in their freedom to interact. Selected
elements could be utilized to further control the disorder of the
material by their interaction with orbitals so as to create the
desired local internal chemical environments. These various and at
least partially unusual configurations generate a large number of
catalytically active sites and hydrogen storage sites not only on
the surface but throughout the bulk of the material. The internal
topology generated by these various configurations allowed for
selective diffusion of hydrogen atoms.
[0009] In general, disorder in the modified material can be of an
atomic nature in the form of compositional or configurational
disorder provided throughout the bulk of the material or in
numerous regions or phases of the material. Disorder can also be
introduced into the host matrix by creating microscopic phases
within the material which mimic the compositional or
configurational disorder at the atomic level by virtue of the
relationship of one phase to another. For example, disordered
materials can be created by introducing microscopic regions or
phases of a different kind or kinds of crystalline phases, or by
introducing regions of an amorphous phase or phases, or by
introducing regions of an amorphous phase or phases in addition to
regions of a crystalline phase or phases. The types of disordered
structures that provide local structural chemical environments for
improved hydrogen storage characteristics include amorphous
materials, microcrystalline materials, multicomponent multiphase
polycrystalline materials lacking long range composition order or
multiphase materials containing both amorphous and crystalline
phases.
[0010] Short-range, or local, order is elaborated on in U.S. Pat.
No. 4,520,039 to Ovshinsky, entitled Compositionally Varied
Materials and Method for Synthesizing the Materials, the contents
of which are incorporated by reference. This patent discloses that
disordered materials do not require periodic local order and how
spatial and orientational placement of similar or dissimilar atoms
or groups of atoms is possible with such increased precision and
control of the local configurations that it is possible to produce
qualitatively new phenomena. In addition, this patent discusses
that the atoms used need not be restricted to "d band" or "f band"
atoms, but can be any atom in which the controlled aspects of the
interaction with the local environment and/or orbital overlap plays
a significant role physically, electronically, or chemically so as
to affect physical properties and hence the functions of the
materials. The elements of these materials offer a variety of
bonding possibilities due to the multidirectionality of f-orbitals,
d-orbitals or lone pair electrons. The multidirectionality
("porcupine effect") of d-orbitals provides for a tremendous
increase in density of sites, the spectrum of types of sites and
hence the presence of active storage sites. Following the teaching
can result in a means of synthesizing new materials which are
disordered in several different senses simultaneously.
[0011] The '597 patent is described in detail above because this
patent represents a starting point for the investigation that
resulted in the present invention. That patent introduced the
concept of making negative electrode material for nickel metal
hydride batteries from multicomponent disordered alloys. This
teaching was diametrically opposed to the conventional "wisdom" of
battery manufacturers at the time. It was not until this concept
was adopted in production processes by said manufacturers that
negative electrode materials with an increased number of
catalytically active sites were produced and nickel metal hydride
batteries became commercially viable. In capsule form, the '597
patent taught that significant additional sites for hydrogen
catalysis (to allow the rapid storage and release of hydrogen and
greatly improve stability) were formed and made available by
purposely fabricating disordered negative electrode material (as
opposed to the homogeneous, ordered polycrystalline material of the
prior art). The '597 patent also proposed that the use of disorder
could be employed to obtain additional hydrogen storage sites
However, it was not appreciated that in order to obtain a
substantial increase in hydrogen storage capacity from such
non-conventional storage sites, it would be necessary to increase
the number of storage sites by approximately 3 orders of
magnitude.
[0012] Not only was the teaching of the '597 patent adopted by all
nickel metal hydride manufacturers, but in recent years some of
these manufacturers have begun to use rapid solidification
techniques (as taught by Ovshinsky) to increase the degree of
disorder within a negative electrode alloy formula. For instance,
battery companies have even gone so far as to rapidly quench
highly-modified LaNi.sub.5-type electrochemical negative electrode
material. By employing nonequilibrium processing techniques, the
resulting negative electrode material includes hydrogen storage
phases and catalytic phases on the order of 2000 Angstroms in
average dimension. The hydrogen storage capacity of the negative
electrode material does not improve significantly, but the
catalytic activity is greatly improved as evidenced by improved
rate capability and stability to oxidation and corrosion, resulting
in increased cycle life.
[0013] As mentioned above, certain battery companies have begun to
investigate the use of rapidly-quenched, highly modified LaNi.sub.5
type hydrogen storage materials for electrochemical applications.
For example, in Phys. Chem 96 (1992) No. 5 pp. 656-667, P. H. L.
Notten, et al of Philips Research Laboratories presented a paper
entitled "Melt-Spinning of AB.sub.55-Type Hydride Forming Compounds
and the Influence of Annealing on Electrochemical and
Crystallographic Properties." In this paper, non-stoichiometric
modified LaNi.sub.55 materials,
La.sub.6Nd.sub.2Ni.sub.3Co.sub.24Si.sub.1 and
La.sub.6Nd.sub.2Ni.sub.26Co- .sub.24Mo.sub.1 were rapidly
solidified. These non-stoichiometric materials were compared to
their stoichiometric counterparts with the result being that the
non-stoichiometric materials demonstrated good, but not unusual
hydrogen storage capacity. However, the non-stoichiometric
compounds did show the presence of additional volume percents of a
catalytic phase, which phase, as predicted by the '597 patent, was
responsible for the improved electrochemical properties as compared
to the properties found in the examples of stoichiometric material.
Once again, and more importantly, no significantly higher density
of non-conventional hydrogen storage sites were obtained. In
research and development activities at Energy Conversion Devices,
Inc. with highly modified TiNi-type electrochemical negative
electrode materials, such as described in U.S. Pat. No. '440 which
is incorporated herein by references, rapidly quenched electrode
materials were melt spun. The work resulted in improved oxidation
and corrosion resistance and cycle life was increased by a factor
of five. On the other hand and as was true in the case of the
aforementioned Japanese work, no significant increase in hydrogen
storage capacity was demonstrated and the phases of the negative
electrode material present were also relatively large.
[0014] Therefore, while the teachings of the '597 patent were
revolutionary for those of ordinary skill in the art in learning to
apply the principals of disorder disclosed therein to negative
electrode materials to obtain commercial batteries with
commercially viable discharge rates and cycle life stabilities
while maintaining good hydrogen storage capacity, the '597 patent
provided for the most part generalities to routineers concerning
specific processes, processing techniques, alloy compositions,
stoichiometries in those compositions, etc. regarding how to
further significantly increase the hydrogen storage capacity (as
opposed to the catalytic activity). It was not until the '440
patent that a teaching was presented of the nature and
quantification of additional active storage sites required to
significantly increase the hydrogen storage capacity of the
negative electrode material through the deliberate introduction of
defect sites and the presence of other concurrent non-conventional
and/or conventional storage sites.
[0015] Despite the exceptional electrochemical performance now
provided by current highly disordered nickel metal hydride systems
(twice the hydrogen storage capacity of conventional NiCd systems)
consumers are demanding increasingly greater run times, safety and
power requirements from such rechargeable battery systems. No
current battery system can meet these demands. Accordingly, there
exists a need for a safe ultra high capacity, high charge
retention, high power delivery, long cycle life, reasonably priced
rechargeable battery system.
[0016] While U.S. Pat. No. 5,840,440 ("the '440patent") represents
innovative ideas with respect to useable storage sites in an
electrochemical negative electrode material due to the use of high
defect density and small crystallite size, the focus of the '440
patent is on the bulk properties of the hydrogen storage alloy.
Significant discussion therein relates to increased surface sites;
however, the additional sites so described relate to the interior
surfaces, or grain boundaries, again within the alloy. The '440
patent does not address the interface between the metal hydride
alloy and the electrolyte at the so-called oxide layer.
[0017] Of most relevance to the present invention is commonly
assigned U.S. Pat. No. 5,536,591 ("the '591 patent") in which the
oxide (metal/electrolyte) interface is addressed in detail and
where teachings on composition, size and distribution of catalytic
sites within the oxide interface was first provided.
[0018] The '591 patent taught that hydrogen storage and other
electrochemical characteristics of the electrode materials thereof
could be controllably altered depending on the type and quantity of
host matrix material and modifier elements selected for making the
negative electrode materials. The negative electrode alloys of the
'591 patent were resistant to degradation by poisoning due to the
increased number of selectively designed storage and catalytically
active sites which also contributed to long cycle life. Also, some
of the sites designed into the material could bond with and resist
poisoning without affecting the active hydrogen sites. The
materials thus formed had a very low self discharge and hence good
shelf life.
[0019] As discussed in U.S. Pat. No. 4,716,088 ("the '088 patent"),
the contents of which are specifically incorporated by reference,
it is known that the steady state surface composition of
V--Ti--Zr--Ni alloys can be characterized as having porous,
catalytic regions of enriched nickel. An aspect of the '591 patent
was a significant increase in the frequency of occurrence of these
nickel regions as well as a more pronounced localization of these
regions. More specifically, the materials of the '591 patent had
discrete nickel regions of 50-70 Angstroms in diameter distributed
throughout the oxide interface and varying in proximity from 2-300
Angstroms or preferably 50-100 Angstroms, from region to region.
This was illustrated in the FIG. 1 or the '591 patent, where the
nickel regions 1 were shown as what appear as particles on the
surface of the oxide interface 2 at 178,000 X. As a result of the
increase in the frequency of occurrence of these nickel regions,
the materials of the '591 patent exhibited significantly increased
catalysis and conductivity.
[0020] The increased density of Ni regions in the materials of the
'591 patent provided metal hydride powder particles having a highly
catalytic surface. Prior to the '591 patent, Ni enrichment was
attempted unsuccessfully using microencapsulation. The method of Ni
encapsulation results in the expensive physical, chemical or
electrochemical deposition of a layer of Ni at the
metal-electrolyte interface. The deposition of an entire layer was
expensive, excessive and resulted in no improvement of performance
characteristics since this kind of encapsulated layer did not
result in the production of the localized, finely distributed
nickel regions of 50-70 Angstrom in a porous matrix.
[0021] The enriched Ni regions of the '591 patent could be produced
via two general fabrication strategies. The first of these
strategies was to specifically formulate an alloy having a surface
region that is preferentially corroded during activation to produce
the described enriched Ni regions. It was believed that Ni was in
association with an element such as Al at specific surface regions
and that this element corroded preferentially during activation,
leaving the enriched Ni regions described in the '591 patent.
"Activation" as used herein specifically refers to "etching" or
other methods of removing excessive oxides, such as described in
the '088 patent as applied to electrode alloy powder, the finished
electrode, or at any point in between in order to improve the
hydrogen transfer rate.
[0022] The second of these strategies was to mechanically alloy a
secondary alloy to a hydride battery alloy, where the secondary
alloy preferentially corroded to leave enriched nickel regions. An
example of such a secondary alloy was given as NiAl. The most
preferred alloys having enriched Ni regions were alloys having the
following composition: (Base
Alloy).sub.aCo.sub.bMn.sub.cFe.sub.dSn.sub.e where the Base Alloy
comprised 0.1 to 60 atomic percent Ti, 0.1 to 40 atomic percent Zr,
0 to 60 atomic percent V, 0.1 to 57 atomic percent Ni, and 0 to 56
atomic percent Cr; b was 0 to 7.5 atomic percent; c was 13 to 17
atomic percent; d was 0 to 3.5 atomic percent; e was 0 to 1.5
atomic percent; and a+b+c+d+e=100 atomic percent.
[0023] The production of the Ni regions of the '591 patent was
consistent with a relatively high rate of removal through
precipitation of the oxides of titanium and zirconium from the
surface and a much lower rate of nickel removal, providing a degree
of porosity to the surface. The resultant surface had a higher
concentration of nickel than would be expected from the bulk
composition of the negative hydrogen storage electrode. Nickel in
the metallic state is electrically conductive and catalytic,
imparting these properties to the surface. As a result, the surface
of the negative hydrogen storage electrode was more catalytic and
conductive than if the surface contained a higher concentration of
insulating oxides. Many of the alloys of the '591 patent include
Mn. The effects of the addition of Mn to these alloys was generally
discussed in U.S. Pat. No. 5,096,667, the disclosure of which is
incorporated herein by reference. The addition of Mn usually
results in improved charging efficiency. This effect appears to
result from the ability of Mn to improve the charging efficiency of
alloys into which it is added by improving oxidation resistance and
oxygen recombination. It has been observed that oxygen gas
generated at the nickel hydroxide positive electrode recombined at
the surface of the metal hydride electrode. Oxygen recombination is
an especially aggressive oxidizer of its environment, even compared
to the alkaline electrolyte.
[0024] It is possible that the modifier elements added to the Base
Alloy of the '591 patent, particularly Mn and Fe, and most
particularly Co, either alone, or in combination with Mn and/or Al
for example, act to catalyze oxygen reduction, thereby avoiding or
reducing the oxidation of the surrounding elements in the metal
hydride alloy. It is believed that this function of the modified
alloys reduces or even eliminates the formation and build up of
detrimental surface oxide, thereby providing a thinner and more
stable surface.
[0025] It is believed that several additional factors may explain
the unexpected behavior of Mn and Fe in the Base Alloys of the
present invention: (1) The combination of Mn and Fe may affect the
bulk alloy by inhibiting the bulk diffusion rate of hydrogen within
the metal through the formation of complex phase structures, either
by effecting the grain boundaries or by affecting the equilibrium
bond strength of hydrogen within the metal. In other words, the
temperature dependence of the hydrogen bond strength may be
increased thereby decreasing the available voltage and capacity
available under low temperature discharge. (2) It is believed that
the combination of Mn and Fe may result in a lower electrode
surface area for metallurgical reasons by increasing the ductility
of the alloy and thereby reducing the amount of surface area
formation during the activation process. (3) It is believed that
the combination of Mn and Fe to these alloys may inhibit low
temperature discharge through the alteration of the oxide layer
itself with respect to conductivity, porosity, thickness, and/or
catalytic activity. The oxide layer is an important factor in the
discharge reaction and promotes the reaction of hydrogen from the
Base Alloy of the present invention and hydroxyl ion from the
electrolyte. It is believed that this reaction is promoted by a
thin, conductive, porous oxide having some catalytic activity.
[0026] The combination of Mn and Fe does not appear to be a problem
under room temperature discharge, but has shown a surprising
tendency to retard the low temperature reaction. The formation of a
complex oxide could result in a subtle change in oxide structure
such as pore size distribution or porosity. Since the discharge
reaction produces water at the metal hydride surface and within the
oxide itself, a small pore size may be causing a slow diffusion of
K.sup.+and OH.sup.-ions from the bulk of the electrolyte to the
oxide. Under room temperature discharge where polarization is
almost entirely ohmic to low temperature discharge where activation
and concentration polarization components dominate the physical
structure of the oxides with Fe and Mn compared to Mn alone could
be substantially different.
[0027] Still another possible explanation is that Mn and Fe have
multivalent oxidation states. Some elements within the oxide may in
fact change oxidation state during normal state of charge variance
as a function of the rate of discharge and can be both temperature,
fabrication, and compositionally dependant. It is possible these
multiple oxidation states have different catalytic activity as well
as different densities that together effect oxide porosity. A
possible problem with a complex oxide containing both Mn and Fe
could be that the Fe component retards the ability of the Mn to
change oxidation state if present in large quantities.
[0028] Throughout the preceding discussion with respect to the
oxide it should be noted that the oxide also contains other
components of the Base Alloy, such as V, Ti, Zr, Ni, and/or Cr and
other modifier elements. The discussion of a complex oxide of Mn
and Fe is merely for the sake of brevity and one skilled in the art
should not infer that the actual mechanism cannot also include a
different or more complex explanation involving other such
elements.
[0029] Deficiencies of the Prior Art
[0030] While prior art hydrogen storage alloys frequently
incorporate various individual modifiers and combinations of
modifiers to enhance performance characteristics, there is no clear
teaching of the role of any individual modifier, the interaction of
any modifier with other components of the alloy, or the effects of
any modifier on specific operational parameters. Because highly
modified LaNi.sub.5 alloys were being analyzed from within the
context of well ordered crystalline materials, the effect of these
modifiers, in particular, was not clearly understood.
[0031] Prior art hydrogen storage alloys, when incorporated into
batteries, have generally exhibited improved performance
attributes, such as cycle life, rate of discharge, discharge
voltage, polarization, self discharge, low temperature capacity,
and low temperature voltage. However, prior art alloys have yielded
batteries that exhibit a quantitative improvement in one or two
performance characteristic at the expense of a quantitative
reduction in other performance characteristics
[0032] Electrical formation is defined as charge/discharge cycling
required to bring the batteries up to their ultimate performance.
For prior art alloys, electrical formation is essential for maximum
battery performance at both high and low discharge rates. For
instance certain prior art VTiZrNiCrMn alloys could require as many
as 32 cycles of charge and discharge at various rates to fully form
the electric vehicle battery. It is believed that this electrical
formation causes expansion and contraction of the negative
electrode alloy material as it alternately stores and releases
hydrogen. This expansion and contraction induces stress and forms
in-situ cracks within the alloy material. The cracking increases
the surface area, lattice defects and porosity of the alloy
material. Heretofore, NiMH batteries have required this electrical
formation treatment.
[0033] There is no "set-in-stone" method of electrical formation.
The reason for this is that different active metal hydride
materials which have been prepared by different methods under
different conditions, and formed into electrodes by different
methods will require different electrical formation processing.
Hence, no detailed method of electrical formation suitable for all
batteries can be described. However, generally electrical formation
involves a relatively complex procedure of cycling the prepared
battery through a number of charge/discharge cycles at varying
rates of charge/discharge to varying depths of
charge/discharge.
[0034] This electrical formation requirement puts an additional
financial burden on commercial battery manufacturers. That is, it
requires the manufacturers to purchase capital equipment in the
form of battery chargers and also requires the cost of labor and
utilities to run the equipment. These costs are significant and are
passed on to the consumer. Therefore, there remains a need in the
art for an electrochemical hydrogen storage alloy which requires
little or no electrical formation.
[0035] The chemical/thermal activation of the electrochemical
hydrogen storage alloys involves a relatively lengthy period of
immersing the alloy material (in powder or electrode form) into a
concentrated potassium hydroxide or sodium hydroxide solution,
preferably at an elevated temperature. In situ treatment of the
electrodes in the battery is limited to a temperature of about
60.degree. C. because of the separators used therein. In powder
form, the temperature limit is higher. The normal maximum
concentration of potassium hydroxide is about 30% by weight KOH in
water. The required residence time depends on temperature and
concentration, but is typically a few days for the finished
batteries. Information on chemical/thermal activation of
electrochemical hydrogen storage alloys is provided in the '088
patent. This again is another added cost for the manufacturer. The
costs of raw materials such as KOH or NaOH and water, the cost of
disposing of spent chemicals, the energy costs to heat the alloy
materials and the KOH solution, the labor and inventory costs, and
the time costs all make it desirable to reduce or eliminate this
activation process. Therefore, there is a need in the art to
develop an electrochemical hydrogen storage alloy which requires
little or no chemical/thermal activation.
[0036] Additionally, prior art alloys have all been designed for
ultimate capacity, and have not been designed for the high rate
requirements of HEV use and the like. Prior art VTiZrNiCrMn alloys
have a specific capacity of 380-420 mAh/g in electrode form.
Recently, there has been increased demand for rechargeable
batteries having higher power and rate capabilities in addition to
the high energy.
[0037] Finally, prior art alloys having high electrochemical
storage capacity have lacked a very highly catalytic surface. Some
prior art alloys had a catalytic surface, but it was limited.
Higher catalytic activity allows for higher exchange currents and
thereby higher rate capabilities. Also, the surface area of the
alloy affects the exchange current. That is, the higher the surface
area, the greater the exchange current. Therefore, there is a need
in the art for electrochemical hydrogen storage alloys which have a
greater surface catalytic activity as well as a greater surface
area.
SUMMARY OF THE INVENTION
[0038] The above deficiencies are remedied by a modified
Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen storage alloy which
has at least one of the following characteristics: 1) an increased
charge/discharge rate capability over that the base
Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen storage alloy; 2) a
formation cycling requirement which is significantly reduced over
that of the base Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy; or 3) an oxide surface layer having a higher
electrochemical hydrogen storage catalytic activity than the base
Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen storage alloy.
[0039] The modified Ti--V--Zr--Ni--Mn--Cr electrochemical hydrogen
storage alloy comprises, in atomic percentage: (Base
Alloy).sub.aCo.sub.bFe.sub.c- Al.sub.dSn.sub.e, where said Base
Alloy comprises 0.1 to 60% Ti, 0.1 to 40% Zr, 0 to 60% V, 0.1 to
57% Ni, 5to 22% Mn and 0 to 56% Cr; b is 0.1 to 10.0%; c is 0 to
3.5%; d is 0.1 to 10.0%; e is 0.1 to 3.0%; and a+b+c+d+e=100%.
BRIEF DESCRIPTION OF THE DRAWINGS
[0040] FIG. 1 depicts Electrochemical capacity of 16 alloys at
discharge currents of 50 and 12 mA/g, without alkaline etching;
[0041] FIG. 2 plots the peak power in W/Kg versus cycle number for
alloys 01, 02, 03, 04, 05, 12, and 13;
[0042] FIGS. 3a-3c plot cell capacity, number of electrical
formation cycles and peak power at 50% depth of discharge (at room
temperature), respectively for alloys 01 and 12; and
[0043] FIG. 4, plots the specific power (W/Kg) at 50 and 80% depth
of discharge for alloys 01 and 12;
[0044] FIGS. 5a and 5b depict the effect of alkaline etching on
alloy-01, specifically depicted are plots of capacity versus cycle
number at 12 and 50 mA/g discharge, respectively;
[0045] FIGS. 6a and 6b depict the effect of alkaline etching on
alloy-012, specifically depicted are plots of capacity versus cycle
number at 12 and 50 mA/g discharge, respectively;
[0046] FIG. 7 plots AC impedance (Nyquist plots) at 85% state of
charge (SOC) of thermal/chemical activated negative electrodes
prepared from base alloy-01 and from alloy-12 of the instant
invention;
[0047] FIG. 8 plots discharge curves at 2C rate for C-cells
manufactured using negative electrodes fabricated from base
alloy-01 and from alloy-12 of the instant invention, it should be
noted that the higher power capability demonstrated by alloy-12 in
half cell measurements was also reflected in the performance of the
C-cells;
[0048] FIG. 9 shows the half cell capacity as a function of
discharge rate, again, the electrodes using alloy-12 electrodes
exhibited better rate capability; and
[0049] FIG. 10 plots the pressure-concentration isotherm (PCT)
curve of alloy-12 cooled in a cylindrical mold verses a flat slab
shaped mold which has a higher overall cooling rate.
DETAILED DESCRIPTION OF THE INVENTION
[0050] The deficiencies of the prior art are overcome by the
instant modified VTiZrNiCrMn electrochemical hydrogen storage
alloy. In order to improve the catalytic activity of the prior art
negative hydride alloys, the base alloy material was modified by
the addition of one or more elements to increase the surface area
of the heat-activated alloys and to enhance the catalytic nature of
the surface of the materials. In addition to VTiZrNiCrMn, the
alloys also contain Al, Co, and Sn. The alloy has an increased
charge/discharge rate capability over that of the base VTiZrNiCrMn
electrochemical hydrogen storage alloy. It also has an electrical
formation cycling requirement which is reduced to one tenth that of
the base VTiZrNiCrMn electrochemical hydrogen storage alloy. A
chemical/thermal activation is required by the base VTiZrNiCrMn
electrochemical hydrogen storage alloy. Finally the alloy has an
oxide surface layer having a higher electrochemical hydrogen
storage catalytic activity and higher surface area than the base
VTiZrNiCrMn electrochemical hydrogen storage alloy. Each of these
properties will be discussed in detail hereinbelow.
[0051] Catalytic Activity
[0052] The '591 patent, discussed hereinabove, represents the best
prior art teaching of the desirable properties of the
metal/electrolyte interface, or surface oxide of the metal hydride
material, providing specific teaching on the role of metallic
nickel sites as catalytic sites. The '591 patent also describes the
nickel sites as approximately 50-70 Angstroms in size, with a broad
proximity range of 2-300 Angstroms. With respect to proximity, the
STEM micrographs provided suggest approximately 100-200 Angstrom
proximity as following the teaching of the '591 patent.
[0053] To distinguish the alloys of the present invention over
those of the '591 patent, the inventors have discovered superior
catalysis and high rate discharge performance can be achieved by
one or more of the following:
[0054] 1) the catalytic metallic sites of the inventive alloys are
formed from a nickel alloy such as NiMnCoTi rather than just
Ni;
[0055] 2) the catalytic metallic sites of the inventive alloys are
converted by elemental substitution to an FCC structure from the
BCC structure of the prior art Ni sites;
[0056] 3) the catalytic metallic sites of the inventive alloys are
much smaller in size (10-50, preferably 10-40, most preferably
10-30 Angstroms) than the Ni sites of the prior art alloys (50-70
Angstroms) and have a finer distribution (closer proximity);
[0057] 4) the catalytic metallic sites of the inventive alloys are
surrounded by an oxide of a multivalent material (containing
MnO.sub.x) which is believed to possibly be catalytic as well, as
opposed to the ZrTi oxide which surrounded the prior art Ni
sites;
[0058] 5) the oxide could also be multiphase with very small (10-20
Angstrom) Ni particles finely distributed in a MnCoTi oxide
matrix;
[0059] 6) the oxide may be a mix of fine and coarse grained oxides
with finely dispersed catalytic metallic sites;
[0060] 7) alloy modification with aluminum may suppress nucleation
of large (50-70 Angstrom) catalytic metallic sites (at 100 Angstrom
proximity) into a more desirable "catalytic cloud" (10-20 Angstroms
in size and 10-20 Angstroms proximity);
[0061] 8) NiMn oxide is the predominant microcrystalline phase in
the oxide and the catalytic metallic sites may be coated with NiMn
oxide.
[0062] The instant alloys, therefore, distinguish over the '591
alloys in that: 1) the catalytic metallic sites are still present
but may be nickel alloy and are much smaller and more finely
divided; 2) the old TiZr oxide support is replaced by a NiMnCoTi
oxide which is more catalytic and more porous; and 3) aluminum
metal doping provides a very fine grain catalytic metallic site
environment.
EXAMPLE I
[0063] Sn, Co, Al, and Fe were considered as additives to a base
AB.sub.2 alloy. Sixteen different chemical formulas were designed
according to the orthogonal array used in the Taguchi method to
minimize the total number of alloys needed to complete the design
matrix. Each element has four different levels; i.e., Sn (0.4, 0.6,
0.8, 1.0), Co (0, 0.5, 1.0, 1.5), Al (0, 0.4, 0.8, 1.2), Fe (0,
0.4, 0.8, 1.2), as shown in Table 1 (all numbers are in atomic
percentages). Alloy-01 is the base formula (control) with only 0.4%
Sn originating from one of the source materials (zircalloy in
replacement of zirconium) to reduce raw materials cost.
1TABLE 1 Alloy Element Concentration (Atom) Capacity # Sn Co Al Fe
V Ti Z (mAh/g) 01 0.4 0.0 0.0 0.0 5.0 9.0 26.6 38.0 5.0 16.0 390 02
0 4 0.5 0.4 0.4 5.0 9.0 26.6 38.0 4.5 15.2 382 03 0.4 1.0 0.8 0.8
5.0 9.0 26.6 38.0 4.0 14.4 375 04 0.4 1.5 1.2 1.2 5.0 9.0 26.6 38.0
3.5 13.6 375 05 0.6 0.0 0.4 0.8 5.0 9.0 26.4 38.0 5.0 14.8 379 06
0.6 0.5 0.0 1.2 5.0 9.0 26.4 38.0 4.5 14.8 387 07 0.6 1.0 1.2 0.0
5.0 9.0 26.4 38.0 4.0 14.8 376 08 0.6 1.5 0.8 0.4 5.0 9.0 26.4 38.0
3.5 14.8 389 09 0.8 0.0 0.8 1.2 5.0 9.0 26.2 38.0 5.0 14.0 401 10
0.8 0.5 1.2 0.8 5.0 9.0 26.2 38.0 4.5 14.0 374 11 0.8 1.0 0.0 0.4
5.0 9.0 26.2 38.0 4.0 15.6 370 12 0.8 1.5 0.4 0.0 5.0 9.0 26.2 38.0
3.5 15.6 385 13 1.0 0.0 1.2 0.4 5.0 9.0 26.0 38.0 5.0 14.4 369 14
1.0 0.5 0.8 0.0 5.0 9.0 26.0 38.0 4.5 15.2 369 15 1.0 1.0 0.4 1.2
5.0 9.0 26.0 38.0 4.0 14.4 335 16 1.0 1.5 0.0 0.8 5.0 9.0 26.0 38.0
3.5 15.2 339
[0064] All sixteen alloys were prepared by induction melting under
an argon atmosphere with commercially available raw materials. The
melt size ranged from 20 to 60 kg depending on the crucible size
been used. After reaching 1600.degree. C., the melt was held at
that temperature for 20 minutes to homogenize it. Afterwards, the
liquid was cooled down to 1300.degree. C. and tilt-poured into a
carbon steel mold. The ingots thus obtained were pulverized by a
hydride/dehydride process without mechanical grinding as indicated
in U.S. application Ser. No. 09/141,668, filed Aug. 27, 1998,
entitled A METHOD FOR POWDER FORMATION OF A HYDROGEN STORAGE
MATERIAL, herein incorporated by reference. Powder of 200 mesh or
smaller was roll-milled onto a Ni-mesh substrate without other
conducting metal powder or inorganic additives. The electrochemical
capacity of each alloy was determined by constructing a flooded
full cell using grafted PE/PP separators, partially pre-charged
Ni(OH).sub.2 counter electrodes, and 30% KOH solution as the
electrolyte The cells were charged at 50 mA/g for 13 hours and then
discharge at 50 mA/g and a final pull current at 12 mA/g.
[0065] The discharge capacity for the third cycle at 50 and 12 mA/g
for each alloy are plotted in FIG. 1. This figure indicates that
alloy-12 shows the smallest differential between capacities at 50
and 12 mA/g, which indicates a good high-rate material.
[0066] Electrical Formation Cycling
[0067] Electrical formation or workup cycling of NiMH type
batteries, as discussed herein above, has previously been a
requirement for alkaline type batteries. This electrical formation
was required to bring the battery to full capacity and especially
full power. Without such formation, the batteries perform below
maximum capability. Typical formation entails cycling the virgin
battery many times at differing charge/discharge rates. For
example, the base alloy, having a nominal composition (in atomic %)
Ti 9.0%, Zr 27.0%, V 5.0%, Ni 38.0%, Cr 5.0%, and Mn 16.0%,
required 32 charge/discharge cycles to achieve full power.
Particularly in the area of electric vehicles, where power
translates into acceleration of the vehicle, formation cycling is
an expensive process with respect to equipment, processing time and
inventory control. Any reduction in the number of cycles required
to form the battery to it's full capability reduces the cost of
manufacturing.
[0068] The instant alloy materials have been specifically designed
to speed up formation. To that end, the instant alloy materials
have reduced the electrical formation requirement thereof to just
three cycles in consumer, cylindrical cells and also EV batteries.
This reduction in formation cycles is ten fold over the prior art
base alloy. Therefore, production times and costs are reduced, and
throughput is increased.
EXAMPLE II
[0069] C-size cylindrical batteries were constructed using the
alloys fabricated from example I as negative electrode. These cells
included paste Ni(OH).sub.2 as the positive electrodes and 30% KOH
solution as electrolyte. The peak power of the battery was measured
by the pulse discharge method and the results of a few key alloys
are plotted in FIG. 2 as a function of cycle number. It is clear
from the figure that alloys -02, -03, -04, -05, -12, and -13 all
have higher peak power than the control (alloy-01). Especially
alloy-12 which reached it's full rate capability after only three
electrical formation cycles. This is a dramatic improvement over
alloy-1 for which more than 15 cycles are needed.
EXAMPLE III
[0070] Both electrodes from alloy-01 and alloy-12 were made into
identical prismatic cells for electrical vehicle application (90 Ah
by design). Testing results for these cells are summarized in FIGS.
3a, 3b and 3c. Both cells reached their designed capacity and power
after 5 days of heat treatment at 60.degree. C. and various number
of mini-cycles for electrical formation. Cell employing alloy-12
showed marginal advantages in both capacity and power. However, the
most significant finding is that the number of mini-cycles needed
to achieve the maximum power was dramatically reduced from 39
(alloy-01) to 9 (alloy-12), which offers a substantial cost
reduction in capital equipment and electricity.
[0071] The electrical formation was further studied to take full
advantage of alloy-12. Instead of the typical 37 hours of
electrical formation for alloy-01, the whole formation process can
be reduced to 12 hours by using alloy-12. The final capacity and
specific power were not affected by this aggressive formation
scheme, ash show in FIG. 4.
Chemical/Thermal Activation
EXAMPLE IV
[0072] All sixteen alloys obtained from Example I were examined
after various etching conditions. This alkaline etch was designed
to simulate the heat formation process during battery fabrication.
Electrodes were cut into proper size (2 by 5 inches) and etched in
a 100.degree. C. 30% KOH solution for 1, 3, and 4 hours. Etched
electrodes together with unetched electrodes were used to construct
flooded full cells using graft PE/PP separators, partially
pre-charged Ni(OH).sub.2 counter electrodes, and 30% KOH solution
electrolyte. The cells were charged at 50 mA/g for 13 hours and
then discharge at 50 mA/g and a final pull at 12 mA/g. The
capacities under various etching conditions are plotted as a
function of cycle number in FIGS. 5a and 5b for the alloy-01 and
FIGS. 6a and 6b for alloy-12, respectively. It is found that
alloy-12 is easier to form (reaching full capacity and rate
capability within fewer cycles) when compared to the alloy-01.
[0073] Catalytic Activity & Rate Capability
[0074] The instant alloy materials have far outdistanced Misch
metal nickel based metal hydride alloys. The rate of catalytic
surface activity of the instant alloys and rate of bulk diffusion
of hydrogen are similar. Therefore, neither process inhibits the
rate of charge/discharge, when compared to the other. In fact,
because of the improvements in catalytic surface activity, the
instant alloys have much improved discharge rate capability, i.e.
as much as 300% greater rate capability. This, as will be discussed
further herein below, appears to be caused by the enhanced oxide
layer of the instant alloy materials.
[0075] Electrochemical studies were conducted to characterize the
newly developed derivative alloys and compare their properties with
alloy-01 material. The study helped to better understand the nature
of the changes occurring at the surface of alloy-01 as a result of
the compositional and structural modifications and the relation of
these changes to the increased catalytic activity and rate
capability of the material.
AC Impedance--Surface Kinetic and Diffusion Properties
[0076] FIG. 7 shows AC impedance plots (Nyquist plots) at 85% state
of charge (SOC) of thermal/chemical activated negative electrodes
prepared from the alloy-01 and alloy-12. The main semicircle in the
impedance plots of FIG. 7 is due to the charge transfer which
occurs at the surface of the MH electrode. The hydrogen species
formed at this step are adsorbed to the electrode's surface. The
diameter of this circle represents the charge transfer resistance
R.sub.ct of the hydride reaction. At frequencies lower than that of
the charge transfer semicircle, the impedance is attributed to the
absorption of the hydrogen below the surface of the metal. This
step is followed by the bulk diffusion step in which the absorbed
hydrogen species diffuse into the bulk of the metal hydride
material. The absorption step gives rise to a small semicircle at
the lower frequency range of the impedance plots and the bulk
diffusion step gives rise to the straight, Warburg, behavior
observed at lower frequency range of the impedance plots. Following
the Warburg region, the impedance turns into a 90.degree.
capacitive line due the fact that the hydrogen diffusion occurs
through a finite length.
[0077] The impedance behavior shown in FIG. 7 therefore support a
three step mechanism as described in the following equations:
2 (A) M + H.sub.2O + e MH.sub.ad + OH.sup.- Charge Transfer (B)
MH.sub.ad .sub.MhabS Absorption (C) MH.sub.abS .sub.MhabB Bulk
Diffusion
[0078] Where MH.sub.ad is the adsorbed hydrogen, MH.sub.abS is the
absorbed hydrogen just below the surface and MH.sub.abB is the
absorbed hydrogen in the bulk of the negative material. The charge
transfer step (A) controls the impedance of the electrode at the
high frequency range. At lower frequencies, the bulk diffusion
process dominates the impedance.
Surface Kinetics
[0079] The surface kinetics of the hydride reaction is measured by
the charge transfer resistance (R.sub.ct) or the exchange current
(I.sub.0). I.sub.0 is related to R.sub.ct by the equation:
I.sub.0=RT/nFR.sub.ct (1)
[0080] Table 2 shows the charge transfer resistances obtained from
the impedance plots of FIG. 7 and the exchange currents for
alloy-01 and alloy-12. As Table 2 shows the exchange current for
the 2.7% Co, Al, Sn modified alloy-12 is 2-3 times larger than that
of alloy-01, indicating faster charge transfer kinetic by the same
proportion.
3 TABLE 2 Alloy R.sub.d(ohm g) Alloy-01 0.4 65 Alloy-12 0.17
155
[0081] The magnitude of the exchange current (I.sub.0) is generally
determined by the catalytic activity of the electrode surface
measured by the exchange current density (i.sub.0) and by the
specific surface area (A) of the electrode. In order to understand
better how each of these parameters contribute to the increase in
charge transfer kinetics observed in the derivative alloys, the
values of i.sub.0 and A for the different electrodes were also
measured. The double layer capacitance (C.sub.dl) of the electrodes
calculated from the ac impedance plots of FIG. 7 were used to
determine the surface area A of the different electrodes. To
calculate the surface area from C.sub.dl, a specific capacitance of
20 uF/cm.sup.2 (a common literature value) was assumed. The
exchange current densities of the different electrodes were
calculated using the relationship
i.sub.0=I.sub.0/A.times.100 (2)
[0082] where i.sub.0 in equation 2 is in mA/cm.sup.2, I.sub.0 is in
mA/g and A is in m.sup.2/g. Table 3 shows the values of I.sub.0,
C.sub.dl, A and i.sub.0 for alloy-01 and alloy-12.
4 TABLE 3 10 .times. Alloy l.sub.a(mA/g) C.sub.d(g)
l.sub.g(mA/cm.sup.2) Alloy-01 65 0.16 0.8 0.8 Alloy-12 155 0.26 1.3
1.2
[0083] As Table 3 shows, both the surface area and the exchange
current densities are higher in alloy-12 electrodes as compared to
alloy-01 electrodes. Calculations from Table 3 show that for
alloy-12 about 50% of the increase in the exchange current
(I.sub.0) can be attributed to the higher surface area of these
materials and about 50% of the increase can be attributed to the
higher exchange current densities of these materials as compared to
alloy-01. The higher surface area of the alloy-12 electrodes with
respect to alloy-01 electrodes occurred during the heat activation
of the electrodes since ac impedance measurements showed that the
surface area of these electrodes in the virgin state were similar
to each other. The heat activation process helps to increase the
surface area of the electrodes. The results presented here show
that the modifying elements added to alloy-01 serve an important
role in creating higher surface area due to their dissolution at
the surface during the heat activation process. The higher exchange
current densities of alloy-12 as compared to the base alloy
indicates the added elements not only contributed to the increase
in surface area of the electrodes but also contribute to the
enhancement in the catalytic nature of the materials by changing
the surface composition.
[0084] Though not wishing to be bound by theory, it is possible the
oxide surface area is substantially increased by a mechanism
similar to that in which Co, Al, and Sn modification causes the
nickel catalytic sites to be reduced from 50-70 Angstroms to 10
Angstroms. It is possible in the VTiZrNiCrMn prior art material,
that two mechanisms control 50-70 Angstroms size sites. First,
microcrystallite size within the bulk alloy may inherently
influence resultant Ni sites after oxidation of the remaining
elements. It is more likely, however, that the dissolution, erosion
and corrosion of the resultant oxide is influenced by its chemical
makeup. For example, in a prior art oxide matrix dominated by TiZr
oxide, which is relatively insoluble, corrosion/erosion may occur
in large chunks of 50-70 Angstrom size, while the Sn, Al, Co
modified materials may be corroding on an atomic basis on the order
of 10 Angstroms. Al and Sn may be particularly crucial in this
regard in that they may be bonded to Ni within the surface oxide
but dissolve in a "finer" or less "chunky" manner than VTiZr
oxides.
[0085] Bulk Diffusion
[0086] From the ac impedance plots of FIG. 7, the diffusion rate of
hydrogen in the bulk material can also be determined. While the
surface kinetics determine the power capability of the electrode
and batteries which use them, the diffusion rate determines the
rate capability. The diffusion rate of the hydrogen species is
reflected in the impedance of the electrodes at the lower frequency
range of the impedance plots of FIG. 7. R.sub.d is related to the
diffusion coefficient of the hydrogen species (D.sub.H) and to the
diffusion length (l) by equation 3:
R.sub.d= V.sub.M/ZFa(dE/dy)(l/3D) (3)
[0087] where V.sub.M is the molar volume of the electrode material,
Z is the charge per hydrogen atom absorbed, F is Faraday number, a
is the geometric surface area of the electrode and (dE/dy) is the
change of equilibrium potential of the electrode per unit change of
hydrogen absorption. This parameter was obtained from measured data
of equilibrium potential versus state of charge of the electrodes
and was calculated to be approximately 0.06V at a state of charge
between 85% to 50%. D and I are the diffusion coefficient and
diffusion length respectively. Assuming a diffusion length equal to
the electrode thickness, the hydrogen diffusion coefficient of the
different alloy materials could be calculated. Table 4 shows the
diffusion resistance obtained from the Nyquist plot of FIG. 7 and
the diffusion coefficient calculated for alloy-01 and alloy-12. The
diffusion coefficient of alloy-12 is larger, giving rise to a
proportionally higher diffusion rate and better rate capability for
alloy-12.
[0088] Though not wishing to be bound by theory, the mechanism
through which the Sn, Co, Al modified alloys may have improved bulk
hydrogen diffusion rates may be related to one or more of the
following:
[0089] 1. refinement of microstructure towards smaller crystallite
sizes, which in turn promotes grain boundaries and hydrogen
transport;
[0090] 2. higher hydrogen equilibrium pressure; or
[0091] 3. finer dispersion of catalytic sites within the bulk
(similar to surface oxide).
5 TABLE 4 Alloy R.sub.d(ohm) (cm/sec) Alloy-01 1.4 3.9 .times.
10.sup.-8 Alloy-12 0.97 5.6 .times. 10.sup.-8
[0092] Cell Performance
[0093] The performance of the different negative electrodes was
also studied in cylindrical C-cells. FIG. 8 shows discharge curves
at 2C rate of C-cells manufactured using negative electrodes
fabricated from alloy-01 and from alloy-12. As FIG. 8 shows, the
cells using alloy-12 exhibited higher operating voltages reflecting
superior power capability. FIG. 9 shows the capacity of electrodes
fabricated from alloys -01 and -12. C-cells using the alloy-12
electrodes exhibited better rate capability.
[0094] Analytical Study--The Oxide Surface
[0095] The oxide surface of the instant alloys is the same
thickness as that of the prior art alloys, however, the instant
inventors have noticed that the modification of the alloys has
affected the oxide surface in several beneficial ways. First the
oxide accessibility has been affected. That is, the additives to
the alloy have increased the porosity and the surface area of the
oxide. This is believed to be caused by Al, Sn and Co. The
modifiers added to the alloy are readily soluable in the
electrolyte, and believed to "dissolve" out of the surface of the
alloy material, leaving a less dense, more porous surface into
which the electrolyte and ions can easily diffuse.
[0096] Second, the inventors have noted that the derivative alloys
have a higher surface area than the prior art alloys. It is
believed that the mechanical properties of the alloy (i.e.
hardness, ductility, etc.) has been affected. This allows the
material to be crushed easier, and allows for more microcracks to
be formed in the alloy material during production and also easier
in-situ formation of microcracks during electrochemical
formation.
[0097] Finally, the inventors have noted that the alloys are more
catalytically active than the prior art alloys. This is believed to
be cause by a more catalytic active oxide surface layer. This
surface layer, as is the case with some prior art materials (see
for example U.S. Pat. No. 5,536,591 to Fetcenko et al.,) includes
nickel particles therein. These nickel particles are believed to
provide the alloy with its surface catalytic activity. In the
instant alloy, the inventors believe there are a number of factors
causing the instant increase in catalytic surface activity. First,
the inventors believe that the nickel particles are smaller and
more evenly dispersed in the oxide surface of the instant alloy
materials. The nickel particles are believed to be on the order of
10 to 50 Angstroms in size. Second, the inventors believe that the
nickel particles may also include other elements such as cobalt,
manganese and iron. These additional elements may enhance the
catalytic activity of the nickel particles, possibly by increasing
the roughness and surface area of the nickel catalytic sites
themselves. Third, the inventors believe that the oxide layer
itself is microcrystalline and has smaller crystallites than prior
art oxide. This is believed to increase catalytic activity by
providing grain boundaries within the oxide itself along which
ions, such as hydrogen and hydroxyl ions, may move more freely to
the nickel catalyst particles which are situated in the grain
boundaries. Finally, the instant inventors have noted that the
concentrations of cobalt, manganese and iron in the oxide surface
are higher than in the bulk alloy and higher than expected in the
oxide layer.
[0098] The surface area of the base alloy increased by about a
factor of two during the activation treatment, the inventive alloy
increases in surface area by about a factor of four. As discussed
earlier, the higher surface area of the inventive alloy is only
partially responsible for the higher catalytic property of these
alloys. As the ac impedance measurements demonstrated, the better
catalytic activity of the surface of the inventive alloy also
contributes to the enhanced catalytic behavior thereof.
[0099] Hence, the improved power and rate capability of the
inventive alloys is the result of the higher surface area within
the surface oxide as well as improved catalytic activity within the
oxide due to the smaller size and finer dispersion of the nickel
catalyst particles compared to prior art materials. Observations
from high resolution scanning transmission electron microscopy
(STEM) included presence of nickel catalyst "clouds" having a size
in the 10-30 Angstrom range and extremely close proximity, on the
order of 10-20 and 10-50 Angstrom distance. Another contributing
factor to the improved catalysis within the oxide is the
transformation of the supporting oxide in which the Ni particles
reside. In prior art materials, the supporting oxide may be
primarily rare earth or TiZr based oxides while in the case of the
inventive materials, the support oxide is now comprised of at least
regions of NiCoMnTi "super catalysts." This could also be NiMn
regions surrounded by TiZr oxide. These super catalysts show a
surprising lack of oxygen based on Electron Energy Loss
Spectroscopy (EELS). It may be possible these regions are partially
metallic or in a low oxidation state.
[0100] Another observation with the inventive materials is that
prior art nickel catalytic regions within the oxide were bcc
crystallographic orientation based on Select Area Electron
Diffraction (SAED), which the inventive materials were observed to
have an fcc orientation. It may be possible that the catalytic
regions of Ni have been partially substituted by Co, Al, Mn, Sn, or
other elements which have shifted the crystallographic orientation.
It is indeed likely the bcc to fcc Ni shift reflects a higher
degree of substitution.
[0101] Though not wishing to be bound by theory, it is also
possible the fcc Ni in conjunction with NiCoMnTi regions and TiZr
oxide may form a super lattice which may further promote ionic
diffusion and reaction. Still another theory based on analytical
evidence suggests that metallic Ni particles reside in a Mn oxide
support. The presence of the Mn oxide is intriguing in that
MnO.sub.x is multivalent and could promote catalysis via changing
oxide states during the charge/discharge reactions.
[0102] Finally, another interpretation of the analytical evidence
suggests even a multiphase surface oxide. In addition to metallic
Ni or Ni alloys, there appears to exist both a fine grained and
coarse grained support oxide. Perhaps the course grained aspect to
the surface is dominated by TiZr prior art style oxide while the
appearance of the fine grained support oxide in the inventive
materials may be the MnOx or NiMnCoTi oxide or a MnCoTi oxide. The
difficulty in assigning these structures more specifically resides
in the very invention itself, i.e. the extremely small size and
fine distribution. Even state of the art analytical instruments
using electron probes, etc., have some kind of analytical region
where averaging is taking place. Difficulty in assignment is mainly
due to overlap of these extremely fine regions with one another
during analysis.
[0103] In this context, one key role of Al, Sn, Co modification in
these alloys may be as a "poison to the surface", inhibiting the
growth of large Ni particles. In other words, these specific
dopants may be viewed as metallic catalysts and support oxide
dispersants.
EXAMPLE V
[0104] Performance of negative electrodes produced with alloys of
the instant invention can be further optimized by adjusting the
melt-casting conditions of the alloys. For example, a flat slab
mold was used to increase the quench rate during casting relative
to the conventional cylindrical mold used in Example 1. The ingots
obtained from the slab mold have an average thickness of less than
about 5 inches and preferably less than about one inch as compared
with 10 inch thick ingots obtained from the cylindrical mold.
[0105] The pressure-concentration isotherm (PCT) curves of alloy-12
from both casting methods are plotted in FIG. 10. From this figure,
one can easily identify the superiority of faster solidification
via the slab mold as compared with the slower cooling of the
cylindrical mold by the extended curve into higher hydrogen
storage. Electrodes were formed from alloy material from both
ingots and subjected to half-cell testing as taught in Example 1.
The results, which are listed in Table 5, confirm the PCT
predictions about capacity. Not only did the full capacity increase
from 355 mAh/g for the cylindrical mold to 395 mAh/g for the slab
mold, but the capacity loss at high rate discharge was less for the
slab mold. These increases in capacity can be directly related to
higher power in the finished battery.
6 TABLE 5 Mold Type Capacity Cylindrical 329 mAh/g 355 mAh/g Slab
376 mAh/g 395 mAh/g
* * * * *