The effects of alloy chemistry on the el

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The effects of alloy chemistry on the electrochemical and hydriding properties of NI-substituted LaNi5
Citation
Witham, Charles Kincaid
(2000)
The effects of alloy chemistry on the electrochemical and hydriding properties of NI-substituted LaNi5.
Dissertation (Ph.D.), California Institute of Technology.
doi:10.7907/w3e0-5779.
Abstract
NOTE: Text or symbols not renderable in plain ASCII are indicated by [...]. Abstract is included in .pdf document.

The primary goal of this work was to verify the hypothesis that alloying LaNi5 with ternary elements that have a large heat of formation with La (and secondarily, with Ni) would slow the kinetics of metal (La) atom diffusion. This would have the effect of stabilizing the Haucke phase crystal structure of LaNi5 during electrochemical and gasphase hydrogen absorption/desorption cycling, and extending the material's useful lifetime.

To test this hypothesis, I prepared a variety of single-phase alloys of composition [...]. Each alloy was annealed to insure equilibrium starting conditions. The lifetimes of these alloys were tested by charge-discharge cycling as the anode of an alkaline Ni-MH rechargeable cell. By characterizing the lifetimes of the alloys as an exponential capacity decay, I was able to determine a trend between the capacity decay and the heat of formation of an average 'B' element with La.

The alloys were further characterized by obtaining gas-phase isotherms and, in the case of the [...] alloys, the thermodynamics of metal hydride formation and decomposition. X-ray diffraction was used to measure the effect of substitution on the alloy and its hydride phase. By examining the data obtained at Caltech as well as data published in the literature, several trends were noted. There is a fairly linear relationship between the solute's expansion of the LaNi5 unit cell and its radius. The total volume expansion an alloy experienced upon absorption of hydrogen was found to decrease with substituted composition. The discrete lattice expansion of [...] alloys was found to to decrease substantially for 0 < x < 0.2, but subsequent substitution had little effect on the volume expansion.

The electrode electrochemical kinetice of charge transfer were investigated for each MH alloy. Measurements of the charge transfer exchange current by micropolarization and AC impedance were similar, while those measured by Tafel polarization did not have a clear relationship.
Item Type:
Thesis (Dissertation (Ph.D.))
Subject Keywords:
Materials Science
Degree Grantor:
California Institute of Technology
Division:
Engineering and Applied Science
Major Option:
Materials Science
Thesis Availability:
Public (worldwide access)
Research Advisor(s):
Fultz, Brent T.
Thesis Committee:
Unknown, Unknown
Defense Date:
24 August 1999
Record Number:
CaltechETD:etd-12272004-145717
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DOI:
10.7907/w3e0-5779
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THE EFFECTS OF ALLOY CHEMISTRY
ON THE ELECTROCHEMICAL AND HYDRIDING PROPERTIES

OF NI-SUBSTITUTED LANI,

Thesis by
Charles K. Witham

In Partial Fulfillment of the Requirements
for the Degree of
Doctor of Philosophy

California Institute of Technology

Pasadena, California

2000
(Submitted August 24th, 1999)

c 2000
Charles K. Witham

iii

Acknowledgments

There were many people who were instrumental in the completion of my graduate
work and this thesis who I would like to thank here. These are by no means listed in
order of importance, as I could not have accomplished as much as I did without each of
them. My advisor Prof. Fultz has been of utmost utility in training me in the use of
laboratory equipment, in applying my own insight, both in and out of the classroom, and
in familiarizing me with the mechanisms of successfully maneuvering in both the
University setting and the current academic funding environment.

Dr. Bob Bowman was particularly helpful in freely transferring to me his vast
knowledge of metal hydride systems, and a bit of his cynicism. He also helped to expose
me to the scientific community and a side of the metal hydride community that I might
never have seen without his introductions. |

In regards to the electrochemical tests, I must acknowledge and show my great
appreciation to Dr. Ratnakumar Bugga of the Jet Propulsion Laboratory Electrochemical
Technologies Group and Adrian Hightower of Caltech. Dr. Bugga trained me on all of
the equipment relevant to the testing of metal hydride alloys in alkaline batteries, and
imparted to me much of what I know about electrochemistry. My fellow graduate student
Adrian Hightower performed some of the tests of the alloys’ electrochemical kinetics.

The work presented in this thesis was carried out under funding by the Department of
Energy, grant DE-FG03-94ER 14493.

And I cannot forget my undergraduate professors, John Hack, Bob Weber, Dave
Keenes, and Alessandro Gomez, who all encouraged me to attend graduate school and
helped give me the courage to achieve.

Finally, I truly appreciate that all of the people who were involved in my work have

also extended to me their friendship.

Abstract —

The primary goal of this work was to verify the hypothesis that alloying LaNi, with
ternary elements that have a large heat of formation with La (and secondarily, with Ni)
would slow the kinetics of metal (La) atom diffusion. This would have the effect of
stabilizing the Haucke phase crystal structure of LaNi, during electrochemical and gas-
phase hydrogen absorption/desorption cycling, and extending the material’s useful
lifetime.

To test this hypothesis, I prepared a variety of single-phase alloys of composition
LaNi,,M,, 0 < x < 0.5, Me {Al, Si, Ga, Ge, In, Sn}. Each alloy was annealed to insure
equilibrium starting conditions. The lifetimes of these alloys were tested by charge-
discharge cycling as the anode of an alkaline Ni-MH rechargeable cell. By characterizing
the lifetimes of the alloys as an exponential capacity decay, I was able to determine a
trend between the capacity decay and the heat of formation of an average “B’ element
with La

The alloys were further characterized by obtaining gas-phase isotherms and, in the
case of the Ge, alloys, the thermodynamics of metal hydride formation and
decomposition. X-ray diffraction was used to measure the effect of substitution on the
alloy and its hydride phase. By examining the data obtained at Caltech as well as data
published in the literature, several trends were noted. There is a fairly linear relationship
between the solute’s expansion of the LaNi, unit cell and its radius. The total volume
expansion an alloy experienced upon absorption of hydrogen was found to decrease with
substituted composition. The discrete lattice expansion of Ge, alloys was found to
decrease substantially for 0 < x < 0.2, but subsequent substitution had little effect on the

volume expansion.

a clear relationship.

Table of Contents

L. Fntroduction.......... 2... cccceceeeec eee eneeeseceecenceseeeeeeeeeneeseeseneesesenseene ees 1
If. Background
A. Short History of Metal Hydrides.................c ccc ece sees ese ee eeee cent ee eees 3

references 8

B. Applications...............cccece eee cc eee eee eee eee e neces eeneeee eases ease sereeneeees 9

references 15

C. LaNi,
1. Phase Diagram and Crystal Structure of LaNi,..............::ceceeeeee eens 18
2. Gas-Phase p-c-isotherm............... cece eee e eee e eee eeeeene eee ceeeneeesseenaees 21
3. Hydride Structure and Lattice Expansion..................cceceeeeeeeeeeeees 27
4. Basic Electrochemical Properties..................ccceseceeeeeeeeeeeneee eens 36
5. Microstrains.............ccecccceceee cece nee ne eee ee ses enseseneesee eee ener nena ees 39

D. Alloy Modifications. ...............cc ccc ccc e cece eee nese ee eeeetenteneen senses anes 47
1. Phase Diagram and Metal-Atom Crystal Structure...............::.00e08 47
2. ISOtherm.............ccecece ese eeceeeeeee ees ceeeeeeensensent cesses reeesnstees eens 51
3. Hydride Structure and Lattice Expansion.................ceceeeee ee ence nene 52

E. Electronic Structure.......... 0.0 c cece cece cece nescence eee eeeeeaseeseessreeeeetes 61
references 62

F. Alloy Degradation
1. Disproportionation...............ccecsecce cece eeeeeteneereneeneeeseeceeneneenens 64
9 @°0) 6 c0)-5 (0) | nen nen 68

3. Approaches to Suppress Degradation. ........... 2.02... ccece eee ec essence eeees 69

references 78

Il. Methods: Experimental Techniques, Equipment, and Data Analysis

A. Alloy Preparation......... 2... ccc sce e cece eee ee ee eece eee eeseeeseeeseeeeteasenneeaeees 82
B. Gas-Phase p-c-isotherms............... ccc ecceesceeccecceeeteeseenceneeessneeeeaees 82
C. X-ray Diffraction. ........ 2... ce cece cece eee ce eee eeseeeeeeaeesseneeeneeees 87

references 110

_ D. Microprobe AnalySis.............:.cccccecececnceeneeene nen eneeeeeeeaeeseensenenenes 111
E. Electrochemical Tests................. ccc eceeccescecesereeseeseeneneeneesenceeesens 111
1. Kinetics. .... 0... cece cece ec ee ence cece cence neon nee eeseeeneneeeneneeeeeaeess 112

2. Cyclic Lifetime Tests............. cc cccece cece eee ee eee nena ceeeeeeeeeaseeennees 123

IV. X-Ray Diffraction

A. Phase CompoSition............ 0... cece cee cc ence nee nee eee eeeeeeneeeeeneeeneeae ences 128
B. Lattice Parameters of Dehydrided Alloys.................. ccc eec eee eceeeeeeeeee 137
C. Hydride Volume Expansion.......... 0.0... cccece cece eeeeen eee eeeeneeeeeseeeeeee 147
D. Microstrain of Dehydrided ANOyS.......... 0... ccc cece eee e ee eeeeeeereetnnenaeees 169

V. Isotherms.............ccc cece cece ccc sev esau ccceeccusscunscecscceuesecesseeceusunesssueesess 178

VI. Electrochemical Kinetics............... 0... cccecee eee cenee eee eeeeeeeceneeeeeeeeeeeeens 204
A. DC Micropolarization cece eee e eee e eee e nese eee e eee eneeeeeeaseeaeeassneeeeseenaeees 204
B. AC Impedance............ cece cece see cence eeceee et eeeeeeetee eerste eeeen seers sence 209
C. Tafel Polarization.............. 00. e cece cece eee eee eee eee eeeceeeeeceeseneeneeees 214
| DR Bec 011.) (0) 0 220

VIL. Electrochemical Cycling
A. Cycling Conditions......... 0.0... ccc cece ec ceeee cece eeee cen eeeeeeeeeeseeeseneneets 232

Vill

B. Effects of Electrochemical Cycling... 2... ccc ccc s ce eceee ese e ee eeneeeeneeeeues 244
C. Hydrogen Absorption Capacity...............cccc cece cece ee ene seer ee eesneneees 254
D. Cyclic Lifetime...........ccccccccccccesesscecessseeeecessesseeccsseeecnsseeseeesass 257

VITI. Conchusions........... 0.0... cece cece cen c cece cc ceneccensceuscccccceeccseeevssseusnesceecs 274

IX. Appendix — Table of solute properties............... cc cee ee cece ceee nsec eeeteeeeees 275

Lists of Tables and Illustrations

Chapter 2: Background

Figure II-1
Figure II-2
Table II-1

Figure II-3
Figure I-4
Figure II-5
Figure 1-6
Figure II-7

Figure II-8
Figure II-9

Figure I-10
Figure I-11
Figure [I-12
Figure II-13
Table I-2

Figure I-14
Figure II-15

Gas-phase isotherm of PdH..................cccccceeeeceeeeeneeeeeeetees
Gas-phase isotherms of archetypal AB, alloys.....................08

Volumetric and gravimetric densities of AB, intermetallic

Detail of Ni-rich half of La-Ni Phase diagram.....................006
Unit cell of LaNiy....... cece cece cence eee ce eee neeenceesseeeneeeeuene
Dependence of LaNi, lattice parameters and isotherm
characteristics on alloy stoichiometry................:.eceeeeeeee eee eeee
Room temperature gas-phase isotherm of LaNi,...................066
Schematic of multiple T isotherms and associated

van’T Hoff plot of MH alloy............. ccc cece ec ee cece cece eee enees
In-situ XRD of LaNi,H, during hydrogen desorption................
i-phase lattice expansion of LaNi,..............cc cece ec eee necro eee
Deuterium site-occupancy in LaNi,H,. 2-site and 5-site models...
D ordering for LaNi,H, for y 2 5.1.0.2... 0:ceeeeeeeeeeeereeeeeneeeeenees
Deuterium site occupancy in B-LaNi,H,...............c:ceeeeereeeees
Electrochemical isotherm of LaNi,...............scscseesseee nese eeeaes
XRD pattern of activated LaNi, depicting anisotropic line

broadening...... 0.0... ccc cece eee e cece e ence nsec nee eec ees eneeneeeeeenees

Williamson-Hall plots depicting LaNi, XRD anisotropic line

Figure II-16
broadening cece eee e ence ence eee ee snes eneceeeeneseeeneceeeeeeeeeeneseeeeeeees 41

Figure I-17 ~ Unit cell volume vs x for LaNi,,M, alloys............::.:::ceeees 49
Figure I-18 Metal site occupancy of LaNi, M, alloys................--seeseeeeees 50
Figure [1-19 Lattice expansion in LaNi, (Co, System..............:seeeeeeeeeee teens 54
Table II-3 Lattice expanding behavior of Ni-substituted LaNi, alloys with

hydrogen F10)<10) 9 0] 5 (0) | Oe nnee 56
Table -4 Deuterium site occupancy in B-LaNi,,M,H)..........-..::::eeeeeeee 57
Figure II-20 Schematic of LaNi, disproportionation.................:ecseeeeeeeeees 66
Figure [I-21 S,,, vs % volume expansion upon hydriding.................+sseeeee 70
Figure II-22 Correlation of AB, alloy cyclic lifetime with metallic radius of

SOlUte AtOM........... cece sees cece eee eee eee e ence nets nee neeeneeneeeaeeeeeees 76
Figure [I-23 Correlation of AB, alloy cyclic lifetime with AH,,,...-.-.------ 77
Chapter 3: Methods
Figure TI-1 Picture of Caltech Sievert’s apparatus................eeeeeeeeeeeeeeen ee 83
Figure II-2 Schematic of Caltech Sieverts’ apparatus................seeeeeese eee 84
Figure II-3 Sample reaction chamber.................ccececeeeencecerenseeeaeeeeeraes 86
Figure [1I-4 Schematic of Inel diffractometer..................cscececeeeeeeneeeeeees 93
Figure I1I-5 XRD pattern of LaB,.................... eee enceneeeeeeeseeeeeesenteneenes 95
Figure II-6a Caltech Inel channel spacing non-linearity..................eeeeee eee 97
Figure IJ-6b Published Inel channel spacing non-linearity.................+.0. e+e 98
Figure II-7 20 geometry, fixed 0..............ccccecececeseceeeeeeeeeeneneeeneenenees 99
Figure 1-8 Radiation path in sample for 20 geometry...............s.ceeceeeeeeee 100
Figure 1-9 Absorption correction in 20 geometry for several incident angles.. 102

Figure I-10

Focusing circle of 20 diffractometer...............ccceeeeeeeeeeeeeeeees

Figure I-11 Peak broadening of 20 diffractometer....................eeeeeeee enone 104
Figure I-12 Peak position errors from sample displacement and transparency
~ for 20 diffractometer................cccccceeccecceeeeceuseeceeeeeeseeeees 106

Figure TH-13 CO surface poisoning stand................ccccceceeeeeeeeeeeeeeeeeeeeess 108
Figure III-14 Picture of half-cell......... 0... cece cece cece eee eeeeee en eeeneee eases 113
Figure I-15 Equivalent circuit for charge-transfer reaction at an electrode...... 118
Figure III-16 Schematic of Cole-Cole plot of MH impedance spectrum.......... 120
Figure I-17 Equivalent circuit for MH electrode.................cc ec eceee eens ne ees 121
Figure I-18 Picture of prismatic cell............. cece ccc e cece tena eeeeeeeeeeeeeee 124
Chapter 4: X-Ray Diffraction
Figure IV-1 XRD patterns of hydrogen activated LaNi, Sn, alloys.............. 129
Figure [V-2a XRD patterns of hydrogen activated LaNi, Ge, alloys.............. 130
Figure IV-2b SEM micrograph of annealed LaNi, Ge, , ingot..................00 131
Table IV-1_ EDAX composition analysis of LaNi, Ge, alloys.................06 132
Figure IV-3a XRD patterns of hydrogen activated single-phase LaNi,,M,,

ALLOYS... eee eee cece eee e eee eee ee eee eeeseeeeeeeneeeeneeeeeseceeseneeenes 133
Figure IV-3b XRD patterns of hydrogen activated multi-phase LaNi, .M,

CC) na 134
Figure IV-4 XRD patterns of hydrogen activated LaNi, M, alloys............... 135
FigureIV-5 Unit cell volumes of LaNi, Sn, alloys vs Sn composition, x....... 138
Figure IV-6 Unit cell volumes of LaNi, Ge, alloys vs Ge composition, x...... 141
Figure IV-7 Unit cell volumes of LaNi, Si, alloys vs Si composition, x.......... 142
Figure IV-8 Unit cell volumes of LaNi, In, alloys vs In composition, x......... 143
FigureIV-9 Unit cell volumes of LaNi, Ga, alloys vs Ga composition, x....... 144
Figure IV-10 Unit cell volumes of LaNi, Al, alloys vs Al composition, x........ 145

Figure IV-11
Figure IV-12

Figure IV-13

Figure [V-14
Figure IV-15

Figure IV-16

Figure [V-17

Figure IV-18

Figure [V-19

Figure [V-20

Figure [V-21

Figure [V-22

Figure IV-23

Figure [V-24

Figure IV-25

Figure IV-26
Figure [V-27

Lattice expanding effectiveness of solutes...............cceeeceee cece
Aala [%] for transition metal-substituted LaNi, M..............--+
Ac/c [%] for transition metal-substituted LaNi, M.,..............++

AV/V [%] for transition metal-substituted LaNi, M

Evolution of LaNi,,,In, ,H, XRD pattern as alloy evolves

Evolution of LaNi,,,Al H, XRD pattern as alloy evolves

4.740 “0.26

Evolution of LaNi, ,Ge,,H

02° "y

XRD pattern as alloy evolves

Evolution of LaNi, Ge, ,H, XRD pattern as alloy evolves
NYATOSEN.... 0. cece cece cee eee cence ene eea ene eeeeneene eens eeeee neces

Evolution of LaNi ,sGe,,H, XRD pattern as alloy evolves

Xil

- Figure IV-28

Figure IV-29:

Figure [V-30

Solute’s effectiveness in suppressing total lattice expansion vs
solute heat of formation with La, AH) jy.....ssceccscsecseseeeeseeeen
Williamson-Hall plots of anisotropic XRD line broadening in

LaNi, Mn,. Data points “1” represent hki families of diffraction

Williamson-Hall plots of anisotropic XRD line broadening in

LaN: i, Al. Data points “i” represent hki families of diffraction

Microstrain. dAk vs Ak for activated LaNi, .M,,......seseeeeseeeees
Microstrain of singly activated LaNi,,M,, alloys vs x M heat

of formation With Law... 2... cece cece cee ccc cee nec cecceeeneeentencees

Hydriding parameters for LaNi, Sn, alloys. Isotherms

measured at 100°C... 0.0... ccc ec cece cee eceeee ee sateeeeeeneneereeneees
Hydriding parameters for LaNi, Al, alloys. Isotherms

measured at 40°C... cece cece cee ce nee ee ee ee ee eeeeeeneneneenenenaes
Hydriding parameters for sp-shell metal substituted LaNi,

alloys. Isotherms measured at 30°C..............cceceeeeeee scence ees
Room temperature gas-phase isotherms of LaNi, Ge, alloys,

08 OSD, 0 ie Sane
Elevated temperature isotherms of LaNi,,Ge,,. ............2ssseee0
Elevated temperature isotherms of LaNi, ,Ge,,. ........:.seseeeeeeee
Elevated temperature isotherms of LaNi,,Ge,,. ............0.c0ee00

Elevated temperature isotherms of LaNi, Ge, ,. ..........seeesere ees

Figure IV-31
Figure [V-32
Chapter 5: Isotherms
Table V-1
Table V-2
Table V-3
Figure V-1
Figure V-2
Figure V-3

Figure V-4
Figure V-5
Figure V-6

Elevated temperature isotherms of LaNi, Ge, ,. ...........-.:.ese0e

183
184
185
186
187
188

xiii

xiv

Figure V-7 Van’t Hoff plots of R in(P,,,) vs 1/T for LaNi, Ge, alloys. ........ 189
Table V-4 Room temperature hydriding parameters for LaNi, Ge, alloys. ... 190
Figure V-8 ~— Room temperature isotherms of LaNi,,M, alloys..................66+ 193
Table V-5 Room temperature Hydriding parameters for LaNi, .M, alloys. ..... 194
Figure V-9 —_Gas-phase hydrogen capacities of LaNi, Sn,, LaNi, Ge, and

LaNi, M, alloys............ cee ce eee e cee ce eee een ec enc eeneeteeeeeteeeeen ness 195
Figure V-10 Juxtaposed isotherms of LaNi, Ge, and LaNi, Sn,

alloys, O< X <0.5 0... cece cece eee e een eneeeeeeneeneneeeees 197
Figure V-11 a) Semi-log plot of plateau pressure vs unit cell volume,

b) Enthalpies of hydride formation and decomposition vs

alloy unit cell volume, for LaNi, Sn, and LaNi, ,Ge,.............4+ 198
Figure V- 12 Enthalpy vs volume parameter plotted with a) solute metallic

radius and b) the heat of formation of the solute with La............ 201

Chapter 6: Electrochemical Kinetics

Figure VI-1 Linear polarization of LaNi, Sn, alloys................cccceeeeeeee ees 205
Figure VI-2 Linear polarization of LaNi, Ge, alloys.................ceseseeeeeeees 207
Figure VI-3__ Linear polarization of LaNi, M, alloys................:.:eceseeseeees 208
Figure VI-4 AC impedance spectra of LaNi, Sn, alloys..............:.ceeeee neon 211
Figure VI-5 AC impedance spectra of LaNi, Ge, alloys.................eceeeeeees 212
Figure VI-6 AC impedance spectra of LaNi, .M, alloys...............csceseeee eens 213
Figure VI-7 Tafel polarization of LaNi, Sn, alloys..............cceceeeeseeee eens 215
Figure VI-8 Mass corrected Tafel polarization of LaNi, Sn, alloys.............. 216
Figure VI-9 Tafel polarization of LaNi, Ge, alloys................ccceceseeeee eens 218
Figure VI-10 Tafel polarization of LaNi, M, alloys................csceceeeeeeeeneees 219

Table VI-1 Kinetic parameters of LaNi, Sn, alloys...............:.csseeeeeeeenes 223

Table VI-2
Table VI-3
Figure VI-11
‘we VI-12
Figure 3
Figure VI-i4

Figure VI-15

Figure VI-16

Kinetic parameters of LaNi, Ge, alloys.............ccc cee eeeeeeeee ees 224

Kinetic parameters of LaNi, M, alloys.................sceeeeeeeereees 225

Exchange currents of LaNi, M, alloys measured by DC

MicTOPOlariZatiON.......... 2... ece cece cence cece ee eeeeeeeeeeseneeeeeenenees 226

Exchange currents of LaNi, M, alloys measured by AC

UMPeMaNce...... 2. cece eee cece eee eee cece teeta eeeeeeneseetneeeeeeenees 227

Exchange currents of LaNi, M, alloys measured by anodic Tafel

00) 6b 2:15 (0) | een ee 228
‘ange currents of LaNi, M, alloys measured by cathodic Tafel

Po. 0) | Rene 229

Transft efficients of LaNi,,M_. alloys measured by anodic Tafel

Polarization..............cceecese ec eeee ee eee nee eeeeseeeeeeeensneeeteeeeees 230

Transfer coefficients of LaNi, M, alloys measured by cathodic Tafel

PolariZation............ ccc cece eee e cece ence cn ceccccecceccesseeeeceseeeeesers 231

Chapter 7: Electrochemical Cycling

Figure VII-1

Figure VII-2
Figure VII-3

Figure VII-4

Figure VII-5

Lifetime of LaNi,,Sn,, during electrochemical charge-discharge
CYCHING.... cece cece cece cence ee eee seen renee enseeeaeeneeeeeeenenetess 233
Dynamic electrochemical discharge isotherm of LaNi,,Sn,,, ...... 236
Variation of LaNi,,Sn,, lifetime with cut-off voltage:

-0.7, -0.75, and -0.8V vs Hg/HgO. . ......... ccc ecc eee ee eee eneeeeeees 237
Variation of LaNi,,Sn,, lifetime with length of open-circuit

stand after discharge. ...............ccccseeceeeceeeeeseeceeeeeeeeeneneees 239
Variation of LaNi,,Sn,, lifetime with corrosion potential (-0.3

and -0.1V vs Hg/HgO) during 15 min. rest period after discharge.. 240

Figure VII-6 Lifetimes trend line for LaNi, Sn, alloys at elevated temperatures.

Sn,, @ 10°C, Sn,, @ 25°C, Sn,,, @ 44°C, Sn,, @ 59°C.......... 242
Figure VII-7 ~ Lifetime of LaNi 4y50,, electrode with Teflonized carbon binder

and diluent compared to standard LaNi,,Sn,, electrode............. 243
Figure VII-8 Linear polarization of LaNi,,Sn,, cell at different stages of

CYCLING. 0... e cece eee ee enc ee cence eens scene es eeeeeene ree eneeeenes 247
Figure VH-9 AC impedance spectroscopy of LaNi,,Sn,, cell at different

Stages Of CYCLING........... ccs e cece eect eee e eect eet eeeeten ence eeeeeeeenes 248
Figure VII-10 Cycling conditions and capacity cycle lifetimes of LaNi,,Sn,,

Cells 6d? - 6g? cece cece eee e cen e eee eee eee eeneceneeeseeeeeeeeeeteeees 249
Figure VII-11 XRD patterns of cycled cells ‘d’ - “g’........ cece esec sense ee nee ee ees 250
Figure VII-12 Cycling conditions and capacity cycle lifetimes of LaNi,,Sn,,

(COMS 6A? ~ Ce eeeeeeeeeeeee cece eee e tenner ee eeee esate teteeeeeeen nner tens 251

Figure VII-13 XRD patterns of virgin LaNi, ,Sn,, and cycled cells ‘a’ - ‘c’...... 252
Figure VII-14 Diffraction peak intensities of corrosion products La(OH), and

Ni(OH), in degraded electrodes.............cccceceeeeseneneeeneneenceees 253
Figure VII-15 Electrochemical and gas-phase capacities of LaNi,,M, alloys..... 256
Figure VII-16 Electrochemical cycle lifetimesof LaNi, Sn, alloys................. 260
Figure VII-17 Electrochemical cycle lifetimesof LaNi,,M,, alloys................ 261

Table VII-1 Linear decay constants of LaNi,,M,, alloys during

electrochemical cycling................cccsceseseeeeeeeeeeeeeeesenee sees 263
Figure VII-18 Electrochemical cycle lifetimesof LaNi, Ge, alloys................ 264
Figure VII-19 Electrochemical cycle lifetimesof LaNi, M, alloys...............+. 266
Figure VII-20 Capacity degradation rate parameter 6 VS X ...........ceceeeeeeeeees 269
Figure VII-21 Capacity degradation rate parameter 5 vs AV/V ..........0:0000006 270

Figure VII-22 Capacity degradation rate parameter 5 vs AH, ,, ........cse-:00eee 271

Xvi

I. Introduction

This section will briefly describe the organization of the thesis.

The reader will note that the “background” section (§II) is quite large, 1/3 the length
of the text. In writing my thesis, I felt it necessary to describe the context within which
my work would exist. The short “History” (§[1.A) and “Applications” (§1J.B) sections
allude to the technological and scientific importance of metal hydrides. The great volume
of applications of this class of material is argument enough that continued research into
the scientific nature of hydrogen absorption in metals will be valuable.

The academic setting for my thesis begins to appear in “LaNi,’ (§I.C) and is
expanded upon in “Alloy Modifications” (§I1.D). These sections describe what is
currently known about how LaNi, and its alloys react in the presence of hydrogen. The
specific dilemma that my work attempts to address is presented in “Alloy Degradation”
(§ILF). Included in this section are the attempts of researchers to prevent LaNi, from
deteriorating while it works. Sections § II.C.1, § IL.C.3, § ILC.5, § ILD, and § II.F will be
most useful in interpreting my work.

“§ IV. X-Ray Diffraction” investigates the physical changes that occur when LaNi, is
alloyed by Ni-substitution, and when these alloys absorb hydrogen. “§ VII. Electro-
chemical Cycling” explores the phase change that takes place during electrochemical
cycling and its dependency on alloy properties. These two sections are integral to
reaching the goal of my thesis. ““§ V. Isotherms” and “§ IV. Electrochemical Kinetics”
are useful in gaining further understanding into the hydriding and bulk electrode
properties of Ni-substituted LaNi,, but their understanding is not essential in increasing

the alloys’ lifetimes during cycling.

In completing this work, I concur with Bob Bowman’s statement that we know a
great deal about LaNi, and other metal hydrides, but what we understand is a much
smaller realm. I hope that the work I have completed will advance our understanding of
the fundamentally interesting natural phenomenon of hydrogen absorption in metal

hydrides.

ll. Background

A. Short History of Metal Hydrides

1. Palladium

In 1866, T. Graham reported the first observation of the absorption of hydrogen by a
metal, namely Pd.' He followed this work with experiments measuring the variation of
several physical properties, including lattice expansion, electrical resistivity, and
magnetic susceptibility, with hydrogen composition.” This discovery of the ability of Pd
to absorb large amounts of hydrogen has been followed by many years of research into
the science and application of this relationship between one of the periodic table’s most
interesting elements and this most versatile class of materials.

A pressure-composition-isotherm, commonly called an isotherm, of PdH at room

temperature is shown in Figure II-1. Pd reversibly absorbs hydrogen up to a

composition of 1.0 H/Pd at a pressure that rises gradually with hydrogen content.

2. Intermetallic Hydrides

The first hydride of an intermetallic compound, cubic ZrNiH,, was discovered in 1958
by Libowitz, et al.* This finding was particularly auspicious because the properties of
hydride forming intermetallic compounds can be controlled. Intermetallic alloys which
form hydrides are traditionally generalized by the chemical formula AB,, where A refers
to the strongly hydride forming element, usually a rare earth (REZ) or early transition
metal, and B usually refers to a late transition metal. Other types of intermetallic
hydrides are FeTi (b.c.c. AB type), Mg,Ni (A,B type),° (Zr,Ti)(Mn,Cr,V), (C14 or C15
Laves phase AB,),” and LaNi, (Haucke phase AB,), the latter of which is the subject of

this thesis.

20 L

16 F-

Pressure, atm

12 -

0) 0.2 0.4 0.6 0.8 1.0

Figure II-1 | Gas-phase isotherm of PdH, 8 = H/M. (reprinted
from J.R. Lacher, Proc. Roy. Soc. (London), Ser. A,
161 (1937): 526.)

‘The hydriding properties of the compound SmCo, were first noticed by Zijlstra and
Westendorp.” They were studying a change in the material’s magnetic properties with
hydrogen content, and were surprised to find a large reduction in coercivity with a
correspondingly large increase in hydrogen content, both of which were reversible.
Neumann conducted the first studies of the hydrogen absorption properties of LaNi,,” but
the importance of this material was not recognized until the work of van Vucht, er al.,
who explored the discovery of Zijlstra and Westendorp with a variety of AB, materials.’
As stated by van Vucht, et al. these metal hydrides are “exceptional in that they absorb
hydrogen at room temperature quickly and reversibly, dependent only on hydrogen-gas
pressure.”" The density of hydrogen in LaNi, at full capacity, ~6.7 H atoms per LaNi,
formula unit (f.u.), is 7.6 x 10” atoms/cm’, or ~1000 cc hydrogen gas per cc LaNi, at STP.
This is much greater than that of liquid hydrogen, 784 cc/cc at -273° C, or hydrogen in
high-pressure storage tanks, ~200cc/cc at 200 atm.

Isotherms of some of these archetype materials are shown in Figure I-2." We can
see that these materials operate at a variety of pressures at room temperature, making
them useful for different applications. They also have a range of volumetric and
gravimetric densities as hydrogen storage materials, seen in Table II-1. The aspect of
intermetallic hydrides that makes them particularly useful and attractive for applications
is that their isotherms, and hence operating pressures as well as several other properties of
these materials, can be modified by alloying them with other metals.

. Although intermetallic compound hydrides are distinct from metal hydrides, I will

call all AB, intermetallic alloys metal hydrides (MH).

100

(40°C)

la Mg5Ni
(300°C) |

Dissociation Pressure [atm]

(-90°C)

O01

05 1
Atom Ratio([H/M]

Figure II-2 | Gas-phase isotherm of archetypal AB, alloys.
(reprinted from a. Percheron-Guégan and J.-M.
Welter in Topics in Applied Physics: Hydrogen in
Intermetallic Compounds I, Chapter 2, L.
Schlapbach, ed., Springer Verlag (Berlin: 1988).)

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T. Graham, Phil. Trans. Roy. Soc., 156 (1866): 399.

T. Graham, Proc. Roy. Soc. (London), 17 (1869): 212, 500; C. R. Acad. Sci., 68
(1869): 101, 1511; Ann. Chim. Phys. (Paris), 16 (1869): 188; Ann. Chem. Pharm.,
150 (1869): 353; 152 (1869): 168.

J.R. Lacher, Proc. Roy. Soc. (London), Ser. A, 161 (1937): 526.

G.G. Libowitz, H.F. Hayes, and T.R.P. Gibb, J. Phys. Chem. 62 (1958): 76.

J.J. Reilly and R.H. Wiswall, Inorg. Chem., 13 (1974): 218; see references for earlier
work on FeTi-H.

J.J. Reilly and R.H. Wiswall, Inorg. Chem., 7 (1968): 2254.

A. Pebler and E.A. Gulbransen, Electrochem. Technology, 4 (5-6) (1966): 211.

D. Shaltiel, I. Jacob, and D. Davidov, J. Less-Common Met., 53 (1977): 117.

H. Zijlstra and F.F. Westendorp, Sol. State Comm.,'7 (1969): 857.

H.-H. Neumann: “Léslichkeit von Wasserstoff und Deuterium in LaNi,; Ph.D.
Thesis, Faculty of Chemistry, Technische Hochschule Darmstadt, 1969.

J.H.N. van Vucht, F.A. Kuijpers, and H.C.A.M. Bruning, Philips Res. Rep., 25
(1970): 133.

K.H.J. Buschow and H.H. van Mal, J. Less-Common Met., 29 (1972): 203.

A. Percheron-Guégan and J.-M. Welter in Topics in Applied Physics: Hydrogen in
Intermetallic Compounds I, Chapter 2, L. Schlapbach, ed., Springer Verlag (Berlin:
1988).

B. Applications

Despite the early discovery of metal hydrides, widespread application of this
phenomenon was not pursued until the 1970’s. That decade heralded a small explosion in
metal hydride research that continues today. This renewed interest is exemplified by the
publication of many texts about the science and technology of metal and intermetallic
hydrides, such as the Topics in Applied Physics (TAP) volumes Hydrogen in Metals I &
il | (1978) '” Hydrogen in Intermetallic Compounds I (1988) & II (1992),** and JI
(1997).”

1. Gas-Phase

There are many characteristics of metal hydrides that are of scientific and
technological interest. Some examples include crystal structure, thermodynamics and
phase diagrams, electronic and magnetic properties, the lattice dynamics of hydrogen and
the host metal, surface properties, and elastic interactions. These properties can be
harnessed in unique ways to create heat pumps and refrigerators, hydrogen purifiers and
isotope separators, temperature sensors, steam generation plants, chilled water production
plants,”* hydrogen sorption cooling for air conditioners,” cryocoolers,’ and many, many
other technological applications. J. J. Reilly provided an extensive, although early,
description of possible technological uses for metal hydrides." I have transcribed some

of his descriptions to give the reader an idea of the usefulness of MH heat pumps:

The rather large latent heat effects of metal-hydrogen systems can be exploited to
design novel heat pump cycles. In essence, the device consists of two metal
hydride beds, an evaporator bed and a condenser bed as shown in Fig. 13. Gruen
et al." have designed and built a demonstration facility (HYCSOS) to test various
alloy hydride and bed configurations. Initial designs incorporated a heat storage
capability which was later eliminated because of its high cost. Present designs
specify two different AB, alloy hydride beds having matched pressure-
temperature characteristics to give optimum performance. The beds are cycled
very rapidly to reduce alloy inventory to a minimum. A mechanical compressor
is not required.

‘Most recently Alefeld has proposed a heat pump topping cycle for power
generation using a high temperature Mg alloy hydride and low temperature
ferrotitanium or AB, alloy hydride. The author envisages an increase in power
generation efficiency from 37 - 49%. The cycle is rather complicated and for a
full discussion the reader is referred to the references cited.

The hydriding and dehydriding reactions of a number of unstable metal-hydrogen
systems are rapid enough to consider their use as hydrogen compressors or
pumps. The pumping action is derived from the alternate decomposition and
reformation of a metal hydride using low grade heat and a heat sink to provide the
driving force. The first such pump using an unstable hydride was built in 1971
for laboratory use and is still in operation. The pumping action was obtained by
the alternate decomposition and reformation of VH, using hot (50° C) and cold
(18° C) water. A more sophisticated and higher capacity compressor has recently
been built using LaNi,H, and its operation is described by van Mal.”*

An interesting variant has been suggested by Powell et al.'’ who are concerned
with highly efficient power conversion systems. It is based on using a low
temperature heat source in combination with a high temperature heat source in a
closed Brayton cycle where hot compressed gas is expanded through a turbine.
The novel feature of the system is the use of an unstable hydride and low
temperature heat to effect the compression of the gas, thereby eliminating
mechanical compressor work and substantially increasing the efficiency of the use
of the high temperature heat.

A later overview of MH applications can be found in TAP v67, chapter 5, written by
Gary Sandrock, S. Suda, and L. Schlapbach."” This compilation adds liquid hydrogen,
catalytic, electrochemical, and permanent magnet production to the long list of MH
technologies. The most recent description of MH applications was written by P. Dantzer
for TAP. v73, chapter 7." One especially useful aspect of this review is that it addresses
the properties of MH alloys that are relevant, generally and in particular, to technological
applications.

Of course, one of the most important applications of metal hydrides is simply the
storage of hydrogen gas for use as a fuel, whether for combustion in a turbine engine or
electrochemical oxidation in a fuel cell. The fossil fuel crisis in the early seventies

brought about a search for alternative fuels, and hydrogen was a prime candidate.'* This

prompted in 1974 discussion of the world’s transforming from a fossil fuel to a
“Hydrogen Economy” at the first international conference on Hydrogen Energy, the
Hydrogen Economy Miami Energy Conference.” This seminal conference was soon
followed with the formation in the same year of the International Association for
Hydrogen Energy.”

The properties of significance to hydrogen storage, whether portable or stationary, are
capacity, sorption temperature-pressure relationship, and cyclic lifetime. These are
depicted by the hydrogen isotherm, the van’t Hoff plot, and capacity-lifetime curves. The
high capacity and convenient sorption pressures of LaNi, makes it the most attractive
system for terrestrial applications. At room temperature, the pressure above LaNi, at full
capacity, ~6.7 H/f.u., is ~2.5 atm. The equipment necessary to store and transport
hydrogen in LaNi, is much simpler, more reliable, and less expensive than that necessary
to store hydrogen compressed at very high pressures or liquid hydrogen at very low
temperatures. One drawback to using LaNi, as a hydrogen storage medium, however, is
its low specific density. The 1.5 wt.% hydrogen density attained by LaNi, is reduced to a
practical value of 1.2 wt.% when the container and regulating equipment are added to the

material weight.”

2. Electrochemical

Another aspect of metal hydrides of great interest to society is their ability to act as
the anode material in an alkaline rechargeable battery. The concept of a MH cell is rather
young, and was first attempted by Lewis, et al. in 1967 with PdH as the anode material.”
Not long after Lewis made his PdH cell, IMCHs were investigated as hydride
electrodes.” When LaNi, was first used as the anode material in an alkaline rechargeable
cell, the researchers could attain a capacity of only 100 mAh/g.” This is ~25% of the
material’s theoretical capacity of 372 mAh/g, corresponding to LaNi,H,. Percheron, et

al., were able to achieve higher capacities with substituted LaNi, alloys, obtaining

capacities as high as 325 mAh/g for LaNi,,,Al,,,, 340 mAh/g for LaNi,,,Ti,,,, and 390

475
mAh/g with LaNi,,.Mn,,.° Subsequent work designing AB, alloys for nickel metal-
hydride (Ni-MH) batteries has been summarized in a recent review by Sakai where he
states, “A tremendous R&D effort has been conducted on the MH electrode and the Ni-
MH battery since 1985, leading to prominent progress in every area of performance.” *
The state-of-the-art, commercially produced nickel metal-hydride battery has a practical
energy density of ~175 mAh/g and can make up to 2000 cycles at 100% depth of
discharge (DOD) before the capacity of the anode dips substantially below that of the
cathode.” Over 200 million Ni-MH cells were manufactured in 1994, and “extensive
research on large Ni-MH batteries is in progress for EV application.” ”

The Ni-MH cell commonly has a metal-hydride anode and a pasted Ni(OH),/NiOOH
cathode. Each electrode material is wrapped in nickel mesh, which serves as the current
collector, and enclosed in some separator material, e.g., woven nylon. A 6 molar KOH
solution serves as the electrolyte to transport the reaction species, OH and H,O. A
schematic of the charging and discharging processes is shown in Figure II-3.

There are several aspects of Ni-MH batteries that make them attractive for
commercial use. These properties will all be compared with Nickel Cadmium (Ni-Cd)
batteries, which are considered the current commercial standard as well as being
prevalent in technological circles. First, the cell components are almost identical to those
in Ni-Cd’s, the only difference being the anode, where Cd is replaced by the MH alloy
and the associated conductive and binder materials. Second, the cell potential is slightly
higher than that of Ni-Cd batteries, giving them greater power densities as well as making
them more resistant to the well known “memory-effect.” (The memory effect is caused
by a reversible phase change in the Ni(OH),/NiOOH cathode which is affected by the
charging potential.” ) Although Cd/Cd(OH), anodes have theoretical densities reportedly
as great as 365 mAh/g,” the practical energy density of Ni-MH batteries is still 20-30

wt.% and 50-100 vol.% greater than that of Ni-Cd’s. The graph shown in Figure II-4

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displays the ranges of power density that a variety of rechargeable battery technologies
can achieve vs discharge current density (with the discharge energies normalized by the
anode active material mass). The gas recombination mechanisms available to Ni-MH
batteries make them resistant to the problems associated with overcharge and discharge. *
Finally, there are little or no environmental or human health hazards associated with the
materials used in Ni-MH batteries, in contrast with the well known hazards associated

with the production and processing of the metal cadmium.

Topics in Applied Physics v. 28 Hydrogen in Metals I: Basic Properties, G. Alefeld
and J. Voikl, eds., Springer-Verlag, Berlin. 1978.

Topics in Applied Physics v. 29 Hydrogen in Metals II: Application-Oriented
Properties, G. Alefeld and J. Volki, eds., Springer-Verlag, Berlin. 1978.

Topics in Applied Physics v. 63 Hydrogen in Intermetallic Compounds I: Electronic,
Thermodynamic, and Crystallographic Properties, Preparation, L. Schlapbach, ed.,
Springer-Verlag, Berlin. 1988.

Topics in Applied Physics v. 67 Hydrogen in Intermetallic Compounds I: Surface
and Dynamic Properties, Applications, L. Schlapbach, ed., Springer-Verlag, Berlin.
1992.

Topics in Applied Physics v. 73: Hydrogen in Metals III, H. Wipf, ed., Springer-
Verlag, Berlin, 1997.

° §. Suda, J. Less-Common Met., 104 (1984): 211.

” _M. Ron, J. Less-Common Met., 104 (1984): 259.

* |. Sheft, D.M. Gruen, G.J. Lamich, L.W. Carlson, A-E. Knox, J.M. Nixon, and M.H.
Mendelsohn, in Proc. Int. Symp. on Hydrides for Energy Storage, Geilo, August 14-
19, 1977, Pergamon, New York, 1978, p. 551.

B.D. Freeman, E.L. Ryba, R.C. Bowman, Jr., and J.R. Phillips, Int. J. Hydrogen
Energy, 22 (12) (1997): 1125; R.C. Bowman, Jr., P.B. Karlmann, and S. Bard,
Advances in Cryogenic Engineering, 43 (1998): 1017.

J.J. Reilly, Zeitschrift fiir Physikalische Chemie Neue Folge, 117 (1979): 155.

D.M. Gruen, T.L. McBeth, M. Mendelsohn, J.M. Nixon, F. Schreiner, and I. Sheft,

Proc. 11th Intersociety Energy Conversion Eng. Conf., Stateline, Nevada, American

. Institute of Chemical Engineers, N.Y., 1976.

G. Alefeld in Topics in Applied Physics: Hydrogen in Metals II, Springer, Berlin-
Heidelberg-New York, 1978.

J.J. Reilly, A. Holtz, and R.H. Wiswall, Jr., Rev. Sci. Instrum., 42 (1972) 10: 1495.
H. H. van Mal, Proc. Int. Symp. on Hydrides for Energy Storage, Geilo, Norway,
1977, A.F. Andresen and A.J. Maeland, eds., pp. 251-261, Pergamon Press, London
1978.

J.R. Powell, F.J. Salzano, Wen-Shi Yu, and J.S. Milau, USERDA Report, BNL
50447, Brookhaven National Laboratory, Upton, N.Y., 1975.

Gary Sandrock, S. Suda, and L. Schlapbach, in L. Schlapbach, ed. Topics in Applied
Physics v. 67 Hydrogen in Intermetallic Compounds I: Surface and Dynamic
Properties, Applications, Springer-Verlag, Berlin. 1992, chapter 5.

P. Dantzer in Topics in Applied Physics v. 73: Hydrogen in Metals III, H. Wipf, ed.,
Springer-Verlag, Berlin, 1997, p. 279.

D.P. Gregory, Scientific American, 228 (Jan. 1973): 13.

Hydrogen Energy: proc. the Hydrogen Economy Miami Energy (THEME)
Conference, Miami Beach, FL, March 18-20, 1974, T. Nehat Veziroglu, ed.(New
York : Plenum Press, 1975).

T. Nehat Veziroglu, International Journal of Hydrogen Energy, 20 (1994): 1.

R.L. Cohen, and J.H. Wernick, Science, 214 (1981): 1081.

F.A. Lewis, The Palladium Hydrogen System, (Academic London 1967).

E.W. Justi, HH. Ewe, A.W. Kalberlah, N.M. Saridakis, and M.H. Schaefer, Energy
Conversion, 10 (1970): 183.

H.H. Ewe, E,W. Justi, and K. Stephan, Energy Conversion, 13 (1973): 109.

A. Percheron-Guégan, L. Schlapbach, J.C. Achard, J. Sarradin, and G. Bronoél,
Journées de |’état solide, Société Chimique de France, Paris, Sept. 1974; G. Bronoél,
J. Sarradin, M. Bonnemay, A. Percheron, J.C. Achard, and L. Schlapbach, Jnt. J. of
Hydrogen Energy, 1 (1976): 251; G. Bronoél, J. Sarradin, M. Bonnemay, A.
Percheron, and J.C. Achard, Mater. Res. Bulletin, 13 (1978): 1265.

T. Sakai, M. Matsuoka, and C. Iwakura, in Handbook on the Physics and Chemistry

of Rare Earths, eds. K.A. Gschneidner, Jr. and L. Eyring, 21 (1995): 133.
H. Bode, K. Dehmeit, and J. Witte, Electrochim. Acta., 11 (1966): 1079.

C. LaNi,

1. Phase Diagram and Crystal Structure of LaNi,

The La-Ni phase diagram, shown in Figure II-5, was first determined experimentally
by Vogel.’ He expressed uncertainty about the compositions of several phases, and
several researchers have since reinvestigated the phase diagram,” with the most recent
study done by D. Zhang, et al.’ As seen in the phase diagram, the alloy LaNi, is a
stoichiometric compound below a temperature of ~950° C.

The original X-ray diffraction (XRD) study of the structure of LaNi, and many other
RE AB, materials was performed by Nowotny in 1942.’ It was correctly indexed as
Haucke phase (CaCu, structure),° but later work by Wernick and Geller in 1959 showed
that Nowotny’s determination of the lattice parameters was in error.’ The alloy unit cell
shown in Figure II-6 is hexagonal with lattice parameters a = 5.017 A and c = 3.987 A
The lanthanum atom lies on the basal plane of the unit cell, and the atoms in its first two
nearest neighbor shells (CN = 18) are all Ni atoms, d,,,, = 2.896 A and 3.204 A. The Ni
atom in the 2c location also lies on the basal plane and has 6 Ni atoms in the Ist nearest
neighbors shell, d,,,, = 2.464 A. The second nearest neighbor shell, lying 17% farther
away (d,,,, = 2.883 A), is occupied by 3 La atoms and 3 Ni atoms. The 3g Ni atom
location lies in the z = 1/2 plane and has 8 Ni atoms in the first two nearest neighbor
shells (1% separation, d,, , = 2.464 A and 2.509 A). Its third nearest neighbor shell, lying
30% farther away than the first (d_,,, = 3.2 A) contains 4 La atoms.

Above ~950° C, this hP¢ ordered alloy is stable for a range of stoichiometry LaNi,,,
that increases with temperature to -0.15 < x < 0.40 at 1200° C, returning to LaNi, at the
melting temperature, 1350° C.’ Buschow and van Mal performed an evaluation of the

alloy structure in this range and found that the lattice parameters depend on the

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stoichiometry as shown in Figure II-7, reprinted from ref. 3. The nature of the resulting
disorder was later investigated by Coene, et al.’ The results of their transmission electron
microscopy (TEM) experiments show that for x < 0, there is a simple substitutional
replacement of the 3g Ni atoms with La. When x > 0, La atoms are replaced by 2 Ni
atoms in a dumbbell configuration with the connecting axis along the crystallographic z-

* * 8
direction.

2. Gas-Phase p-c-isotherm

When LaNi, metal is in the presence of hydrogen, the hydrogen is absorbed into the
metal lattice at a hydrogen partial pressure (equivalent to chemical potential) determined
by the temperature and the heat of formation of the metal hydride. This reaction follows

equation II.1:

LaNi, + y/2 H, > LaNi,H,. [1.1]

The relationship between hydrogen composition in the metal hydride and hydrogen
pressure in the gas phase is usually depicted as a hydrogen pressure-composition-
isotherm, or isotherm. A room temperature isotherm of binary LaNi, can be seen in
Figure II-8. There are three sections of note in this diagram. In the first, the hydrogen
pressure and composition are both quite low, and the pressure rises rather steeply with
composition. Here, the hydrogen gas is in equilibrium with hydrogen in solution in the
metal lattice, which is called the a-phase. In the a-phase, the hydrogen pressure is
proportional to the square root of the hydrogen stoichiometry, following what is known
as Sievert’s law.’ As the hydrogen composition increases, attractive hydrogen-hydrogen
interactions begin to dominate, and the hydrogen begins to cluster in the metal lattice.
These regions of high hydrogen composition are called the hydride, or B-phase. In the

portion of the isotherm called the plateau region, the a&- and B-phases are in equilibrium

5.03 , 4.01
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Figure H-7 Dependence of LaNi, lattice parameters and isotherm
characteristics on alloy stoichiometry. (reprinted
from K.H.J. Buschow, and H.H. van Mal, J. Less-
Common Met., 29 (1972): 203.)

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Figure II-8

- with each other and with hydrogen in the gas-phase. When both the a- and B-phase
hydrides are present, the additional variable of the relative phase fractions removes one
degree of freedom from the system, a consequence of Gibbs’ phase rule. In a closed,
isothermal homogeneous system, the variation of the phase fractions is allowed because
the pressure in the 2-phase region becomes constant. The end of the plateau region in the
room temperature isotherm corresponds to a fB-phase metal hydride holding
approximately 5.9 hydrogen atoms per LaNi, formula unit.” Finally, when all of the
metal hydride in the o-phase has been converted to the B-phase, the hydrogen-hydrogen
interactions begin to be repulsive, and the pressure rises. Isotherms of LaNi, have been
measured to a pressure of 1650 atm. Corresponding to a composition of LaNi,H,,..

The effect of increasing temperature is shown schematically in Figure II-9.” The
plateau pressure increases and the width of the plateau decreases, a result of the
narrowing of the miscibility gap between the o- and B-phases. At some critical
temperature for B-phase formation, the miscibility gap and plateau disappear and the a-
phase transforms continuously into the B-phase. An interesting phenomenon that can
sometimes be seen in the LaNi, isotherm is the appearance of a second plateau that
appears between 3 and 3.5 H/f.u.” This has been attributed to the existence of a second
hydride phase. This hydride phase was first noted in the isotherm of the metal
hydride ° and later confirmed by in-situ XRD. Recent isotherm measurements
confirm that the plateau splitting is stabilized at higher temperatures (60 - 90° C) and by

cycling.’®

a. Thermodynamics

The functional relationship between the free energy of hydrogen in the metal lattice
and the absorption or desorption pressures is easily determined by equilibrium

thermodynamics. The hydrogen in the hydride phase is in equilibrium with hydrogen in

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the gas phase, so the free energy of hydrogen in each phase is equal. At room
temperature, the chemical potential of an ideal gas is W = wp, + kT Inp. The
thermodynamic relationships Nu = G and AG = AH - TAS make it possible to find the
enthalpy and entropy of hydride formation by plotting R Inp vs 1/T in a van’t Hoff plot,
shown schematically in Figure II-9.° The heat of formation of the intermetallic hydride
of LaNi, is ~31.6 kJ/mol H,, and the entropy loss associated with this phase change is
~108 J/°C/mol H.,.

Entropy changes in reaction [II.1] are almost entirely dominated by the reduction of
the entropy of gaseous hydrogen from 130.8 J/°C/mol H, "’ to that of essentially metallic
hydrogen. The contribution of H atoms to the lattice vibrational entropy of the solid AB,
hydride is probably quite small near room temperature, the main effect being the addition

of an optical phonon branch at fairly high frequencies.

b. Hysteresis

A pressure hysteresis can be noted between the absorption and desorption plateaus in
the isotherm of LaNi,, as seen in Figure II-8. To some degree, hysteresis accompanies
all first-order phase transitions. In metal hydrides, the transformation from the o-phase
(hydrogen in solution in the MH lattice) to the B-phase MH (hydrogen ordering on the
MH lattice at a composition of ~1 H/metal atom) is a first-order phase transition. A first-
order phase transition is characterized by a discontinuity in the first derivative of the free
energy at the transition point. Evidence of this in the hydride transition can be seen in the
volume discontinuity (explained in §I1.C.3.). Above the critical point for the hydride
transition (see Figure II-9), the volume discontinuity and hysteresis disappear. The

magnitude of the energy dissipated (or entropy produced) in one hysteresis cycle is:

1/2 RT In (p/p,) [11.2]

Reviews of hysteresis in metal hydrides are available by Qian and Northwood * and
by Flanagan and Park,” among others. Ever since hysteresis was first noticed in the Pd-H
system by Lambert and Gates, ” researchers have formulated models to explain the cause
of this effect. The magnitude of the hysteresis has been empirically correlated to the
discontinuity in the volume expansion and the width of the pressure plateau cP - c*,” and
we have found a correlation to the density of crystalline defects as determined by XRD.
Justification has been given for using as the equilibrium hydriding pressure the
desorption pressure, the arithmatic mean of the absorption and desorption pressures, and
the geometric mean of the pressures." Recent experiments to determine the nature of
hysteresis in LaNi, indicate that an equilibrium pressure between the absorption and
desorption plateaus does not exist, rather that the metal hydride exhibits a
thermodynamically important difference during absorption and desorption.” Their
neutron powder diffraction (VPD) measurements of lattice parameter variation during and
after activation indicate that a- and B-phases coexist in LaNi, particles, implying that the
hysteresis is caused by misfit stresses at the o/B interface.” Existing theories adequately
describe some aspects of hysteresis in some MHs, but there is no theory that sufficiently

explains all of the experimentally observed behavior in each MH system.

3. Hydride Structure and Lattice Expansion

When LaNi, absorbs hydrogen, the crystal lattice must expand to accommodate the
hydrogen. At room temperature, the unit cell volume of activated LaNi, first expands
isotropically and continuously ~1.5%. The lattice then experiences a discontinuous, near-
isotropic volume expansion of 21.5% when it transforms from the a@- to the B-phase. In
Figure II-10, showing in-situ XRD by Notten, et al.,” the shift in diffraction peak
positions associated with the lattice expansion can be seen. This figure also explicates
the 2-phase nature of LaNi,. It can be seen that the diffraction peak positions, and hence

the lattice parameters, do not increase continuously when the system absorbs hydrogen,

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but material in the o-phase changes discontinuously to the B-phase.” When all the
available o-LaNi, has been hydrided to the lower limit of the B-phase, the unit cell
volume expands continuously with hydrogen composition. Figure [I-11 depicts lattice
expansion as B-LaNi,H,,, absorbs hydrogen. The a-axis expansion is initially greater than
the c-axis expansion, but subsequent absorption makes the expansion isotropic. At the
largest hydrogen concentration measured, 6.7 H/f.u., the lattice has isotropically
expanded 25.75%.”

Under certain circumstances (elevated temperature, microstrain), the unit cell volume
of a-LaNi, expands 11.8% discontinuously and anisotropically (Aa/a = 5.2%, Ac/c =
1.3%) to the limit of a third phase called the y-phase, and subsequent expansion is
continuous with hydrogen composition. The nature of the volume expansion has been
empirically linked to characteristics of the plateau in the isotherm. A particularly flat
plateau is indicative of the discontinuous volume expansion associated with the
miscibility gap of the first-order phase transition. A sloping nature to regions of the
plateau indicates the alloy-hydrogen system is not in a 2-phase region, and the resulting
volume expansion will be continuous. This can be interpreted as the alloy having higher
temperature and hydrogen composition than the critical temperature and composition for
hydrogen ordering on the Haucke-phase lattice.

The structure of the B-phase hydride is a complex question, and was for some time a
point of contention in the metal-hydride community. Geometric considerations by
Westlake and Switendick determined that hydrogen atoms would occupy interstitial sites
with a minimum radius of 0.4 A ~ and a minimum separation of 2.1 A from another
hydrogen atom.” The best methods to probe hydrogen utilize coherent scattering of
neutrons by deuterium and their large incoherent scattering by hydrogen.” Neutron
diffraction is used to determine the locations of deuterium in LaNi,D, compounds and
quasi-elastic neutron scattering from deuterium or incoherent inelastic scattering from

hydrogen is used to measure the dynamics of hydrogen motion in the alloys. These data

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can also be reinforced by proton NMR measurements which investigate the time scales of
hydrogen motion in the hydride.” An extensive review of H locations in LaNi, and its
substituted alloys LaNi, M_ was performed by Percheron-Guégan, et al.”

The first neutron powder diffraction (NPD) pattern of LaNi,D,, was recorded and
analyzed by Bowman, ef al.” The pattern was indexed to the P31m space group, and
hydrogen atoms were determined to fully occupy 3g (z = 0.1) sites and partially occupy
6d (z = 0.5) sites, a model later confirmed by other groups.” When Percheron-Guégan,
et al. performed Rietveld refinement on a high-resolution neutron diffraction patterns
taken on the D1B instrument at the Institut Max von Laue-Paul Langevin, Grenoble, they
saw no evidence of a symmetry reduction and concluded that the deuteride structure was
in the P6/mmm symmetry group.” In this model, hydrogen partially occupies 5 sites:
3f (1/2 O 0), 12n (1/2-6, 0 8),4h (1/3 2/3 1/2-8),6m (xk 2x = 1/2), and 120
(x 2x 1/2-6,), where 6,.is a small quantity. Figure I-12 shows the LaNi, crystal
structure with hydrogen locations of both models. The models are similar, but while the
2-site model utilizes anisotropic thermal parameters to describe the regions of deuterium
occupation, the 5-site model creates/describes 2 “cages” containing multiple deuterium
sites. In this way, the 5-site model gives more information on the structure of the
potential wells nearby the 3c (corresponding to the 3f, 12n cage) and 6d clusters
(corresponding to the 6m, 4h, 120 cage).

Further study by Percheron, et al. on NPD patterns of LaNi,D,, 5.0 < y < 6.7, noted a
small diffraction peak at the d = 1.805A reflection position in their experiments as well as

in previously published diffraction patterns,”

the intensity of which seemed to
decrease with decreasing D composition. Including this and two other {(201) and (205)}
superlattice reflections in the refinement of their diffraction patterns required a doubling
of the LaNi, unit cell along the c-axis, reducing the crystal symmetry to P6,mc. As seen

in table II-2, it was determined that in this range of D composition, the ratio of deuterium

in the 3c cluster of sites to that in the 6d cluster stays approximately constant. However,

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ordering does occur within the 6d cluster as the D composition increases. Repulsive
interactions between D atoms in adjacent cells forces them out of high symmetry 6m
positions to 4h and 120 positions in order to maximize the H-H separation. The ordering
parameter corresponds to n’ = (1-a)n, n = (x, 2x, z) occupation, n’ = (x, 2x, z+1/4)
occupation. Figure II-13 from ref. 10 shows a schematic representation of the ordering
of the D atoms on (x, 2x, z) positions. In this new model, occupation of the 3f site has
been abandoned. This study was soon followed by an investigation by P. Thompson, et
al., who utilized the different NPD scattering lengths of Ni and Ni to discover a similar
deuterium ordering and symmetry reduction consistent with the 2-site model.”

Statistical thermodynamic calculations have also been used to support both the 2- and
5-site models.” The entropy loss upon hydride phase formation is dominated by the
entropy of gaseous hydrogen, 130.8 J/°C/mol H. Early calculations of the configurational
part of this entropy loss neglected to consider the constraints caused by the interaction

between protons in B-LaNi,.“””

Percheron, et al. calculated entropy and enthalpy changes
for hydrogen absorption using both the 2- and 5-site models in which the geometric
considerations of Westlake are used to provide constraints to a statistical mechanical
partition function for occupation of a multiple interstitial site structural model. Their
conclusions showed support for the 5-site model: “The calculated configurational
entropies for both structures are very close: 3.6 cal (mol H,)" and 2.6 cal (mol H,)' for
the P6/mmm and the P31m space groups respectively.” However, the calculation of the
enthalpy change AH, (-11 kcal (mol H,)" and -8.6 kcal (mol H,)" for P6/mmm and P31m
respectively) shows that hydrogen absorption is more favourable in the structure with
space group P6/mmm.”*

Other support for structural determinations of LaNi,H, comes from hydrogen
diffusion and lattice dynamics measurements. Quasielastic neutron scattering “ and

nuclear magnetic resonance “ experiments have determined two time scales for hydrogen

motion in LaNi,H,. Hydrogen diffusion with low activation energy and diffusion

| coefficients measured to be on the order of 10” cm’/s have been assigned to short range
- hydrogen motion within the cages and between the z = 1/2 cages forming a hexagonal
ring. Long-range hopping between the two sets of interstices has higher activation
energies and corresponds to a diffusion coefficient of ~10° cm’/s.” Although these
measurements do not distinguish between the 2-site or 5-site models, they do give general
information about the hydrogen occupation in the hydride phase.

As mentioned above, the lattice parameters of the intermediate y-hydride phase have
been studied by XRD *“" and NPD “. The most recent work by Gray, et al., shows that
the lattice expansion between the o- and y-hydride phases is discontinuous and
anisotropic (Aa/a > Ac/c) at all temperatures studied. The y- to B-hydride phase
transition, however, is discontinuous at 30° C and continuous at 65° C and 100° C.
“There appears to be a thermodynamic critical temperature for the B + y system in our
sample in desorption between 30 and 65° C.” The existence of the y-hydride phase and
information about its structure imply that the stability of one of the hydrogen
environments in the LaNi, cell is particularly sensitive to temperature and structure

16
changes.

4. Basic Electrochemical Properties

LaNi, is particularly attractive for use in alkaline rechargeable cells because it is
simple to design a LaNi,-based alloy with a plateau pressure below 1 atm., which can
then be easily used in sealed cells. The electrochemical potential of the metal hydride cell

at room temperature is related to the hydrogen partial pressure by the Nernst equation:

E.,,(F(ws. HgO/Hg) = [E°(H) ~ £° (HgO/Hg)] + <= nf 9D)

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yet RT
= [E°(H) E’ (Hg0/Hg)] +> na ry

= -0.9324 - 0.0291 log (P,,,)." [11.3]

E°(H) = standard electrode potential of the H,O/H couple.

E°(HgO/Hg) = standard electrode potential of the HgO/Hg couple.

a(i) = activity of species i. Y(H,) = fugacity coefficient of water.
R= gas constant. T = cell temperature [K].
F =Faraday constant.

Because a change of one order of magnitude in hydrogen pressure only corresponds to
a potential change of 29 mV, Ni-MH batteries have a very stable voltage profile over
their capacity range. In electrochemical isotherms, the cell potential is commonly
converted to pressure by the Nernst equation, and hydrogen composition is displayed as
an electrochemical capacity in mAh/g. As in gas-phase isotherms, electrochemical
isotherms are quasi-equilibrium, usually pulse charged with a waiting period of 20-60
min to allow the potential to reach its rest value before measurement. Electrochemical
isotherms of a Ni-MH rechargeable battery using AB, as the anode material are shown in
Figure II-14.

A hydrogen composition of 6.7 H/LaNi, would give the anode material a maximum
capacity of 372 mAh/g. Because of the weight of the cathode, electrolyte, casing, and
other inactive components, the largest energy density that can be realized in a battery is
approximately 175 mAh/g.” This, combined with a cell potential in the range of 1.2 -
1.5V gives the Ni-MH battery a fairly high power density when compared to other types
of rechargeable battery technologies.” The voltage and specific power of Ni-MH
rechargeables makes them attractive for use in many applications. The similarity in
voltage and benefits in specific energy and environmental compatibility makes Ni-MH
the number one choice for replacement of Ni-Cd batteries for single cell (AAA - D sizes)
and battery pack (camcorder, power tools) applications. For low current applications,

such as laptop computers and portable phones, Li-ion rechargeables will be optimal once

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applications such as electric vehicles.

5. Microstrains

It is well known that the discontinuous lattice expansion accompanying hydrogen
absorption in metal hydrides creates great stress at the interface between the a- and B-
phases because of the misfit in the lattice parameters of the 2 phases. This misfit strain
provides the activation energy to nucleate dislocations and induce planar defects in the
Haucke phase crystalline lattice.” An obvious macroscopic result of the misfit strains is
the decrepitation of ingots of LaNi, to fine powders.

The microstructural effects of the misfit strains are clearly evident in the XRD pattern
of a LaNi, sample which has been activated by gas-phase hydrogen absorption, as seen in
Figure II-15. The diffraction peaks of activated LaNi, are broadened anisotropically,
with large lorentzian (crystallite size) and gaussian (microstrain) components to the
broadening in all directions and additional gaussian broadening in directions orthogonal
to (001). This broadening was first interpreted using the simple method of Williamson
and Hall.” For each family / of reflections (hkl), the broadening in K-space (integral
breadth, Bcos@) was plotted vs the K-vector of the diffraction peak (2sin@/A). An
example of this anisotropy can be seen in Figure II-16, showing Williamson-Hall plots
of peak broadening in an activated sample of LaNi,. The apparent crystallite size in each
family/direction is given by A/B, and the microstrain Ad/d in that direction is given by the
slope of this plot.” (For further explanation of microstrain and the Williamson-Hall
technique, see HI.C.b.) Subsequent works implemented modifications of the Rietveld
method to refine the severe anisotropic strain seen in diffraction patterns of hydrogen
activated LaNi, alloys.”

There have been a number of TEM studies of activated LaNi, alloys undertaken to

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determine the microstructural defects responsible for the diffraction broadening. Kisi, et
al.* attempted a TEM study on bulk LaNi, powder particles, but the severe distortion
defects and lack of sample preparation prevented them from identifying the types of
defects present in activated LaNi,. G.-H. Kim, et al.” used TEM on LaNi, powders
embedded in epoxy and ion beam milled to study the change in the surface of LaNi,
induced by hydrogen activation. They found regions of LaNi, crystals containing
microtwins near the surface of the grains and the decomposition products. Notten, et al.”
studied similarly prepared samples of unhydrided and dehydrided LaNi, and found sets of
planar defects in the dehydrided powders. The displacement vectors of the defects were
determined to be in the (001) plane, which is consistent with the anisotropic diffraction
broadening seen in powder diffraction patterns. In addition, the average spacing between
the defects was consistent with mean particle sizes obtained by Williamson-Hall analysis
of NPD,” rather than the average grain size obtained by SEM. High-resolution TEM
(HRTEM) images showed further that the habit planes of planar defects were actually
stepped (100) plane interfaces having an average inclination of 10° to the (100) plane.
The 180° C cure of the TEM specimens might have modified the defect structures of the
activated powders. However, similar treatments do not affect the severe diffraction peak .
broadening mentioned above, so correlations with these measurements should be
consistent.

Further work by G.-H. Kim, et al.” and H. Inui, et al.” have identified dislocation
structures in hydrogen activated LaNi,. In LaNi, hydrided to the limit of the ot-phase,
they detected “ripple” dislocations thought to be produced by the lattice distortion
during hydrogen penetration into the matrix. In addition, they found <00I> dislocation
loops that could have been produced by the lattice misfit accompanying hydride
precipitation. After complete hydride formation, misfit super dislocations were
found to pile up, sometimes forming hexagonal grid twist boundaries. These types of

dislocations are consistent with the maximum strain to the plane and minimum

‘strain to <00I> plane measured by powder diffraction techniques. They also found some

structural disorder and micro-twins.” Inui, et al. found dislocations formed at

cracks and equal fractions of and <00I> grown in dislocations. Observation of the

facile motion of only the c-type dislocations led them to conclude that the a-type

dislocations contribute to decrepitation in LaNi,.” These determinations of dislocation

direction and burgers vectors are consistent with the anisotropic broadening seen in

powder diffraction patterns,

R. Vogel, Z. Metallkd., 38 (1947): 97.

A. V. Virkar and A. Raman, J. Less-Common Met., 18 (1969): 59; K.A. Gschneidner,
Jr., Rare Earth Alloys, D.D. van Nostrand, New York, 1961, p.221.

K.H.J. Buschow and H.H. van Mal, J. Less-Common Met., 29 (1972): 203.

D. Zhang, J. Tang, and K.A. Gschneidner, Jr., J. Less-Common Met., 169 (1991): 45.
H. Nowotny, Zeitschrift Fiir Metallkunde, 34 (1947): 247.

J.L.C. Daams, P. Villars, and J.H.N. van Vucht, Atlas of crystal structure types for
intermetallic phases (Newbury, OH : ASM International, 1991).

J.H. Wernick and S. Geller, Acta Crystallogr., 12 (1959): 662.

W. Coene, P.H.L. Notten, F. Hakkens, R.E.F. Einerhand, and J.L.C. Daams, Phil.
Mag. A, 65 (6) (1992): 1485.

R. Speiser in Metal Hydrides, W.M. Mueller, J.P. Blackledge, and G.G. Libowitz,
eds. (Academic Press, New York: 1968), p. 72.

C. Lartigue, A. Percheron-Guégan, J.C. Achard, and J.L. Soubeyroux, J. Less-
Common Met., 113 (1985): 127.

J.F. Lakner, F.S. Uribe, and S.A. Steward, J. Less-Common Met., 74 (1980): 13.

L. Schlapbach in Topics in Applied Physics: Hydrogen in Intermetallic Compounds I,
Chapter 1, L. Schlapbach, ed., Springer Verlag (Berlin: 1988).

P.D. Goodell, J. Less-Common Met., 99 (1984): 1.

S. Ono, K. Nomura, E. Akiba, and H. Uruno, J. Less-Common Met., 113 (1985): 113.
T. Matsumoto and A. Matsushita, J. Less-Common Met., 123 (1986): 135.

S. Luo, LD. Clewley, T.B. Flanagan, R.C. Bowman Jr., and J.S. Cantrell, J. Alloys
Comp., 253-254 (1997): 226.

R. Griessen and T. Riesterer, chapter 6 in Topics in Applied Physics vy. 63 Hydrogen
in Intermetallic Compounds I: Electronic, Thermodynamic, and Crystallographic
Properties, Preparation, L. Schlapbach, ed. Springer-Verlag, Berlin. 1988, p. 219.

S. Qian and D.O. Northwood, Int. J. Hydrogen Energy, 13(1) (1988): 25.

T. Flanagan and C.-N. Park, Mat. Sci. Forum, 341 (1988): 297; T.B. Flanagan, C.-N.
Park, and W.A. Oates, Prog. Solid St. Chem.., 23(4) (1995): 291.

B. Lambert and S.F. Gates, Proc. Roy. Soc., Ser. A, 108 (1925): 456.

E.MacA. Gray, E.H. Kisi, and R.I. Smith, J. Alloys Comp., in press.

M.P. Pitt, E.MacA. Gray, E.H. Kisi, and B.A. Hunter, J. Alloys Comp., in press.
P.H.L. Notten, J.L.C. Daams, A.E.M. De Veirman, and A.A. Staals, J. Alloys Comp.,
209 (1994): 85.

D.G. Westlake, J. Less-Common Met., 90 (1983): 251; 91 (1983) 1.

A.C. Switendick, Theoretical Study of Hydrogen in Metals; Current Status and
Further Prospects Report, Sandia Lab, SAND 78-0250 (1978).

A. Percheron-Guégan and C. Lartigue, Materials Science Forum, 31 (1988): 125.

D. Richter, R. Hempelmann, and R.C. Bowman, Jr. in Topics in Applied Physics:
Hydrogen in Intermetallic Compounds II, Chapter 3, L. Schlapbach, ed., Springer
Verlag (Berlin: 1988), p. 97.

A.L. Bowman, J.L. Anderson, and N.G. Nereson, in Proc. 10th Rare Earth Research
Conf., Carefree, AZ, 1973, C.J. Kevane and T. Moeller (eds.), p. 485.

P. Fischer, A. Furrer, G. Busch, and L. Schlapbach, Helv. Phys. Acta, 50 (1977): 421.
A.F. Andresen, in Proc. Int. Symp. on Hydrides for Energy Storage, Geilo, Norway,
1977, A.F. Andresen and A.J. Maeland, eds., Pergamon Press, London 1978, p. 61.

D. Noréus, L.G. Olsson, and P.E. Werner, J. Phys. F: Met. Phys., 13 (1983): 715.

A. Percheron-Guégan, C. Lartigue, J-C. Achard, P. Germi, and F. Tasset, J. Less-
Common Met., 74 (1980): 1.

P. Thompson, J.J. Reilly, L.M. Corliss, J.M. Hastings, and R. Hempelmann, J. Phys.
F: Met. Phys., 16 (1986): 675.

D.M. Gruen and M. Mendelsohn, J. Less-Common Met., 55 (1977): 149.

W.E. Wallace, HE. Flotow, and D. Ohlendorf, J. Less-Common Met., 79 (1981):
157.

J.C. Achard, C. Lartigue, A. Percheron-Guégan, J.C. Mathieu, A. Pasturel, and F.
Tasset, J. Less-Common Met., 79 (1981): 161.

A. Pasturel, B. Brion, P. Hicter, A. Percheron-Guégan, and J.C. Achard, J. Less-
Common Met., 86 (1982): 19.

A. Pasturel, B. Brion, P. Hicter, A. Percheron-Guégan, and J.C. Achard, J. Less-
Common Met., 88 (1982): 442.

D. Ohlendorf and H.E. Flotow, J. Chem Phys., 73 (1980): 2937.

C. Lartigue, A. Percheron-Guégan, J.C. Achard, M. Bee, and A.J. Dianoux, J. Less-
Common Met., 101 (1984): 403.

R.C. Bowman, D.M. Gruen, and M.H. Mendelsohn, Solid State Commun., 32 (1979):
501.

F.E. Spada, H. Oesterreicher, R.C. Bowman, Jr., and M.P. Guse, Phys. Rev. B, 30
(1984): 4909. |

C. Iwakura, T. Asaoka, H. Yoneyama, T. Sakai, H. Ishikawa, and K. Oguro, Nippon
Kagaku Kaishi, 8 (1988): 1482.

T. Sakai, M. Matsuoka, and C. Iwakura, in Handbook on the Physics and Chemistry
of Rare Earths, eds. K.A. Gschneidner, Jr. and L. Eyring, 21 (1995): 133.

G.K. Williamson and W.H. Hall, Acta Metail., 1 (1953): 22.

350

C. Lartigue, ‘A. Le Bail, and A. Percheron-Guégan, J. Less-Common Met., 129
(1987): 65.

P. Thompson, J.J. Reilly, and J.M. Hastings, J. Less-Common Met., 129 (1987): 105.
E.H. Kisi, C.E. Buckley, and E.Mac A. Gray, J. Alloys Comp., 185 (1992): 369.
G.-H. Kim, C.-H. Chun, S.-G. Lee, and J.-Y. Lee, Scripta Met. et Mat., 29 (1993):
485.

A.E.M. De Veirman, A.A. Staals, and P.H.L. Notten, Phil. Mag. A., 70 (1994): 837.
E. Wu, E.H. Kisi, and E.MacA. Gray, J. Appl. Cryst., 31 (1998): 363.

G.-H. Kim, C.-H. Chun, S.-G. Lee, and J.-Y. Lee, Acta Met. Mat., 42 (1994): 3157.
G.-H. Kim, S.-G. Lee, K.-Y. Lee, C.-H. Chun, and J.-Y. Lee, Acta Met. Mat., 43
(1995): 2233.

H. Inui, T. Yamamoto, Z. Di, and M. Yamaguchi, J. Alloys Comp., 269 (1998): 294.

D. Alloy Modifications

One aspect of LaNi, that makes it very attractive for reversible gas-phase and
electrochemical hydrogen storage is that its hydriding and cycling stability are easily
altered by alloying substitutions for both La and Ni. Even in the first examinations of the
reversible hydriding of LaNi, by van Vucht, et al., attempts were made to change its
sorption properties by partial substitution of La with other metals.’ They studied the
variations of the plateau pressures and lattice parameters of La,,Ce.Ni, with Ce
composition. The desorption rate of LaNi, was also studied at various temperatures and
was compared to that of La,,Zr, ,Ni,.. While many investigations have been performed on
the alloying of LaNi,, systematic variations of alloy composition were not routinely
performed, and it is still difficult to extract trends from published data that can be used in
designing LaNi, alloys for use in alkaline rechargeable batteries. The alloys used in my
research were prepared by systematically varying the chemical composition of the solute

element. In all cases, third element substitutions were made for Ni.

1. Phase Diagram and Metal-Atom Crystal Structure

Many metal elements have a range of solubility in which they can be substituted for
Ni in equilibrium LaNi, alloys, and some (Co, Cu) form solid solutions with Ni in the
CaCu, lattice. A good review of the preparation of substituted LaNi, alloys was written
by A. Percheron-Guégan and J.-M. Welter.” Phase composition and crystallographic
states of order have proven to be sensitive to preparation and subsequent heat-treatment
conditions. This fact complicates comparison of published results because alloys of
nominally identical composition but different heat treatment (resulting, e.g., from batch
size) can have different physical properties. Researchers have empirically determined

solubility limits for some elements by the presence of other phases in a CaCu, matrix. No

one has examined the nature of the solubility limits to determine guidelines such as a
Hume-Rothery rule.’

The most immediate effect of substitution on the crystal structure of the alloy is a
change in the lattice parameters and hence the unit cell volume. This is fairly easy to
measure with diffraction methods, and it has been shown that lattice expansion (or
contraction) is in most cases linear with solute composition x. Simple geometrical
considerations outlined by J. Shinar, et al. show that in LaNi,, the z = 0 lattice plane is
close packed while the z = 1/2 plane is not.’ As a result, elements with metallic radii
larger than Ni should have a steric preference to occupy the 3g Ni sites on the z = 1/2
plane.

The hydriding properties and crystal structures of the systems LaNi,,Co, and
LaNi, Cu, have been studied by van Mal, et al.’ and J. Shinar, et al.,’ respectively . Each
system is isostructural with CaCu, for the entire range of substitution. The lattice
parameters of LaNi, Cu, follow Vegard’s law, but those of LaNi,,Co, do not. The
discontinuity in the LaNi, Co, lattice parameters with x has been attributed to magnetic
effects ** and a preferential ordering of Co.’ A variety of researchers have confirmed that
Co atoms do indeed favor the 3g sites, e.g., 88%/12% ’ or 75%/25% * occupation of the
3g/2c sites, rather than the statistical 60%/40%. Cu, however, has been found to have a
slight preference for the 2c site.”

Refinement of the crystal structures of LaNi, .M, alloys has been performed by NPD
and XRD, and by Méssbauer spectroscopy in the case of Fe. Metal atom site occupation
has been determined for Mn,'"” Fe,” Co,” Cu,’ Al,””® Si,’ and Sn. Mn was found to
substitute on both the 2c and 3g sites with a preference for the 3g sites,’ while Fe
occupies almost exclusively (95%) the 3g site.” Noble, or p-shell, metals show strong
preference for the 3g sites, occupying them almost exclusively.” Figures I-17 and
II-18, reprinted from ref. 14, show the unit cell volume and metal site occupancy of some

substituted LaNi, M, compounds.

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et al., J. Less-Common Met., 109 (1985): 287.)

2. Isotherm —

Ternary substitutions for Ni affect the isotherm in four ways: (i) change in plateau
pressure, (ii) reduction of total H capacity, (iii) change in pressure hysteresis, and (iv)
change in shape (reduction of plateau width or increase in slope). These effects are
important in gas phase applications such as heat pumps and cryocoolers, and less so in
rechargeable batteries. For most gas-phase hydride applications, for example, a near-zero
plateau slope is desirable and the pressure-temperature relationship should be precisely
controlled throughout the MH lifetime.” In rechargeable batteries, the plateau pressure
affects the cell potential, self-discharge characteristics, and the ease of testing in the
slightly pressurized configuration often used in the laboratory environment. Of course,
alloys with plateau pressures greater than a few atmospheres can not easily be made into
commercial sealed cells. The total capacity affects the alloy charge capacity, usually
measured in mAh/g. The pressure hysteresis seems to be connected to the alloy’s volume
expansion and cyclic stability. A sloping nature in the isotherm characteristics will affect
the cell potential, but because 1 order of magnitude change in the plateau pressure
corresponds to a potential difference of only 29 mV, Ni-MH cell potentials are rather
stable. The plateau width does seem to be related to the size of the miscibility gap and
corresponding discrete volume change between the a-and B-phases, and this has an effect
on the alloy cyclic stability. Isotherms of LaNi,,Co,H, are unique in the splitting of the
pressure plateau at y = 3.5 for x = 0.5.”

Many models have been proposed to predict the relative stability of MH compounds,
or more precisely the hydride phase heat of formation.’® A particularly useful trend that is
used to design AB, metal hydrides is the linear dependence of the hydride heat of
formation,’”” and thereby the logarithm of the alloy plateau pressure,” on the size of the

18,19

interstices occupied by hydrogen in the cell *”’ and the alloy’s unit cell volume.” There

are many examples of how this has been exploited to prepare complex alloy hydrides

having predictable and desirable temperature-pressure characteristics for specific
applications."” Other effects of substitution on the isotherm are discussed in the

following sections.

3. Hydride Structure and Lattice Expansion

The decrease in hydrogen storage capacity mentioned above (§II.D.2) is a result of the
blocking of hydrogen occupation in specific interstitial sites. Most of the work on
hydrogen (deuterium) site occupation in ternary substituted alloys has been performed by
Percheron-Guégan, et al., and is summarized in Table II-3.” The D-site occupation of
LaNi,,M,H, alloys exhibit characteristics similar to that of LaNi,H,..,.. The ratio of
occupied sites in the 3c cluster to those occupied in the 6d cluster stays approximately
constant. However, deuterides of Ni-substituted LaNi, do not exhibit the deuterium
ordering observed in LaNi,D,,,. The 4h, 120, and 3f sites are found to have progressively
lower occupations, dependant upon substituent and substituted composition, while the 6m
and 12n sites do not experience much decrease in hydrogen occupation. Again, what is
happening is that D is prevented from occupying high symmetry positions in the 6d
clusters. Therefore, the D-D spacing is not allowed to be maximized between adjacent
unit cells. The 6m and 12n interstitial sites are larger than the other sites, and would
therefore be less likely to become too small to be occupied by protons when Ni is
substituted by a larger element. Hydrogen occupation has also been postulated to be
affected by the electron state of the substituting metal. The 120 and 4h interstitial sites
are in close contact with at least two metal atoms in the 3g sites, where Ni substitution
usually takes place. These interstitial sites would be more strongly affected by the
number of valence electrons of the substitutional atoms as suggested by Gschneidner, et
al.” This exclusion of hydrogen sites also precludes the hydrogen ordering found in
LaNi,H,, y 2 5. Other than that, the hydrides of Ni-substituted LaNi, alloys have been

found to retain the P6/mmm symmetry for all substituents but Co,’ as explained below.”

Another effect of Ni substitution is a change in the lattice expansion of the alloy upon
hydrogen absorption. In all cases, substitution reduces the total volume expansion of the
alloy. This is consistent with a reduction in hydrogen capacity, except in the case of Mn
where there is little or no capacity reduction for x < 0.7." The molar volume of hydrogen
in the hydride phase is also often reduced. While the lattice expansion of LaNi, is
approximately isotropic, the substitution of Ni with other metals often induces anisotropy
in the lattice expansion in which the expansion in the direction of the a-axis is greater
than in the direction of the c-axis. This anisotropy is consistent, though more
pronounced, with the anisotropy in lattice expansion seen in LaNi,H,. In LaNi,,Si,H,,,
for example, the a-axis expands 6.4% while the c-axis expands only 2.1%.’ As
mentioned above, the reduction in volume expansion is accompanied by a reduction in
pressure hysteresis.

The lattice expansion and crystal structures of LaNi,,Co,H, alloys were studied by
XRD for 0 < x <5 by van Mal, et al.’ LaNi,CoD, was studied by NPD by Gurewitz,
et al.’ and Latroche, et al. *” at a variety of D compositions. The diffraction patterns of
Gurewitz, et al. exhibited the presence of both the o- and B-phase hydrides.’ The hydride
structures were found to be hexagonal for 0 < x < 3 and orthorhombic for 3 < x < 5, with
LaNi,Co,H, changing from orthorhombic to hexagonal for y = 3.2.° An additional phase
intermediate between the a- and B-phases appears that is coincident with the second
plateau in the isotherm. The lattice expansion caused by hydrogen absorption in the
LaNi, Co,H, system is depicted in Figure II-19. In this figure, open symbols denote
discontinuous volume expansion and closed symbols denote lattice expansion continuous
from the previous discontinuous expansion value. When an intermediate hydride phase
does appear, its lattice expansion is anisotropic, Aa/a > Ac/c. As the hydrogen
composition increases, the anisotropy becomes less pronounced, but it does not

disappear. NPD measurements have determined that in LaNi,CoD,,, which is where the

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“intermediate phase appears, D atoms are spread approximately equally between 6m and
12n sites, with minor occupation of the 4h sites.” Further deuterium loading does not
change the relative occupation of these sites, but does induce the ordering mentioned
above (§II.C.3), displacing deuterium from 6m to 120 sites.”

NPD and XRD measurements have rarely been performed on Ni-substituted LaNi,
alloys at intermediate stages of hydrogen (deuterium) occupation, but there are enough
measurements available to get a good idea of how the lattice behaves when it absorbs
hydrogen. The results of these studies, shown in Table II-4, will be discussed below.
All lattice expansion values are with respect to the dehydrided lattice parameters, unless
otherwise stated. Latroche, et al. performed in-situ NPD on an electrode of LaNi,,Al,,
during charge-discharge cycling. The initial lattice expansion was found to be
anisotropic (Aa/a:Ac/c::5:1) as was the discontinuous lattice expansion (Aa/a:Ac/c::8:5).
Subsequent continuous expansion was greater in the c-direction than in the a-direction.
The volume expansion is reduced appreciably from that of LaNi,, but we do not see
discontinuous expansion being replaced by continuous expansion.”

Latroche, et al. also measured lattice expansion in LaNi,CoD, and multi-component

LaNi Mn,,,Al,,C0,,sD,. These measurements were not performed in-situ, but were

3.55
taken at several deuterium compositions. As seen before, the initial and “discontinuous”
lattice expansions are anisotropic, and subsequent expansion reduces the anisotropy. The
volume of the o-phase at maximum hydrogen composition was not measured, but
assuming ~1% expansion in the o-phase and <1% expansion in the B-phase before the
y = 4.4 measurement, this corresponds to a discontinuous volume expansion of ~14%,
which is consistent with measurements by Gurewitz, et al.’ It is not clear whether
subsequent expansion to 22.4% is continuous or discontinuous. The flat plateau of the
isotherm implies that the volume expansion is discontinuous. However, the in-situ

measurements by Ono, et al.” and Gray, et al.” show that LaNi,H, can have a flat plateau

for y > 4.0 and still expand continuously in that range, implying that LaNi,CoH, will

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expand continuously between y = 4.4 and y = 6.11. This being the case, we do see
discontinuous lattice expansion being replaced by continuous expansion for these two
alloys. In LaNi,Co, the discontinuous volume expansion is large, but is only 63% of the
total volume expansion. The discontinuous lattice expansion is reduced in
LaNi,,,Mn,,Al, ,Co,,, and because total volume expansion is large, it accounts for only
47% of the total volume expansion.

Nakamura, et al. performed in-situ XRD on La,,Sm,,Ni,,Fe,, and LaNi,,Fe,,. They
measured a slightly anisotropic (Aa/a = 6.9%, Ac/c = 5.8%), discontinuous volume
expansion of 20.9% in LaNi,,Fe,, followed by a continuous expansion of 1.4%. They
made no mention of a-phase expansion, but if we presume a 1% a-phase expansion, the
discontinuous expansion is 19.9%, 89% of the total. Overall, the lattice expanding
behavior of LaNi,,Fe,, was not unlike from that of LaNi,. The La,,Sm,,Ni,,Fe,, sample
showed more pronounced changes. The discontinuous volume expansion was found to
be quite reduced (12.9%, 77% of the total) and much more anisotropic (Aa/a = 4.9%,
Ac/c = 2.4%), and the total volume expansion (16.7%) is one-third smaller than that of
LaNi,. In this sample, some discontinuous lattice expansion is replaced by continuous
expansion, but it is not as prominent as in even the LaNi,Co alloy. Fe is known to
preferentially occupy the crystallographic 3g-site, but it does not induce a second plateau
in the isotherm as Co does.’ It is more likely that Fe-substitution promotes the exclusion
of hydrogen sites, and hence reduces lattice expansion, rather than slightly changing the
occupation energies of hydrogen sites.” The latter situation would be consistent with a
lattice expansion continuous with hydrogen composition.

Another outcome of the reduction in hydriding volume expansion is a reduction in the
microstrain and the associated diffraction peak broadening. Oddly enough, the
anisotropy in the diffraction peak broadening of dehydrided alloys disappears at moderate

substitution levels.” This will be further discussed with the results.

J.H.N. van Vucht, F.A. Kuijpers, H.C.A.M. Bruning, Philips Res. Rep., 25 (1970):
133. |

A. Percheron-Guégan and J.-M. Welter in Topics in Applied Physics: Hydrogen in
Intermetallic Compounds I, Chapter 2, L. Schlapbach, ed., Springer Verlag (Berlin:
1988).

William Hume-Rothery and G.V. Raynor, The Structure of Metals and Alloys,
(Institute of Metals: London) 1962.

J. Shinar, D. Shaltiel, D. Davidov, and A. Grayevsky, J. Less-Common Met., 60
(1978): 209.

H.H. van Mal, K.H.J. Buschow, and F.A. Kuijpers, J. Less-Common Met., 32 (1973):
289.

M. Brouha and K.H.J. Buschow, J. Phys. F, 5 (1975): 543.

E. Gurewitz, H. Pinto, M.P. Dariel, and H. Shaked, J. Phys. F: Met. Phys., 13 (1983):
545.

M. Latroche, J. Rodriguez-Carvajal, A. Percheron-Guégan, and F. Bourée-Vigneron,
J. Alloys Comp., 218 (1995): 64.

J.C. Achard, A.J. Dianoux, C. Lartigue, A. Percheron-Guégan, and F. Tasset, in The
Rare Earths in Modern Science and Technology, vol. 3, eds. G.J. McCarthy, Silber,
and J.J. Rhyne (Plenum Press, New York, 1982), p. 481.

A. .Percheron-Guégan, C. Lartigue, J.C. Achard, P. Germi, and F. Tasset, J. Less-
Common Met., 74 (1980): 1.

C. Lartigue, A. Percheron-Guégan, J.C. Achard, and F. Tasset, J. Less-Common Met.,
75 (1980): 23.

J. Lamloumi, A. Percheron-Guégan, J.C. Achard, G. Jehanno, and D. Givord, J.
Physique, 45 (1984): 1643.

J.-M. Joubert, M. Latroche, R. Cerby, R.C. Bowman, Jr., A. Percheron-Guégan, and

K. Yvon, , J. Alloys Comp., in press.

A. Percheron-Guégan, C. Lartigue, and J.C. Achard, J. Less-Common Met., 109
(1985): 287.

R. Griessen and T. Riesterer, chapter 6 in Topics in Applied Physics v. 63 Hydrogen

in Intermetallic Compounds I: Electronic, Thermodynamic, and Crystallographic

Properties, Preparation, L. Schlapbach, ed. Springer-Verlag, Berlin. 1988, p. 219.
J.C. Achard, A. Percheron-Guégan, H. Diaz, and F. Briaucourt, in F. Demany (ed.),
Proc. 2nd Int. Cong. on Hydrogen In Metals, Paris, June 1977, Vol. 3, Pergamon,
Oxford, 1978, Paper 1E12.

C.E. Lundin, F.E. Lynch, and C.B. Magee, J. Less-Common Met., 56 (1977): 19.
D.G. Westlake, J. Less-Common Met., 91 (1983) 1.

M.H. Mendelsohn, D.M. Gruen, and A.E. Dwight, J. Less-Common Met., 63 (1979)
193; D.M. Gruen, M.H. Mendelsohn, and A.E. Dwight, Advances in Chemistry Series
167 Transition Metal Hydrides, R. Bau, ed., American Chemical Society, Wash. DC,
(1978): 327.

A. Percheron-Guégan and C. Lartigue, Materials Science Forum, 31 (1988): 125.
K.A. Gschneidner, Jr., T. Takeshita, Y. Chung, and O.D. McMasters, J. Phys. F, 12
(1982): L1.

M. Latroche, A. Percheron-Guégan, and F. Bourée-Vigneron, J. Alloys Comp., 265
(1998): 209.

M. Latroche, A. Percheron-Guégan, Y. Chabre, C. Poinsignon, and J. Pannetier, J.
Alloys Comp., 189 (1992): 59.

S. Ono, K. Nomura, E. Akiba, and H. Uruno, J. Less-Common Met., 113 (1985): 113.
E.MacA. Gray, E.H. Kisi, and R.I. Smith, J. Alloys Comp., in press.

7 Y, Nakamura, K. Sato, S. Fujitani, K. Nishio, K. Oguro, and I. Uehara, J. Alloys
Comp., 267 (1998): 205.

E. Electronic Structure

There has been speculation about the influence of the electron structure of LaNi, on
its hydriding properties since work on this material began. Van Vucht hypothesized that
that La “induces nickel to become also a hydrogen-bonding element.” The capacity of
the hydride phase has been associated with the valence electrons of the alloy.” Griessen
and Driessen have developed a semi-empirical band-structure model correlating the heat
' of formation of the hydride and a characteristic band-structure energy parameter.”

There have only been three band-structure calculations of LaNi, alloys, and only two
of these include their hydrides. Malik, et al.’ used the augmented plane wave (APW)
method ° to calculate energy bands and densities of states for LaNi, and the isoelectronic
GdNi,. M. Gupta used a tight-binding recursion scheme for the calculations, and
included calculations for LaNi,H,.° Later calculations by Gupta used the density
functional theory in the local density approximation, using the self-consistent linear
muffin tin orbital method within the atomic sphere approximation. The latter work
included LaNi,M, M € {Fe, Co, Mn, Cu, Al}, LaNi,,Sn,,, and LaNi,CoH,.’ Results from
both authors match well with experimental data.

Experimental results on the electronic structure of LaNi, include measurements of

heat capacity “"° and magnetic susceptibility, as well as direct DOS measurements by

12-13 10,16

core-level and valence band “” photoemission spectra. Heat capacity and

magnetic susceptibility

measurements have been made for B-LaNi,H,. However, the
tendency of AB, MH alloys to desorb hydrogen in vacuum has made direct measurements
of their electronic structure impossible. Therefore, the best information available about

the hydrides of these alloys comes from band-structure calculations.

Band structure calculations and core-level photoemission spectroscopic
measurements confirm that there is little or no charge transfer between Ni and La in the
binary alloy. The band structure looks very similar to those of metallic Ni and La, with
the major differences resulting from the decrease in metal-atom coordination and increase
in Ni-Ni and La-La distances. The Fermi energy E, falls close to the end of the Ni 3d
bands, leaving them unfilled. Substitutions for Ni with metals having s-p valence
electrons leads to a decrease of the DOS at E, associated with a progressive filling of the
Ni-d states.’ In these ternary alloys (Sn,, and Al,), low energy structures are found at
~-9 eV (Sn) and ~-7 eV (Al) that are attributed to the s states of the substituent.’

In the hydride, calculations predict a charge transfer between La and H, inducing a
hybridization of the Ni-d and H-s states. The metal-hydrogen bonding states appear
5-10 eV below E,. The Fermi is found at a higher energy than in the alloy, and the DOS
is found to be smaller at E, than in the alloy.’ It is possible that the s-derived states block

the metal-bonding states by filling the DOS at that energy.

' J.HLN. van Vucht, F.A. Kuijpers, and H.C.A.M. Bruning, Philips Res. Rep., 25
(1970): 133.

* _ K.A. Gschneidner, Jr., T. Takeshita, Y. Chung, and O.D. McMasters, J. Phys. F, 12

(1982): L1.

R. Griessen and T. Riesterer, chapter 6 in Topics in Applied Physics v.63 Hydrogen

in Intermetallic Compounds I: Electronic, Thermodynamic, and Crystallographic

Properties, Preparation, L. Schlapbach, ed. Springer-Verlag, Berlin. 1988, p. 219.

* §.K. Malik, F.J. Arlinghaus, and W.E. Wallace, Physical Review B, 25 (1982): 6488.

* FJ. Arlinghaus, Physical Review, 186 (1969): 609.

° M. Gupta, J. Less-Common Met., 130 (1987): 219; M. Gupta and L. Schlapbach,

chapter 5 in Topics in Applied Physics v. 63 Hydrogen in Intermetallic Compounds I:

Electronic, Thermodynamic, and Crystallographic Properties, Preparation, L.
Schlapbach, ed. Springer-Verlag, Berlin. 1988, p. 139.

M. Gupta, °in Hydrogen in Semiconductors and Metals, eds. N.H. Nickel, W.B.
Jackson, R.C. Bowman, and R.G. Leisure (Materials Research Society, Warrendale,
PA 1998) 79.

S. Nasu, H.H. Neumann, N. Marzouk, R.S. Craig, and W.E. Wallace, J. Phys. Chem.
Solids, 32 (1971): 2779.

T. Takeshita, G. Dublon, O.D. McMasters, and K.A. Gschneidner, Jr., in The Rare
Earths in Mod. Sci. and Technol., ed. G.J. McCarthy, J.J. Rhyne, and H.B. Silber
(Plenum Press, New York 1980) p. 563; T. Takeshita, K.A. Gschneidner, Jr., D.K.
Thome, and O.D. McMaster, Phys. Rev. B, 21 (1980): 5636.

D. Ohlendorf and H.E. Flotow, J. Chem. Phys., 73 (7980): 2973.

L. Schlapbach, J. Phys. F., 10 (1980): 2477; L. Schlapbach, C. Pina-Perez, and T.
Siegrist, Solid State Commun., 41 (1982): 135.

L. Schlapbach, Solid State Commun., 38 (1981): 117.

F.U. Hillebrecht, J.C. Fuggle, P.A. Bennet, Z. Zolnierek, and Ch. Frieburg, Phys. Rev.
B, 27 (1983): 2179.

J.C. Fuggle, F.U. Hillebrecht, R. Zeller, Z. Zolnierek, P.A. Bennet, and Ch. Frieburg,
Phys. Rev. B, 27 (1982): 2145.

J.H. Weaver, A. Franciosi, D.J. Peterman, T. Takeshita, and K.A. Gschneidner, Jr., J.
Less-Common Met., 86 (1982): 195.

Y. Chung, T. Takeshita, O.D. McMasters, and K. A. Gschneidner, Jr., J. Less-
Common Met., 74 (1980): 217.

J. Palleau and G. Chouteau, J. de Phys. Lett., 41 (1980): L227.

F. Alloy Degradation

1. Disproportionation

One of the main barriers to the widespread use of LaNi, for reversible hydrogen
storage is the fast degradation of the active material with cycling. The hydride of LaNi,,
or any intermetallic hydride, is not in thermodynamic equilibrium. The heat of formation

of H with La is great, causing the reaction

LaNi,H, > LaH, + 5Ni + 2H, [11.4]

called a disproportionation reaction, to be thermodynamically favorable. The heat of

formation of this reaction is -224.6 kJ (mol H,)’,’ compared with AH’... = -273 kJ (mol

LaNis

LaNi,)' * and AH,

hydride

= -15.6 kJ (mol H)’. With extended cycling, LaNi,
disproportionates into LaH,, 2 < y < 3, and Ni metal. This can also take place during
thermal aging of the hydride, when B-phase material is kept at high temperatures for
extended periods.

Disproportionation was described by P. D. Goodell as follows:

Alloy disproportionation is the result of the metastable character of many
pseudobinary hydride phases.’ It is a separation of alloy components which is
induced by hydrogen: the more electropositive alloy elements (lanthanum,
calcium, titanium, etc.) tend to form regions of increased hydride stability. This
leaves the less electropositive alloy elements (nickel, cobalt, iron, etc.) as regions
of decreased hydride stability. The regions in question may be of atomic
dimensions as in lattice disorder or clustering, or they may be much larger if
phase separation occurs. The hydriding properties may be expected to vary
accordingly, but a general result is the loss of useful reversible hydrogen storage
capacity. In some cases the capacity can be recovered by removing the hydrogen
during mild thermal treatment.’

The first experimental verification of the disproportionation of LaNi, was obtained by
Cohen, et al. The hydrogen content and residual hydrogen over La, Eu, ,Ni,.Mn,, during

1500 thermally induced hydrogen absorption-desorption cycles were found to be

‘consistent with reaction 11.4. In addition, Eu Mossbauer spectroscopic measurements of
the cycled sample found a stable phase containing divalent Eu that did not return to Eu”
upon hydrogen desorption. This “fixed” europium hydride coexisted with the reversibly
hydrided material.' Subsequent experiments including “Ni Méssbauer ° and
magnetization ° measurements of thermally cycled LaNi, confirmed the existence of Ni
microprecipitates. The presence of trapped hydrogen in thermally cycled LaNi, was later
confirmed by inelastic neutron scattering, but there was not strong evidence for the
presence of LaH,.’ It was not until Bowman, et al. performed XRD and TEM on
thermally cycled LaNi, that evidence of the LaH, product of the disproportionation
reaction was found.’ XRD patterns of alloys degraded over 50% in reversible H capacity
contained diffraction clearly indicating the presence of LaH,. Dark-field TEM images
using the (111) and (200) f.c.c. diffraction rings revealed diffractions from regions only a
few nanometers in size.’

The fact that LaNi, does not immediately decompose in the presence of hydrogen gas
rests on the large energy barrier that must be overcome to dissociate the Haucke phase
CaCu, structure. For a hydrogen atom to bond with La, a free La atom must be present.
Because no unbonded La atoms are present in LaNi,, the first atoms to disproportionate
will be those in the least stable environments, at the alloy surface.

When the La atoms at the alloy surface have left the alloy and bonded with hydrogen,
a La depleted region is created near the surface. This concentration gradient drives more
La atoms to diffuse to the surface, continuing the disproportionation reaction. A
schematic of this process is depicted in Figure I-20, reprinted from L. Schlapbach, et
al’ Once an appreciable amount of La has left the lattice, the large concentration of
vacancies will destabilize the remaining Ni structures. The added strain associated with
lattice dilatation and La atom diffusion will break these structures off, exposing fresh
alloy surface to hydrogen. The rate limiting step in this process should involve the

kinetics of metal-atom diffusion.

CCpl (0861) EL “IAW UounUoD
-SSIT ‘f “Te 19 ‘yoeqdeyyos “] wor poyurder) ‘uoneuonsododstp *INe’y] Jo neways §=7Z-]] aNs1y

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<— UOUONUIIUOD

Metal-atom diffusion should follow Fick’s First law:

J= >. [11.5]

That is to say that atoms in a potential gradient, in this case a gradient in chemical
potential caused by the concentration gradient of the species, will diffuse in a direction
that will oppose the concentration gradient. There are several ways in which the flux of
La atoms to the alloy surface can be enhanced. Increasing either the diffusion coefficient
D or the potential gradient do/dx will enhance the diffusion of lanthanum.

In LaNi, La has a lower surface energy than Ni, providing a driving force for La
diffusion to the alloy surface.””” Therefore the situation described here is that of diffusion
in a potential field. Evidence of this has been given in the form of Auger Electron
Spectroscopy (AES) depth profiling.’ In this experiment a sample of polycrystalline
LaNi, previously exposed to air was introduced to an ultrahigh vacuum of 10° torr. The
sample was then sputter cleaned until the surface composition was approximately equal to
that of LaNi,. Immediately after sputter cleaning, the surface concentrations of La and O
were found to increase while the Ni concentration decreased slowly. After 20 minutes,

the La to Ni concentration ratio at the surface was ~1:1.7.""

“Surface segregation is a well known phenomenon in_ intermetallic
compounds.” The free energy of an intermetallic compound is minimized if the
surface concentration of the element with the smaller surface energy is increased.

In LaNi, stability increases if La segregates to the surface...”

The second factor enhancing the diffusion of La to the surface of LaNi, is the
presence of large strains at the interface between the region of hydrogen in solution in the
metal lattice and the metal hydride. As discussed in §.II.C.3, LaNi, experiences strains of
up to 7.5% during hydrogen absorption. The presence of such large strains during
hydriding will facilitate the atom movements required during La diffusion, thereby

creating a larger, “dynamic” diffusion coefficient. For the large (1.83A radius) La atom

to hop from its lattice position to a possible vacancy position, the Ni atoms separating the

‘two locations must move far from their equilibrium positions. The largest hole in the
equilibrium LaNi, lattice has a radius of 0.75 A. Each Ni atom would have to move ~1.1
A to accommodate the La atoms. This corresponds to an activation energy of ~10
eV/atom. At the interface between the a- and B-phase, the lattice misfit strains should be
large enough to enhance the diffusion of La atoms. In addition, the strains induce
microstructural defects such as dislocations and stacking faults, as explained above (§
IL.C.5). The presence of dislocations is known to enhance the kinetics of metal
diffusion, so the high concentration of such defect structures in activated LaNi, will

enhance the diffusion of La.

2. Corrosion

If hydride disproportionation limits the attractiveness of LaNi, for consumer hydrogen
storage applications, corrosion by OH ions in an alkaline electrolyte almost prevents its
usefulness in Ni-MH rechargeable batteries. During cycling in the 6 molar KOH solution

electrolyte of Ni-MH batteries, LaNi, corrodes according to:

LaNi, + 3 OH” — La(OH), + 5Ni + 3e [IT.6]

LaNi, + 3 H,O —> La(OH), + 5Ni + 3H [11.6]

The free enthalpy of formation of this reaction is -472 kJ (mol LaNi,)’, ° which is
~4.4 times the heat of hydride formation assuming 7 H (mol LaNi,)” and almost 2 times
the formation enthalpy of the compound LaNi,, -273 KJ H (mol LaNi,)" *. It is likely that
the very large heat of formation driving the transformation to La(OH), means that the
kinetics of the alloy degradation is no longer simply controlled by metal atom diffusion.

_ Hydroxide ions quickly form a La depletion layer that may be thick enough to spall or
crack with each new hydriding cycle. This decrepitation prevents the non-conducting

hydroxide films from inducing a charge transfer resistance at the particle surface. If

La(OW, developed more slowly, it might not spall off during subsequent cycles. The
charge transfer resistance would then create an over-potential on discharge, preventing
active hydrogen storage capacity from being utilized, and the corrosion layer would be
enhanced. It is ironic that the fast corrosion of LaNi, in KOH is partially responsible for
its usefulness in a battery.

The corrosion is thought to take place during discharge, and mostly when the
hydrogen content of the anode is low. It is then that the potential of the anode is lowered,
thermodynamically favoring corrosion, and the kinetically favorable reaction of OH with
H’ is hindered by the low concentration of H at the anode. The corrosion products
La(OH),, and Ni have been identified by Sakai, et al. in cycled electrodes.’ Boonstra, et
al. detected Ni(OH), as well, and proposed a corrosion mechanism whereby LaNi,

corrodes to La(OH), and Ni(OH),, followed by a reaction of LaNi, with Ni(OH),.”°

3. Approaches to Suppress Degradation

Many approaches have been explored to slow the degradation of LaNi, during
hydrogen absorption/desorption cycling. These have included microencapsulation of the
LaNi, powder with Ni or Cu, ” altering the concentration ° or composition ” of the

22-24

electrolyte, pretreatment of the alloy powder, * changing the alloy microstructure,
Pp ging y

2,15,25-28

alloying substitutions for Ni and La, and optimizing the alloy stoichiometry.” The
lifetime stabilizing mechanisms associated with many of these methods have been

proposed, and sometimes refuted, and some of these will be discussed below.

a. Alloy Composition

As stated in the introduction, Willems, et al. decreased the capacity fade of LaNi,-
based MH rechargeable batteries by alloying substitutions for Ni and for La. To quantify
the improvements in lifetime, they used the parameter S,,,, the relative fraction of the

residual storage capacity left after 400 electrochemical cycles. Figure I-21, reprinted

CEL (L861) 621 “J2W uowMmo)-ssaT ‘f ‘Moyosng
‘CHS pue sursy[iy “9 'f'f Wor poyurider) “SurpiupAy uodn uotsuedxa oumnjoa % sa °”’G [Z-J] aN

(H)AIAV 72 CC 02 Bl SL 71 ras Ot

Li LU Ul v ig Li

‘from ref. 15, shows Sin VS % volume expansion upon hydriding. They noticed that
reduced volume expansion in LaNi,,Co, alloys containing small amounts of silicon or
aluminum (dashed line) was twice as effective in promoting electrochemical stability as
the same reduction in LaNi, Co, alloys (solid line). (It should be noted that although they
did differentiate between discontinuous and continuous volume expansion, the expansion
values in Figure II-21 refer to total volume expansion.) To put this another way, adding
small amounts of Si or Al to the Co-substituted compounds did not appreciably change
the nature of the volume expansion, but did increase the stability of the compounds
during electrochemical cycling. To justify this difference, they hypothesized that, “[t]he
beneficial influence of small amounts of aluminum or silicon on the electrode stability
was interpreted as being due to the formation of a more or less closed surface layer
surrounding the fine particles, which gives them additional protection against the
electrolyte. This surface layer can be thought to be built up of islands of cobalt or nickel
interspersed with aluminium- or silicon-containing oxides. The protective effect of the
surface layer will be present only when the volume expansion is not too large.” Their
final alloy, La, ,Nd,,Ni, Co, ,Si, ,, retained a residual storage capacity of 88% of the initial
capacity of 292 mAh/g after 400 cycles. They did find more stable compounds, but the
low initial capacity of these alloys was considered to limit their use.”

T. Sakai, et al. tested LaNi,.M,, and LaNi,M alloys, M € {Mn, Cu, Cr, Al} and
LaNi, .Co,, and determined that their effectiveness in improving the cycle life increased
in the order Mn, Ni, Cu, Cr, Al, and Co.” They found some correlation between the
capacity fade and the total volume expansion, but this was clearly not the only effect.
Correlations were also found between capacity fade and both the alloy hardness, and its
BET surface area after cycling.

Reducing the volume expansion upon hydride formation has been followed by many

27,28

researchers in the pursuit of long-lived MH alloys. It is clear that the problem is not

the volume expansion itself. A uniform and gradual dilatation will not produce

decrepitation and defect structures. It is the sharp strain gradient at the o-/B- interface
associated with the discontinuous volume expansion that causes lattice misfit strains.
One method used to extend MH lifetimes is not to reduce the total volume expansion,
which often reduces the total hydrogen capacity, but rather to reduce the discontinuous
volume expansion. As explained above (§ I.D.3) alloys of LaNi, will often continue to
expand in the B-phase, absorbing more hydrogen as they expand. In this way, the misfit
strain at the interface between the o- and B-phases will correspond to the smaller
discontinuous lattice expansion while the total lattice expansion continues to be large.
Notten, et al. found that a reduction in the discrete volume expansion was more
important than a reduction in the alloy’s maximum volume expansion. In-situ XRD
determined that the volume expansion of LaNi, CuH_, alloys were all approximately
equivalent, but the difference in unit cell volumes between the a- and B-phases of these
alloys varied greatly.” On electrochemical cycling of these and LaNi,, Cu, alloys, they

found that alloys with the smallest discrete volume expansions had the longest lifetimes.”

b. Mischmetal

In addition to altering the energetics or time scale of the phase decomposition during
cycling, the possibility of alloying LaNi, increases the likelihood of consumer use of AB,
Ni-MH batteries from an economic standpoint. La is a relatively abundant metal, the
least expensive of the rare earths at $150/kg 1994 prices (99.9% purity),” but the
purification of La from: mischmetal (Mm) is an energy intensive and costly processing
step. Mischmetal is a naturally occurring ore containing a mixture of the 4f rare earth
metals. Mischmetal ores from different mines contain different proportions of the rare
earths present in the purified mischmetal. As of 1993, China was thought to hold 51% of
the world’s store of RE ores. The predominant RE minerals found there, bastnaesite and

monazite, have La:Ce:other RE ratios of 23:47:30 and 32:50: 18.

The large amount of Ce present in mischmetal raises the plateau pressure of the metal
hydride by slightly less than an order of magnitude, depending on the Ce composition.”
This effect can easily be offset by alloying with metals such as Mn, Al, or Sn.” The

*°° chemical composition,” and solidification

lanthanide ratio,” alloy stoichiometry,
technique (and resulting microstructure) “ of Mm-based alloys have all been varied to
optimize their lifetimes during electrochemical cycling. The alloy composition
commonly used in the commercial production of Ni-MH cells is MmNi,,,Co,..Mn,,Al,,,
often denoted MmB,, (B = Ni,..Co,,,Mn, ,Al,,). Cells made with these alloys experience
a capacity reduction of only 10% after 2000 cycles.”

The volume expanding behaviors of LaB, and MmB, alloys have also been measured.
LaB, alloys have a miscibility gap between y = 0.6 D/f.u. and y = 3.4 D/f.u. and
corresponding discrete volume expansion of 9.1%.” The isotherms of MmB, alloys have
almost no plateau region, and therefore a smaller miscibility gap, but well-annealed alloys
still experience a discrete volume expansion of ~9.5%.* Lichtenberg, et al. “ and Ziittel,
et al. “ measured discrete volume expansion in a number of multicomponent MmB,
alloys. They found a general trend linking cyclic stability to low volume expansion, but
the substituted elements used and their compositions had an equally large effect.
Solidification techniques such as gas atomization or melt spinning can have the effect of
further reducing the miscibility gap.”

The increased lifetimes of these alloys is likely not only due to chemical effects from
the alloying elements. Including a variety of elements in the alloy will create
configurational disorder on the LaNi, lattice. Disorder on the crystalline lattice would
create disorder in the o- to B-hydride transition by creating a continuum of chemical
environments for H to occupy in the Haucke phase lattice. This would replace the
singularity in hydrogen concentration at the alloy/hydride interface with a hydrogen
concentration gradient. This effect could be produced with many metal elements,

exclusive of their other effects on the hydriding properties of the alloy. Creating

configurational disorder as outlined here would thus lower the critical temperature for
hydrogen ordering, reducing the miscibility gap between the a- and B-hydride phases.
There are several ways to lower the critical temperature for hydrogen ordering. All of
the possible methods. involve changing the regularity of the sites available for hydrogen
occupation. This can be performed by alloying substitutions for Ni and/or La, by
quenching in atomic disorder in the lattice, and by preparation of a non-stoichiometric
alloy. The third method, preparing non-stoichiometric alloys, is actually a combination
of the other two methods. Changing the La:Ni ratio means that La is a substituent for Ni,
or Ni (a pair of Ni atoms actually) is a substituent for La. LaNi, is a stoichiometric
compound at room temperature. This means that non-stoichiometric alloys must be
quenched below the alloys’ glass transition temperature so that the additional
' stoichiometry will not precipitate out of the LaNi, matrix. The precipitate would be Ni in
the case of Ni-rich alloys, or a lanthanum-rich binary compound (La,Ni, or LaNi,) in the
case of La-rich alloys. Alloying different metals for Ni creates disorder in the hydrogen
sublattice, but sometimes it changes the environment of the hydrogen sites too drastically.
This may result in the loss of hydrogen capacity instead of resulting in a region of the

isotherm that is outside the 2-phase region.

c. Thesis Approach

The methods of reducing the volume dilatation and lowering the critical temperature
for hydride formation both help extend the lifetimes of AB, alloys during cyclic hydrogen
absorption/desorption. However, these are both mainly mechanical effects. If these results
can be generated with alloying elements that also increase the chemical stability of the
alloy (i.e., through the thermodynamics or kinetics of phase transformation), the lifetimes
of the MH alloys would be greatly enhanced. There has been some evidence of these
chemical effects in the literature, but the sources of these effects are extremely hard to

extract from other causes. Most chemical differences will be associated with mechanical

effects that result from the substitution of a metal with a different metallic radius than Ni.
This will change the plateau pressure, volume expansion, and H environments, each of
which may affect the cyclic lifetime.

Willems, et al. saw evidence of chemical effects with Al and Si substitutions in
LaNi, ,Co,,, but incorrectly attributed the effects to the formation of structures on the
alloy surface.” XPS experiments by Schlapbach, et al. determining surface compositions
have refuted the existence of an Al or Si rich surface layer in similarly substituted alloys.’
Sakai, et al. correlated an increase in lifetime with a decrease in the alloy hardness.

To determine the chemical effects of Ni substitution with respect to the cyclic
lifetimes of LaNi, MH alloys, trends in cyclic lifetime of Ni-substituted LaNi, alloys were
extracted from the pre-1993 literature. The lifetime effect of each Ni substitution was
quantified by taking the ratio of the number of cycles the substituted alloy completed
before 50% degradation to the number of cycles the binary alloy from the same
publication/research group completed before 50% degradation. This ratio was plotted vs
a variety of parameters that described the alloy, such as hardness or total volume
expansion, or the substituting element, such as metallic radius or the heat of formation of
the element M with La, AH, ,,,, Two examples of these correlations are shown in Figure
II-22, vs metallic radius, and F igure II-23, vs AH,,,. The best trend was found in the
case of AH, >

Because the AH“*”” of reactions II.4 and II.6 are much larger than AH,,,,.’, it is not
likely that the alloying will appreciably affect the thermodynamics of the degradation of
LaNi, during cycling. Indeed, calorimetry by Pasturel, et al. has shown for LaNi,M
alloys M € {Mn, Cu, Fe}, (AH ay - AHiw)/ AH, = 8-6%." To understand this
correlation, we borrowed insight from work on the kinetics of metal atom diffusion.

Our approach to reducing the alloy degradation is to reduce the kinetics of metal-atom

diffusion by substituting for Ni metals which have a large heat of formation with La. Dr.

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Correlation of AB, alloy cyclic lifetime with metallic radius of solute atom.

Figure IJ-22

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B. Fultz used Monte Carlo simulations to study the kinetics of short- and long-range
ordering in binary and ternary alloys.”*“° He found that small substitutions with a ternary
element could affect the activation barrier for metal atom diffusion and the time scales for
order-disorder transitions. *

This was first tested in the LaNi, system by thermally cycling LaNi,, LaNi,,Sn,,, and
LaNi,,Sn,, alloys from room temperature to greater than 200° C. It was found that even
these dilute substitutions of Sn for Ni provide dramatic improvements to the cyclic
lifetime of the alloys." However, a recent measurement by Joubert, et al. shows that the
total volume expansion of Sn,, is 22.4%, a reduction of only 2.2% from that of LaNi,.”
In addition, the wide plateau of Sn,, (Ay = 5.3) implies that almost all of the volume
expansion is discontinuous. It is reasonable to think that this small reduction in lattice
expansion and the associated misfit strains alone should not produce such a marked
improvement in cyclic lifetime. This positive result encouraged us to test this hypothesis
with other metal elements. Because we used electrochemical cycling, the faster

degradation mentioned above enabied us to test many alloys.

* RL. Cohen, K. W. West, and J. H. Wernick, J. Less-Common Met., 70 (1980): 229.

* _ T. Sakai, M. Matsuoka, and C. Iwakura, in Handbook on the Physics and Chemistry
of Rare Earths, eds. K.A. Gschneidner, Jr., and L. Eyring, 21 (1995): 133.

° K.H.J. Buschow and A.R. Miedema, in Proc. Int. Symp. On Hydrides for Energy
Storage, Geilo, A.F. Andresen and A.J. Maeland, eds., (Pergamon Press: Oxford,
1978), p. 235.

* PP. D. Goodell, J. Less-Common Met., 99 (1984): 1.

° —_H. Rummel, R.L. Cohen, P. Gutlich, and K.W. West, Appl. Phys. Lett., 40 (6) (1982):
477.

° R.L. Cohen, R.C. Sherwood, and K.W. West, Appl. Phys. Lett., 41 (10) (1982): 999.

719

M.J. Benham and D.K. Ross, Zeitschrift fiir Physikalische Chemie Neue Folge, Bd.,
147 (1986): 219.

RC. Bowman, Jr., C.H. Luo, C.C. Ahn, C.K. Witham, and B. Fultz, J. Alloys Comp.,
217 (1995): 185.

L. Schlapbach, A. Seiler, F. Stucki, and H.C. Siegmann, J. Less-Common Met., 73
(1980): 145.

L. Schlapbach, chapter 2 in Topics in Applied Physics v. 67 Hydrogen in
Intermetallic Compounds II: Surface and Dynamic Properties, Applications, L.
Schlapbach, ed. Springer-Verlag, Berlin. 1992, p. 72.

L. Schlapbach, A. Seiler, H.C. Siegmann, T. van Waldkirch, and P. Ziircher, Int. J.
Hydrogen Energy, 4 (1979): 21.

S. H. Overbury, P.A. Bertrand, and G. A. Somorjai, Chem. Rev., 75 (1975): 547.

J. J. Burton, E. Hymann, and D. G. Fedak, J. Catalysis, 37 (1975): 106.

A.D. Vasilyev, Thin Solid Films, 289: (1996) 205.

J.J.G. Willems and K.H.J. Buschow, J. Less-Common Met., 129 (1987): 13.

A.H. Boonstra and T.N.M. Bernards, J. Less-Common Met., 155 (1989): 119.

H. Ishikawa, K. Oguro, A. Kato, H. Suzuki, and E. Ishii, J. Less-Common Met., 107
(1985): 105; 120 (1986): 123.

T. Sakai, H. Ishikawa, K. Oguro, C. Iwakura, and H. Yoneyama, J. Electrochem. Soc.,
134 (3) (1987): 558.

A.H. Boonstra and T.N.M. Bernards, J. Less-Common Met., 161 (1990): 355.

S. Mukerjee, J. McBreen, G. Adzic, J.R. Johnson, J.J. Reilly, M.R. Marrero, M.P.
Soriaga, M.S. Alexander, A. Visintin, and S. Srinivasan, J. Electrochem. Soc., 144
(1997): L258.

N. Kuriyama, T. Sakai, H. Miyamura, H. Tanaka, I. Uehara, F. Meli, and L.
Schlapbach, J. Alloys Comp., 238: (1996): 128.

L. Zhang, TJ . O’Hara, and G.M. Michal, in Hydrogen and Metal Hydride Batteries,
P.D. Bennett and T. Sakai, Ed., PV 94-27, p 45, The Electrochem. Soc. Proceedings
Series, Pennington, NJ (1994).

W.K. Hu, J. Alloys Comp., 279 (1998): 295.

R.C. Bowman, C. Witham, B. Fultz, B.V. Ratnakumar, T.W. Ellis, and I.E. Anderson,
J. Alloys Comp., 253 (1997): 613.

B.V. Ratnakumar, C. Witham, R.C. Bowman, Jr., A. Hightower, and B. Fultz, J
Electrochem. Soc., 143 (1996): 2578.

C. Witham, A. Hightower, B. Fultz, B.V. Ratnakumar, and R.C. Bowman, Jr., J
Electrochem. Soc., 144 (1997): 3758.

M.P. Sridhar Kumar, K. Petrov, W. Zhang, A.A. Rostami, S. Srinivasan, G.D. Adzic,
J.R. Johnson, J.J. Reilly, and H.S. Lim, J. Electrochem. Soc., 142 (1995): 3424.

F. Meli, A. Ziittel, and L. Schlapbach, J. Alloys Comp., 231 (1995): 639.

P.H.L. Notten, R.E.F. Einerhand, and J.L.C. Daams, J. Alloys Comp., 210 (1994):
221.

J.J.G. Willems, Philips J. Research, 39 (1) (1984): 1.

T. Sakai, K. Oguro, H. Miyamura, N. Kuriyama, A. Kato, and H. Ishikawa, J. Less-
Common Met., 161, 193 (1990).

P.H.L. Notten, J.L.C. Daams, and R.E.F. Einerhand, J. Alloys Comp., 210 (1994):
233.

J.B. Hedrick, J. Alloys and Comp., 250 (1997): 471.

P. Falconnet, J. Alloys and Comp., 192 (1993): 114.

K.R. Clay, A.J. Goudy, R.G. Schweibenz, and A. Zarynow, J. Less-Common Met.,
166 (1990): 153.

R. Balasubramaniam, M.N. Mungole, and K.N. Rai, J. Alloys Comp., 196 (1993): 63.
G.D. Adzic, J-.R. Johnson, J.J. Reilly, J. McBreen, S$. Mukerjee, M.P.S. Kumar, W.
Zhang, and S. Srinivasan, J. Electrochem. Soc., 142 (1995): 3429.

T. Sakai, H. Miyamura, H. Kuriyama, H. Ishikawa, and I. Uehara, J. Alloys Comp.,
192 (1993): 155.

Y. Fukumoto, M. Miyamoto, H. Inoue, M. Matsuoka, and C. Iwakura, J. Alloys
Comp., 231.(1995): 562.

A. Ziittel, D. Chartouni, K. Gross, P. Spatz, M. Bachler, F. Lichtenberg, A. Folzer,
and N.J.E. Adkins, J. Alloys Comp., 253 (1997): 626.

T. Sakai, H. Miyamura, N. Kuriyama, H. Ishikawa, and I. Uehara, Zeitschrift fiir
Physikalische Chemie, Bd., 183 (1994): 333.

J. Latroche, A. Percheron-Guégan, and G. Bourée-Vigneron, /. Alloys Comp., 365
(1998): 209.

F. Lichtenberg, U. Kohler, A. Félzer, N.J.E. Adkins, and A. Ziittel, J. Alloys Comp.,
253-254 (1997): 570.

A. Pasturel, C. Chatillon-Colinet, A. Percheron-Guégan, and J.C. Achard, J. Less-
Common Met., 84 (1982): 73.

B. Fultz, J. Chem. Phys., 87 (3): (1987): 1604.

B. Fultz, J. Mater. Res., 7 (4) (1992): 946.

J.-M. Joubert, M. Latroche, R. Cerby, R.C. Bowman, Jr., A. Percheron-Guégan, and

K. Yvon, J. Alloys Comp., in press.

lll. Methods: Experimental Techniques, Equipment, and Data

Analysis

A. Alloy Preparation

Alloys were made by melting in a silver-boat induction furnace stoichiometric
amounts of La (99.99%), Ni (99.999%), and one of the solute elements Sn, Ge, In, Al, Si,
Ga, Sb, or Bi (99.999%) purchased from Alfa-Aesar, in an argon atmosphere. Alloy
ingots were then sealed in quartz ampoules that had been purged and back-filled with
argon. Annealing was performed at a temperature of 900° C for 3 days, after which the
alloys were allowed to slowly cool inside the furnace. Alloys were then activated by gas-
phase hydrogen absorption. Some alloys were further activated by the measurement of
gas-phase isotherms or by thermally driven absorption/desorption cycling. All handling
after annealing was performed in an argon glove box, except electrode preparation which

was done in air. In the text, alloys of composition LaNi, M_ will be denoted as M..

B. Gas-Phase p-c-isotherms

Unless otherwise noted, the isotherms shown here were measured with an automated
Sievert’s apparatus designed, constructed, and operated at Caltech. The equipment was
also used for hydrogen sorption activation, thermally induced hydrogen cycling, and
hydrogen charging of the alloys before poisoning the alloy surface with CO. It has
experienced several design as well as mechanical changes, and I will only describe the
most recent version here.

A picture of the Sievert’s apparatus appears in Figure III-1, and a schematic drawing
in Figure III-2. All tubing and fittings used were 1/4” 316 stainless steel with Swagelok

VCR fittings unless otherwise noted. Manual valves were Nupro bellows valves, VS-

S apparatus.

1eve

Picture of S

Figure II-1

Reactor>

Temperature
Bath

500 O

Copper Reactor
~35 cc

SS Reactor
~28.4 cc

CO Reactor
~ cc

Figure III-2. Schematic of Caltech Sievert’s apparatus.

SIC. Pneumatic valves were actuated by small solenoid valves. Pressure measurements
were made with SETRA pressure gauges. Isotherms could be measured at high pressures
(up to 500 psia) at 0.1 psi accuracy or at low pressures (100 psig) at 0.01 psi accuracy.
These gauges were calibrated by mounting a 600 psig manual Heise gauge on the system.
Zero pressure was read by a Granville-Philips convectron gauge, giving a zero reading of
< 1x 10° torr. Reactor and room temperatures were monitored by K-type thermocouples
from Omega. Isotherms at elevated temperatures were measured with the sample reactor
immersed in water in a Polystat Model 1268-44 Circulator, accurate to +0.5° C.

The manifold and reactor volumes were calibrated with high-purity argon or helium
gas. V,,, was initially calibrated at Aerojet Aerospace Corp. and was used as the
reference volume in gas-expansion volume measurements. It was possible to vary the
manifold volume by opening manual valves isolating calibrated volumes. The hydrogen
used for measurements was Matheson ULSI hydrogen (99.9999%). It was connected to
the Sievert’s apparatus with electropolished SS tubing to a brass regulator. A Tribodyn
oil-free shaft driven mechanical pump manufactured by Danielson was used to evacuate
gas from the system. A hot air gun was used to heat the sample reactor to a maximum of
300° C for sample bake-out.

The analog readings from the pressure and temperature gauges were converted by a
Strawberry Tree PCI board. This was interfaced to a Macintosh SE computer. The SE
read and recorded the pressure data and operated the relay-sclenoid-pneumatic-valve
chain used to control the pressure of the hydrogen gas in the apparatus. The program

used to operate and read the equipment was written in C.

Figure II]-3

Sample reaction chamber.

The sample reactor, shown in Figure HI-3, was a double walled tube made of 1/16”
copper pipe electron-beam welded to a 3/4” 316 stainless steel male Swagelok VCR
flange. Several VCR porous metal filter gaskets with pore size 0.5 um were used to
prevent the sample from migrating into the system and contaminating the valves.
Although the copper reactor has much better thermal conductivity than previous stainless
steel sample reactors, the kinetics of hydrogen sorption in this system was dominated by
thermal transport in the metal hydride particles and gas diffusion through the filter
gaskets. When measuring isotherms, sufficient time (20-40 min) was allowed for the

hydrogen pressure in contact with the sample to equilibrate before recording the pressure.

C. X-ray Diffraction

X-ray diffraction (XRD) was used to obtain information about alloy lattice
parameters, phase composition, grain size, and retained lattice strain. Diffraction patterns
were taken from activated samples and electrochemically cycled materials. Most of the
work was performed with an Inel CPS-120 diffractometer using Co Ka radiation
(Aq, = 1.788965A). In the following section, I will discuss some of the physics that goes
into the intrinsic diffraction profile. Then I will explain the characteristics that are
particular to the diffractometer used in the work shown here. Finally, I will discuss the
methods of data analysis I used to extract information about the sample structure and
microstructure from the diffraction patterns.

X-ray diffraction is one of the most powerful techniques available to the materials
scientist because the amount of information gained is relatively high when compared to
the amount of effort put forth to analyze the material. However, many physical processes
affect the diffraction pattern, and these must be taken into account when interpreting the
resulting data.

An experimentally measured XRD pattern h(x) is considered to be the convolution of

two components, the intrinsic diffraction profile f(x) and the instrumental factors g(x):

88
h(x) = f(x) * g(x). [10.1]

1. Pure diffraction profile

When a monochromatic, unpolarized electric field is incident on a crystalline lattice, the
scattered wave can be calculated as:
lattice basis

y (Ak)= »y exp{- i 2nAk er, »> atom Te) Setom Ne) exp(— i 2nAk er, ), (.2]’

tT, k

This is the diffraction equation, and can be separated into a structure factor:

basis

F (Ak)= 9! Baum F,) Foam (hi, JexP(-i 20 Ak er, ), [uu.3)°
and a shape factor:
lattice
S (Ak)= Sexp(-i 2nAk e(r, -s)) . [111.4]?

Ak = diffraction vector (hkl).
r, =unit cell basis vector (X,..VeomZatom)*

Som = Occupancy of unit cell site at r,.

From = atomic scattering factor of atom at r,,.

r, covers all unit cell positions in the defect-free crystal.

s =deviation, or defect, vector.

In XRD, the structure factor controls the magnitude and peak positions of the

diffracted intensity, while the shape factor describes the distribution of the diffracted

intensity in k-space.

a. Structure

Information about the atomic structure of a crystalline sample can be extracted from
the positions and intensities of a set of diffraction peaks. The wavelength of the radiation
used and the unit cell basis vectors determine the locations of the diffracted intensity.
From the locations, 8, of the peaks, the spacing of the diffraction grating generating these

peaks can be calculated using Bragg’s Law:
4./2=d,, * sin), [0.5]

where A = x-ray wavelength and d,,, = crystalline lattice spacing. In a hexagonal lattice,
the lattice spacing for a particular set of planes (hkl) is calculated from the unit cell

parameters as:
d,,. = (h’ + h*k*V3/2 +k’) a’ +c’. [11.6]

When the intensities of the scattered wave are measured by an X-ray diffractometer, a
variety of geometric factors must be included to calculate the measured intensity of the

scattered wave: >
1.=S*F,.P*T*LP*J*P [TH.7]
‘hki hkl

S =scaling factor (incident beam intensity, time length of pattern, etc.)
F,,, = structure factor

g, = lattice site occupation factor

f, = atomic scattering factor

T = temperature factor = exp(-B/sin@/A)
1+ cos” (26)cos (26,,)

> 2
sin’ @cos@

LP = Lorentz-polarization factor =

J = multiplicity of crystal plane

P = Preferred orientation (factor)

: a
The atomic scattering factor, temperature factor and the Lorentz-polarization factor all

induce asymmetries in the resulting peak profiles.

i. Lattice parameter extrapolation

The effect of errors in the diffraction peak position on the calculated lattice parameter

can be determined by taking the derivative of [TII.5):
-d,,, * cos(0) * AO = Ad,,, * sin(®)

-cot(@) * A®@ = Ad, / d,,. {HI.8]

So, lim 6x, Ad,,, = 0. In a well-aligned asymmetric diffractometer, the largest error in 20

position results from sample displacement and transparency (see § III.C.2.e), A26 = s *

sin(2@), so [III.8] becomes:
- cot(®) * 2s sin(@)cos(8) = -2s cos (0) = Aa,,,/a,,.- [111.9]
by extrapolating values of d,,, vs cos’(®), we should obtain a line whose intercept with

cos’(8) = 0 is the lattice parameter a,,..

b. Microstructure
The kinematic theory of x-ray scattering shows that particle size

and lattice distortion are diffraction order independent and
dependent, respectively, enabling the separation of the two
effects...

T. Ungar, A Borbély, Appl. Phys. Lett., 69 (21) (1996):

3173.
I open the section on microstructure with this optimistic statement because that is the
approach that I took when I began the analysis of the XRD patterns included here.
Unfortunately, I had neither the experience nor equipment with enough precision

necessary to analyze the MH samples tested here to this extent. I did accomplish some

crystallite size/microstrain analysis, so I will include the formalism here.

The formalism of diffraction broadening resulting from crystallite size and

microstrain is easily seen through manipulation of Bragg’s law. Taking the derivative of

[11.5],
0 = Adsin@-—dcos@.

= = A@cot 0 [I11.10)
A A A2
and - OcosO cos@ _ A cos 0 [0.11]

dsin@ A 7 A

In this case, A stands for the breadth of a parameter, not the deviation from its true value,
as in [11.8]. Strain broadening can be directly related through Ad/d, while size

broadening must involve the shape function of the diffraction equation, eq III.4.

i. Crystallite Size

Scherrer first pointed out that as crystallite size decreased below about lum, the

integral breadth (8) of the diffraction profile would increase according to

Boe

tTcos@.

[T1.12]

where 7 is the apparent X-ray crystallite size.”

ii. Microstrain

There are many methods to measure crystallite size and strain in polycrystalline
materials by X-ray and neutron powder diffractometry. The first measurements of strain
by XRD peak broadening were performed by Scherrer * and Warren.’ The understanding
of lattice strain and size broadening effects were later simplified and expanded by

Williamson and Hall.°

It is now routine to determine lattice strain by Rietveld refinement of diffraction
‘patterns. The most important step in determining strain from powder diffraction patterns
is to create a physical model by which lattice strain can be connected to diffraction peak
broadening. Small crystallite size broadens diffraction peaks by reducing the size of the
coherent domain that scatters X-rays. The effect of crystallite size on diffraction patterns
can be seen in equation [111.8]. The extent of this microstructural damage can be
measured with XRD by plotting the peak broadening in k-space (dAk) of diffraction
peaks vs the diffraction vector Ak.

Microstrain also broadens the specimen profile according to

B= ketanO [u1.13]

where € represents the microstrain and k is a constant whose value depends on the

definition of microstrain used.

2. Inel CPS-120 Instrument Broadening Effects

Several characteristics inherent to the Inel diffractometer, shown in Figure II-4,
make acquisition and analysis of its diffraction patterns different from those of patterns
recorded with a conventional Bragg-Brentano diffractometer. Some of the differences
that will be addressed here include: the effects of absorption by the sample, sample

displacement, and the lack of a focusing or even a parafocusing geometry.

a. Background for Curved Position Sensitive detectors

To date, many attempts have been reported to build cylindrical or curved position-
sensitive gaseous detectors for X-ray crystallography.” Of these, Ballon, et al." were
the first to employ a thin metal blade as anode, which has the advantage of being rigid
and relatively easy to construct. Their detector, of 20 cm radius of curvature 60° angular

opening, works in ‘self-quenching streamer’ mode in place of the normal proportional

Picture of Inel diffractometer.

Figure III-4

. iS . . . . . .
mode as in other examples.” A good signal-to-noise ratio obtainable in this mode of

gas ionization has provided a high angular resolution of some hundredths of a degree.”
b. Speed of acquisition

The gas-filled proportional counter used to detect diffracted photons allows data to be
collected simultaneously over 120 degrees 20.'' This means that diffraction patterns can
be acquired in a fraction of the time necessary with traditional 6-26 diffractometers. This
is particularly useful for ascertaining what phases are present in a sample and
qualitatively examining variations in microstructure within a set of similar samples. In
addition, the Inel is rather versatile. Changing sample configuration is elementary, and
changing source radiation only slightly more complicated. However, the Inel also
introduces several uncertainties and instrument effects that make quantitative analysis of

XRD patterns quite difficult.

c. 20calibration

The detector has 4096 electrical bins or channels within its ~120° 20 range. This
gives it an approximately constant step size of 0.032 °20/channel. Unfortunately, the
spacing of the channels on the detector knife edge is not constant, and varies from
detector to detector. This means there is some deviation from linearity in the 20
calibration. To determine this deviation, I used a LaB, standard, NIST SRM 660, to
calibrate the 20 range. An example of the LaB, raw data is shown in Figure II-5. |

The first step in the 26 calibration is to fit the diffraction peaks to gaussian profiles vs
channel number. The LaB, peak positions in 20 angle as determined by Bragg’s law with

A < = 1.79026 A are then fit vs peak position in channel to a sixth order polynomial.
This was used as an initial 26 calibration and Rietveld refinement was performed. Once

the instrument function (including X-ray incident angle and monochromator setting ee
al

discussed in § III.C.2.d.i and § IlI.C.2.f) and sample microstructure were determined, an

XRD pattern was calculated using these parameters. The peaks in this calculated pattern

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“were fit to gaussian profiles to determine their experimental peak positions in 20 angle.
These peak positions were plotted vs the previously obtained positions in channel, and a
final 20 calibration curve was determined. This process was repeated if the
diffractometer alignment was changed. This procedure was a bit lengthy, and was
automated by code written in the programming environment of the Mac application Igor
v. 3.13. I should stress that this only need be followed for the extraction of quantitative
results from the XRD patterns.

Figure III-6a shows the deviation from linearity of the 20 values of the Inel used in
this work, and Figure III-6b shows a published example of Inel detector non-linearity."
The 26 scaling wave is approximated well by the sixth order polynomial:

20 = 0.52 -—1.1 x 10°x’+6.8 x 10%%*-1.5x 10°x'+1.1x10"x", [1.14]
but a higher order polynomial was often used. This shape is preserved when the 20 range

of the detector is changed.
d. 26 geometry

i. absorption correction

The sample goniometer of the Inel is in a 20 geometry. As seen in Figure III-7, the
angle of incidence, o, is fixed, while the direction of the diffracted beam, 20, is variable.
One effect of this is that the path length of radiation through the sample, absorption of the
radiation, and hence the intensity of the diffracted beam, has a functional relationship on
the incident angle of the source radiation and on the direction of the diffracted beam.

Figure ITII-8 shows the radiation path for a beam diffracted at angle 20. The volume
irradiated in an element of thickness dz at a depth z is AV = (A/sina)*dz, where A is the
area of the incident beam. Absorption along the path length z/sina reduces the incident
intensity at z to Lexp(-z/sina), where 1. = sample absorbance. The diffracted beam path
length is z/sin(20-a), so the diffracted beam intensity is further attenuated by the factor

exp(-yz/sin(20-a)). If b denotes the fraction of the incident beam that is diffracted, the

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2000

1000

Channel #

Non-linearity of channel spacing for Caltech Inel detector.

Figure III-6a

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101

intensity of the radiation diffracted from volume element AV at a depth z at the

diffraction angle 26 is then:

HZ HZ
= bly exo( — =) = sin(20 — a) Je U.S]

If the sample thickness is much larger than the X-ray extinction length, we can let
t — oo, The effect of sample absorbance can be found by integrating the diffracted

intensity from z = 0 + co:

I= yb f=

1 1)
A ex -ue ‘(+ sin(26 — a } ig f (1.17)

1,4b ( sin(26-a) _\
u \sina+ sin(26- a)

In =

[Tl.18]

The effect of absorption in the 26 sample geometry can be seen by plotting the

absorption correction factor

(8c -( sin(29- a) _)

sin a + sin(20 — a) )

{Tl.19]

vs 20. Figure III-9 displays the absorption correction factors for several glancing angles.

ii. peak broadening from reflection asymmetric geometry with flat-plate sample

Because @, the angle of incidence, does not change (~15°), the diffractometer is not in
a parafocusing geometry. Instead, the geometry is what is known as reflection
asymmetric. As shown in Figure III-10, radiation that is diffracted at angles less than o
is focused in front of the detector and that diffracted at greater angles is focused behind
the detector. This adds an additional instrumental peak broadening function, shown in
Figure III-11 in the parafocusing approximation, to be convoluted with the sample
intrinsic diffraction profile. There are no Rietveld refinement codes that include this type

of instrument effect, and I did not attempt to change any code to offer this capability.

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103

Focusing

R41/(2sina) = Re = L/(2sin(26—-«)

. . ~
incident X-ray
beam

Figure III-10 Focusing circle of 20 diffractometer.

A=h’ (L-R3)/L

Figure Ifl-11

h’ =h sin(20—a)

Peak broadening of 20 diffractometer.

104

105

However, it is reasonable to assume that this effect can be taken care of by its inclusion in

the instrumental broadening function.

é. flat-plate Bragg-Brentano sample orientation.

; The flat plate approximation to the Bragg-Brentano orientation was chosen as the best
sample configuration for this work. There are several outcomes of this arrangement that
will change the positions and shapes of the diffracted lines. Sample transparency and
displacement from the goniometer center will both increase the path length of the incident
radiation before diffraction, resulting in a 20 offset of the diffraction peaks that is
proportional to sin(26), as can be seen in Figure III-12. Axial divergence contributes to
an asymmetric broadening of the diffraction peaks that will have a small effect on the
position of the diffraction peak. The first effect is easily treated in Rietveld refinements
by refinement of a sample displacement parameter. Axial divergence is related to the
vertical divergence of the incident and diffracted beams, and is best treated by a first
principles approach. Larry W. Finger, et al. have developed an analytical approximation
to the convolution of the axial divergence with a diffraction profile which is available in

some Rietveld refinement codes.”

f. monochromator

In this section I will discuss the horizontal and spectral divergence of the incident
beam. The incident beam monochromator on the Inel is pyrolytic graphite, which has a
mosaic spread of ~0.5°. This is useful for separating the Kq radiation from the
Bremstrahlung and Kg energies, but the Ky, radiation cannot be separated from the Kg,
radiation. This results in a horizontal divergence of the radiation. This addition to the
peak broadening is included in the instrument function. The Ky,/Kq, intensity ratio is
dependant upon the monochromator setting, and is not necessarily equivalent to the
theoretical value. This value was refined from the LaB, calibration pattern and the

refined value was fixed in subsequent data refinement.

901

‘Ig\UIOJOVIFJIP
gz 10j Aouaredsuey pur yuouisoedsip ajduues wroy sio1o uomsod yesg = Z[-[]] omnsty

107

3. CO Poisoning

An important property of the alloys studied here is the volume expansion upon
hydrogen absorption. Unfortunately, equipment that would allow measurement of XRD
patterns at high hydrogen pressures was not available. To measure the volume
expansion, the method of Johnson and Reilly ” was used to “poison” the surface layers of
hydrided alloys with CO. This creates a lanthanum carbonate film on the powder surface,
allowing the alloys to retain their hydrogen composition when exposed to air. After the
poisoning treatment, the alloys lost hydrogen slowly so that diffraction patterns could be
recorded at various hydrogen compositions.

Poisoning the alloys was performed by first activating them by hydrogen absorption,
then desorbing hydrogen from the samples to specific compositions. The sample
chamber was then removed from the Sievert's and relocated to the poisoning stand,
shown in Figure III-13. There it was immersed in liquid nitrogen. While the sample
cooled, the dead space of the poisoning apparatus was back-filled and purged with CO at
~5 atm. After sufficient cooling, the sample was exposed to the 300cc volume filled with
CO at 5 atm. The sample chamber was closed when the CO was done condensing,
decreasing the effective volume by a factor of 10. The I-N, was allowed to boil off,
slowly bringing the sample back to room temperature. When the sample had reached
room temperature, the pressure of CO in the chamber was 40-45 atm.

When the sample had achieved room temperature, XRD patterns were taken. The
samples lost hydrogen gradually, and many quick diffraction patterns were initially taken
after air exposure to ensure that hydrogen was not being lost during measurement. This
would have the effect of changing the position of the diffraction peak during

measurement, making it broader than the sample intrinsic profile.

tand.

isoning $

CO po

Figure III-13

108

109

4. Data Analysis

Most of the data were analyzed by hand, by fitting the diffraction peaks to gaussians
to determine their peak positions, half-widths, and intensities. Lattice parameters were
obtained by performing a extrapolating vs cos (@). The uncertainties mentioned above
prevent quantitative phase analyses from being performed, but peak intensities could be
extracted to determine the growth of the hydride or other phases.

The program Rietan “ was used to determine the lattice parameters of hydrided
samples. The variation in hydrogen composition in the sample prevented other
microstructural information from being extracted, but I have confidence in the lattice
parameters obtained.

The Macintosh application "Igor" was used to fit the diffraction peaks to Voigt
functions and to perform subsequent analyses on peak positions and shapes. The native
Igor peak-fitting procedure was extensively modified and linked with compiled c-code to
facilitate the fitting of XRD patterns. Lattice parameters were determined by the standard
extrapolation technique of plotting d,,, vs cos @,,,- Lattice parameters were determined to
be the y-intercept of a linear fit to these data.

Because LaNi,-based alloys have a hexagonal lattice, an iterative procedure was used
to determine the lattice parameters: First, (hkO) peaks were used to calculate the a-lattice
parameter. This value was then used with (hkl) peaks to calculate the c-lattice parameter.
When the c-parameter was then used to calculate the a-parameter, it was found that the
original estimate was over- or under-estimated. The new a-parameter was used to correct
the original estimate, and a new c-parameter calculated. This continued for several
iterations and the final lattice parameter was determined by fitting successive estimates to
an exponential decay. These final lattice parameters were used to calculate their

complement to insure consistency.

110

H.P. Klug and L.E. Alexander, X-Ray Diffraction Procedures, (John Wiley & Sons,

New York; 1974) p. 618.

B. Fultz and J. Howe, Transmission Electron Microscopy and Diffractometry_of

Materials, (in preparation).

The Rietveld Method, ed. R.A. Young, (Oxford University Press, New York: 1995).

P. Scherrer, Nachr. Ges. Wiss. Géttingen, 26 September (1918): 98; Zsigmondy
Kolloidchemie, 3rd edn., (1920) p. 387.

B.L. Averbach and B.E. Warren, J. Appl. Phys., 20 (1949): 1060.

G.K. Williamson and W.H. Hail, Acta Metail., 1 (1953): 22.

S.K. Byram, B. Han, G.B. Rothbart, R.N. Samdahi, and R.A. Sparks, Advances in X-
ray Analysis, 20 (1977): 529.

D. Ortendahl, V. Perez-Mendez, and J. Stoker, Nucl. Inst. and Methods, 156 (1978):
53.

T. Izumi, Nucl. Inst. and Methods, 177 (1980): 405.

E.R. Woelfel, J. Appl. Cryst., 16 (1983): 341.

J. Ballon, V. Comparat, and J. Pouxe, Nucl. Inst. and Methods, 217 (1983): 213.
J.H. Dujin, C.W-E. van Ejik, R.W. Hollander, and R. Marx, Trans. IEEE Nucl. Sci.,
33 (1986): 388.

G.D. Alekseev, N.A. Kalinina, V.V. Karpukhin, D.M. Khazins, and V.V. Kruglov,
Nucl. Inst. and Methods, 177 (1980): 385.

M. Atac, A.V. Tollestrup, and D. Potter, Nucl. Inst. and Methods, 200 (2-3) (1982):
345.

S. Shishiguchi, I. Minato, and H. Hashizume, J. Appl. Cryst., 19 (1986): 420.

M. Evain, P. Deniard, A. Jouanneaux, and R. Brec, J. Appi. Cryst., 26 (1993): 563.
L.W. Finger, D.E. Cox, and A.P. Jephcoat, J. Appl. Cryst., 27 (1994): 892.

F. Izumi, The Rietveld Method, ed. R.A. Young, (Oxford University Press, New

York: 1995), Chapter 13.

111

» JR. Johnson and J.J. Reilly, Inorganic Chemistry, 17: 3103 (1978).

D. Microprobe Analysis

Microprobe analysis was performed on some alloys in the Geology department at
Caltech by Paul Carpenter and myself. Small ingots were embedded in epoxy and
polished in air before inserting into the SEM. Chemical analyses were performed with a

JEOL Superprobe 733 electron microprobe.

E. Electrochemical Tests

The electrochemical cell is a unique place to test chemical environments because the
polarization behavior of an electrode gives important information about the chemical
reactions taking place. Chemical reaction rates are strong functions of electrochemical
potential, and the current resulting from a reaction is proportional to the net rate of the
reaction. Interpretation of this information is complicated by the fact that several
reactions can be taking place simultaneously at different rates, but we can still extract
useful information about the chemical reactions taking place at the electrode. Several
researchers have developed kinetic models for the MH electrode.'*

As mentioned in the introduction, an alkaline rechargeable battery is also a useful
place to study the stability of MH alloys during hydrogen absorption/desorption cycling.
LaNi, degrades much faster during electrochemical cycling in an alkaline electrolyte than
in a gaseous hydrogen based environment. However, the mechanisms involved in
hydride formation and decomposition in the two situations seem to be the same. In each
case, hydrogen absorption at the alloy surface is a function of its chemical potential.
Electrochemical isotherms whose potentials have been converted to pressures via the
Nernst equation are equivalent to gas-phase isotherms.’ Hydrogen diffusion coefficients
measured by electrochemical methods are of the same order of magnitude as those

measured by proton NMR.” The faster degradation in the electrochemical environment is

112

a result of the much larger heats of formation of La with the OH” ion than with hydrogen.
We can therefore obtain information about the durability of alloys much more quickly by
cycling them electrochemically than by either thermal or pressure induced hydrogen
absorption desorption cycling.

Unless otherwise noted, all electrochemical testing was done on the following
mixture including the MH alloy. After gas-phase activation, 76 weight % (w%) MH
alloy powder was mixed with 19 w% INCO 255 filamentary nickel powder (<1 ym) as a
conductive diluent and 5 w% PTFE as binder material. This will be called the anode
mixture. Because Hg/HgO reference electrodes were used in all electrochemical tests, all
potentials given in this text will be with respect to Hg/HgO. Conversion to potentials

with respect to a standard hydrogen electrode can be performed by adding 320mV.

1. Kinetics

The half-cell is often used to make electrochemical measurements because it
eliminates or decreases effect from reactions at the cathode. This is accomplished by
using a cathode with an electrochemical capacity several times that of the anode, placing
the two electrodes far enough apart that the anode experiences a homogeneous potential
from the cathode, and using a stable reference electrode to measure the potential of the

anode.

a. Half-cell construction

A picture of the half-cell used for the measurements of kinetic parameters is shown in
Figure III-14. To construct this cell, ~100mg of the anode mixture was pressed into the
cylindrical cavity of a BAS (Bio-Analytical Systems) electrode. This electrode was
positioned as close as possible to the tip of a Luggin capillary connected to the Hg/HgO
reference electrode. The cathode was flight quality NiOOH from Eagle-Picher wrapped

in Ni ExMet (expanded metal screen). This cathode was placed at an appreciable

113

Figure III-14 —_ Picture of half-cell.

114

‘distance from the anode so that the anode would see a uniform counter potential. A
stirring bar was inserted into the half-cell and it was filled with 31 w% KOH. The half-
cell was positioned above an active stirring plate during testing to reduce concentration
gradients and so that bubbles from evolving gasses did not affect the measurements by

interference with the MH/electrolyte interface.

b. Kinetics

In studying the kinetics of Ni-MH batteries, we are investigating the reactions that
take place at the anode and their individual time constants. To begin, it is instructive to

examine the steps that take place during charge transfer to (and from) the MH anode:

1. Ionic transport (OH’) in the electrolyte and electronic transport (e) in the solid
phase (electrode).
2. Charge-transfer and hydrogen transfer reactions at the surface of the MH particles:
M* +e +H,O > MH,,, + OH’. (WLE.1]
The reduction of hydrogen from water at the alloy surface creates a concentration
gradient that drives hydrogen diffusion into the particle, leading to
3. A surface to bulk hydrogen atom transfer:
M‘H,,+M’ > M°+M’H,, [II-E.2]
where M® and M” are the surface and bulk species of the MH.
4. The diffusion of hydrogen in metal, involving nucleation and growth of the B-
phase from the a-phase:
M"H®,, > MHP... (IIL.E.3]

Secondary reactions that are not usually included in kinetic analyses are:
5. MH alloy corrosion or disproportionation:

LaNi, + 3 OH” — La(OH), + 5Ni + 3e, (TI.E.4]

115

Ni + 2 OH’ — Ni(OH), + 2e. {T.E.5}
6. Hydrogen evolution reactions:

2H,0 +2e + H,+20H, [TT.E.6]
2 M’H.,, > 2M + H,. (IILE.7]
It is generally accepted that the rate limiting step under most conditions is charge-
transfer at the MH alloy surface. Because of the high concentration of OH’ ions, charge
transport in the electrolyte will only become important at very high currents. The transfer
of surface adsorbed hydrogen to bulk absorbed hydrogen is not well understood.
However, because the activation energy for hydrogen absorption in the a-phase (~ -61 kJ
/ mol H)’ is much greater than the activation energy for the c- to B-phase transition (~-16
kJ / mol H),’ it is reasonable to assume that reaction [III.E.2] is faster than reaction
{I1.E.3]. The diffusion coefficient of hydrogen in the bulk alloy is important for some
MH electrodes.” In LaNi,, however, it is on the order of 10° cm’/s," and its effects will
not dominate in the range of the time constants of the other electrode processes. The
mechanisms of nucleation and growth of the hydride (B-) phase are not formally
understood, but it is believed that they are of the same order as the hydrogen diffusion
coefficient in the hydride. Other elements that are present in the electrochemical circuit
are the electrolyte resistance between the electrodes, the contact resistance and
capacitance between the current collector and the sample (alloy + diluent + binder), and

the contact resistance and capacitance between alloy particles.
The kinetic models developed at least include steps 1-4 outlined above. The current
density due to the charge-transfer reaction is usually assumed to be given by the Butler-

Volmer equation:

ui _ C,(0,t et nn /RT C,(0,t ole aFy /RT (WLE.8] 2

C,,(0,t) = oxidant and reductant surface concentrations.

or = Oxidant and reductant bulk concentrations.

116

a .
i, = exchange current density. F = Faraday constant.
n = number of electrons. R= gas constant.
1) = electrode over-potential. T = temperature [K].

o.= transfer coefficient (reaction reversibility).

If the solution is well stirred or currents are kept so low that the surface

concentrations do not differ appreciably from the bulk values, then [III-E.8] reduces to:

i 1
JL 8 nF (RT _ (aay RT [III.E.9] *

iy
Other phenomena accounted for include surface hydrogen coverage, modeled by a

4,16

Langmuir isotherm;'*”* particle size; spherical diffusion;*”* electrode porosity; and alloy
electrocatalytic activity.’ For our purposes, however, we will merely use the unmodified
Butler-Volmer equation to obtain the kinetic parameters, as outlined below. Three
different techniques were used to characterize the electrochemical kinetics of the alloys:

DC micro-polarization, Tafel polarization, and AC impedance spectroscopy.

i. Micro-Polarization

If the electrode over-potential is small, 1 << RT/omF, the current in eq. [III.E.9] can

be approximated through a Taylor’s expansion of the exponential to

i/i,=nFn/RT. [T.E.10]
The ratio -n/i has units of resistance and is often called the charge transfer resistance,

R, = RT / nF i,. [T.E.11)

This parameter characterizes the electrode’s resistance to charge transfer in the absence of
chemical reactions. By the same token the exchange current, i, characterizes the

electrode’s ability to pass current under near equilibrium conditions.

117

DC micropolarization measurements were performed on the alloys under
' potentiodynamic conditions at scan rates of 0.02 mV/s. The scan rate was so chosen to
provide near-steady state conditions with minimal changes in the state of charge of the
electrode or its surface conditions. Although these tests were done in open cells, the
results should be valid for sealed cells as well, because additional experiments on the
x,, = 0.2 alloy have demonstrated that the kinetics as measured by micro- and Tafel

polarization are fairly independent of the state of charge of the material.

ii. AC Impedance Spectroscopy

Electrochemical (or AC) Impedance Spectroscopy (EIS) is another powerful
technique that may be used to characterize electrochemical systems." One way to
implement this method is to control the magnitude of an applied AC potential and
measure the phase and magnitude of the resulting current for a range of frequencies. If
measurements are made over a wide-enough frequency range, then different time
constants, resistances, and capacitances corresponding to different physical processes can
be extracted. An equivalent circuit is assigned to this preliminary data. Individual circuit
elements can be assigned to particular physical phenomena by varying the electrode
construction or its conditions during testing.

The problem of determining the AC response of a simple electrochemical system with
either charge transfer or diffusion-limited kinetics has been considered by several authors.
The equivalent circuit for this situation was first proposed by Randles ” and is shown in
Figure III-15. In this figure, R, represents the electrolyte resistance, C,, is the double
layer capacitance, and Z,,* is the Warburg impedance, resulting from diffusion of the
electroactive species. EIS was first used on an ingot LaNi, electrode by Agarwal and
Orazem, who used a transport based mathematical model to analyze the results of the EIS

data.”

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119

The procedure described above was followed by Kuriyama, et al., who used the AC
impedance technique to investigate deterioration of MH electrodes during charge-

21,22

discharge cycling. A schematic Cole-Cole plot of the impedance spectrum from the
electrodes is shown in Figure III-16, reprinted from ref. 22, and the equivalent circuit for
this data is shown in Figure IlI-17. The capacitive components labeled by Q are
modeled as constant-phase elements (CPE) to describe the depressed nature of the semi-
circles." The circuit elements were assigned to their various physical effects by
analyzing impedance spectra of several series of electrodes: series 1 electrodes with
constant alloy weight and binder fraction were tested at varying depths of discharge
(DOD) and temperatures; series 2 electrodes contained different alloy weights with a
constant fraction of binder; and series 3 electrodes contained a constant alloy weight and
varying binder content.

From the results of these tests, Kuriyama, et al. determined that R, corresponded to
the resistance of the electrolyte between the working and the reference electrode, R, was a
result of the contact resistance between the current collector and the alloy, and R, and C,
were related to the alloy particle-to-particle contact. These were all effects of the battery
construction. The reaction resistance for the hydrogen transfer and double layer
capacitance were responsible for elements R, and Y,,. Zhang, et al. concluded later that

the linear region at low frequencies was an effect of the porous structure of the electrode

rather that being a Warburg impedance.’

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iii. Tafel Polarization
When the electrode over-potential is large, the second term of eq. [III-E.9] becomes

very small, and

Sa mrai Rr (II.B.12]

ig
Under this condition, the electrode potential has the exponential relationship to current

first given by Tafel in 1905

N=atblogi. [NI.E.13]
We can rewrite eq. [III-E.12] as
In— =2.303log— =anFn/ RT (IILE.14]
lo Io

In this case, the effective charge transfer resistance is the semi-log slope, 2.303RT/omF of
the asymptote of the Tafel curve at infinite current. The effective exchange current is the
intercept of the asymptote with the current axis. The Tafel exchange current can be
considered the maximum exchange current for all electrochemical processes at high over-
potentials. The transfer coefficient is the inverse semi-log slope.

The following description of Tafel polarization from Bard & Faulkner is helpful in

understanding the type of kinetic phenomena that it measures:

When electrode kinetics are sluggish and significant activation overpotentials are
required, good Tafel relationships can be seen. This point underscores the fact
that Tafel behavior is an indicator of totally irreversible kinetics. Systems in that
category allow no significant current flow except at high overpotentials, where the
faradaic process is effectively unidirectional and, therefore, chemically
irreversible.”

Tafel polarization measurements can be used to independently determine the kinetics
of charge and discharge, or absorption and desorption in the case of MH electrodes.

However, the presence of mass transfer effects in the electrolyte should dominate

123

cathodic Tafel polarization measurements. For the Tafel polarization experiments
performed here, the potential was scanned from extreme anodic (positive) values to the
cathodic (negative) values, to avoid the uncertainties arising from hydrogen bubbles
adhering to the surface of the MH electrode. The data deviate from linear behavior as 1
approaches zero, because the back reactions can no longer be regarded as negligible. The
data indicate strong mass transfer effects at high currents. The cathodic Tafel plot of
LaNi, appears to show a different slope at high overpotentials, possibly corresponding to
hydrogen evolution.

In the analysis of the Tafel polarization data, the limiting current is used to correct for
mass transport effects. The limiting currents (i,,) are measured in a separate
potentiodynamic experiment at a potential 400 mV more positive than the equilibrium
potential. The Tafel plots can be corrected for the mass transfer effects by plotting the
logarithm of i,,,/(1-i,,,/i,,,) against the electrode potential. The exchange current density
and transfer coefficients for hydrogen absorption are then calculated from the intercept
and inverse slope of the corrected cathodic Tafel plots, respectively. The corresponding
coefficients for hydrogen desorption are calculated from the corrected anodic Tafel plots.
The exchange current density is the current at which the electrode begins to exhibit Tafel,
or irreversible, behavior. “The transfer coefficient, o, is a measure of the symmetry of

yo 12

the energy barrier. It quantifies the change in the activation energy of the charge

transfer reaction for a given electrode potential.

2. Cyclic Lifetime Tests

a. Prismatic cell construction

A picture of the prismatic cell used for lifetime measurements is given in Figure III-
18. The anode was constructed by spreading an amount of the anode mixture containing

approximately 1g MH alloy onto 1 sq. inch Ni ExMet and hot pressing at 150° C and 69

124

Picture of prismatic cell.

Figure III-18

125

MPa for 30 minutes. The cathodes consisted of two ~4 sq. inch flight quality NiOOH
cathodes from Eagle-Picher wrapped in Ni ExMet, each bagged with woven nylon paper.
The anode was then bagged with woven nylon paper and inserted between the cathodes.
Teflon shims were used for compaction. A Hg/HgO reference electrode was inserted in

the cell next to the electrodes.

b. Cycling conditions

Cells were initially charged at 60 mA/g for 7 hours, ~150% of their theoretical
capacity. After charging, the cells were left at open circuit for 15 minutes to allow their
potentials to equilibrate. Then the cells were discharged at C/2 rate, or 150 mA/g, to
-0.5V, a particularly low cut-off potential for AB, alloys that insures 100% DOD. After
another 15 minutes at open circuit, the cells were charged at C/5 rate, or 60 mA/g, to
115% of their previous discharge capacity. These conditions were designed to accelerate
the capacity loss of the anodes. The -0.5V cutoff potential insured some MH alloy
oxidation at each cycle, and the high discharge current created a large over-potential that

would induce hydrogen evolution and MH alloy oxidation.

' MM. Viitanen, J. Electrochem. Soc., 143 (1993): 936.

“ W. Zhang, M.P.S. Kumar, K. Petrov, and S. Srinivasan, Abstract No. 593, May 22-
27, 1994 ECS meeting.

W. Zhang, M.P.S. Kumar, S. Srinivasan, and H.J. Ploehn, J. Electrochem. Soc., 142
(1995): 2935.

L.O. Valgen, S. Sunde, and R. Tunold, J. Alloys Comp., 253-254 (1997): 656.

G. Zheng, B.N. Popov, and R.E. White, J. Electrochem. Soc., 143 (3), 834 (1996).
B.S. Haran, B.N. Popov, and R.E. White, J. Electrochem. Soc., 145 (12) (1998):
4082.

126

T. Sakai, M. Matsuoka, and C. Iwakura, in Handbook on the Physics and Chemistry
of Rare Earths, K. A. Gschneidner, Jr. and L. Eyring, eds., Vol. 21, Elsevier Science
B. V., Amsterdam (1995), p. 133.

B.V. Ratnakumar, A. Hightower, C. Witham, R.C. Bowman, Jr., and B. Fultz, in

Aqueous Batteries, P.D. Bennett and S. Gross, Eds., PV 96-16, abs. # 0053, The

Electrochem. Soc. Proceedings Series, Pennington, NJ (1997).

F.D. Manchester and D. Khatamian, Mat. Sci. Forum, 31 (1988): 261

S. Wakao and Y. Yonemura, J. Less-Common Met., 89 (1983) 481.

D. Richter, R. Hempelmann, and R.C. Bowman, SJr., in Topics in Applied Physics:
Hydrogen in Intermetallic Compounds II, Chapter 3, L. Schlapbach, ed., Springer
Verlag (Berlin: 1988), p. 97.

Electrochemical Methods, A.J. Bard and L.R. Faulkner, Eds., John Wiley & Sons,
Inc., New York (1980).

G. Zheng, B.N. Popov, and R.E. White, J. Electrochem. Soc., 143 (1996): 435.

Y. Leng, J. Zhang, S. Cheng, C. Cao, and Z. Ye, Electrochim. Acta, 43 (1998): 1

J.M. Heikonen, H.J. Ploehn, and R.E. White, J. Electrochem. Soc., 145 (1998): 1840.
P. De Vidts, J. Delgado, and R.E. White, J. Electrochem. Soc., 142 (1995): 4006.

W. Zhang, S. Srinivasan, and H. Ploehn, J. Electrochem. Soc., 143 (1996): 4039

J.R. Macdonald, Impedance Spectroscopy, John Wiley & Sons, New York, 1987.
J.E.B. Randles, Discuss. Faraday Soc., 1 (1947): 11.

P. Agarwal, M.E. Orazem and A. Hiser, in Hydrogen Storage Materials, Batteries,
and Electrochemistry, D.A. Corrigan and S. Srinivasan, eds., PV 92-5, p. 120, The
ECS Proc. Series, Pennington, NJ (1992).

N. Kuriyama, T. Sakai, H. Miyamura, I. Uehara, H. Ishikawa, and T. Iwasaki, J.
Electrochem. Soc., 139 (1992): L72; N. Kuriyama, T. Sakai, H. Miyamura, I. Uehara,
and H. Ishikawa, J. Alloys Comp., 192 (1993): 161.

127

2 N. Kuriyama, T. Sakai, H. Miyamura, I. Uehara, H. Ishikawa, and T. Iwasaki, J.
Alloys Comp., 202 (1993): 183.

128

IV. X-ray Diffraction

_ A. Phase Composition

. One of the first things learned in materials science is that chemical composition alone
does not decide what a material’s properties will be. Hence, the first test to be performed
before alloy properties were measured was to confirm that we had Haucke phase material
with no secondary phases present. This was determined by performing X-ray diffraction
(XRD) and examining the patterns for diffraction peaks that did not correspond to a
Haucke phase crystalline lattice. Diffraction patterns for LaNi,,Sn,, LaNi,,Ge,,
LaNi,,M,,, and LaNi, M_, can be seen, respectively, in Figures IV-1, IV-2a, IV-3a and
b, and IV-4. In some cases, the alloy compositions were verified by Scanning Electron
Microscopy (SEM) and Energy Dispersive X-ray spectroscopic microprobe analysis
(EDAX).

Most of the alloys melted at Caltech were found to be single-phase. This observation
was confirmed by microprobe analysis, which found a minute precipitation of LaO
particles in all samples. Diffraction peaks from secondary phases are denoted by an
asterisk (*).

All Ge, alloys but Ge,, were found to be single phase. An example of the
microstructure of the Ge,, alloy after annealing can be seen in the SEM micrograph
shown in Figure IV-2b. The second phase, identified as having approximately LaNiGe
composition, is present as small equi-axed inclusions in the LaNi,,Ge,, matrix. The
EDAX results for Ge, alloys are displayed in Table IV-1.

The Bi,, and Sb, alloys were decidedly multiphase, as can be seen in Figure IV-3b.
Microprobe analysis revealed the Bi,, alloy to have two phases, one equivalent to LaNi,

Bi, and one of approximate composition LaNiBi. The Sb,, alloy (not included in Figure

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136

IV-3b) had three phases present, which were identified as LaNi,,Sb,, LaNi,, and
~LaNiSb. Muilti-phase alloys were not included in the subsequent testing.
XRD can also be used to identify the presence of corrosion products and the extent of

decomposition. This technique will be discussed further in section VII.B for the case of

LaNi,,Sn, ,.

3. Discussion

Included in Table IV-1 are the ratios of the measured Ge composition () and the
Ge composition scaled by the alloy stoichiometry ( * 5 / B) to the nominal
composition (x,,””). The composition of the Ge,, matrix measured by SEM microprobe
was lower in all cases than the composition trend set by the other Ge, alloys. Because the
Ge,, alloy was not single phase, we believe the solubility limit of Ge for Ni, under the
annealing conditions used here, is in the range 0.45 < x,, < 0.5. The inclusions present in
Figure IV-2b are in rows, implying that they were ejected from the melt as the alloy
solidified, resulting in a eutectic structure. During annealing, these lamellar grains would
break up and become spherical to reduce surface energy effects. Because this phase is
present is such a small quantity, we assumed that it would have a small effect on the
alloy’s properties, and Ge,, was used in further measurements.

There is an interesting trend that can be seen in the equilibrium solubilities of the
p-shell metals for Ni in LaNi,. Of the group IIIA metals, Al has a solubility limit of x ~
1.5,' Ga has a solubility limit somewhere within 1 * < x < 2,’ and In has a solubility limit
of at least x = 0.4.° In group IVA, Si,’ Ge,° and Sn’ all have solubility limits of x ~ 0.5.
In group VA, Sb has a solubility less than x = 0.1 and Bi has a solubility less than x = 0.2.
It seems then, that as the number of valence electrons (electron-to-atom ratio) and
“electronegativity” of the solute atom increases, its solubility for Ni in LaNi, decreases.
This should be strongly dependant on the heat of formation of the solute with La (AH_,,,),

which would drive the precipitation of LaM, binary alloys for those solutes with large

137

AH, ay It is also intriguing that AH,,,, is negatively correlated with the metallic radius of
the solute. Within each of the groups mentioned above, increasing the atomic number of
the solute results in an increase in AH, ,,, making the solute less stable in LaNi, (decrease
in solutility limit). However, the simultaneous decrease in the solute excess metallic
radius, (t,, - Ty)» decreases the misfit energy and lattice strain in the alloy, making the
solute more stable. It is likely that the solubility limits of elements within each group of

the periodic table stays approximately constant because of these competing effects.

B. Lattice Parameters of Dehydrided Alloys

1. Literature Survey

A review of the literature was made to find previous measurements of the lattice
parameters of LaNi,.M, metal alloys. The unit cell volumes of transition metal
substituted LaNi, has been covered in the introduction, §11.D.1. Of the non-transition
metals, Al was extensively studied as a solute as early as 1977 by the groups
Mendelsohn, et al.* and Diaz, et al.’ It is quite effective in expanding the lattice, and the
unit cell volume expands linearly with composition to x,, = 1.0. Mendelsohn, et al. made
alloys with other non-transition metals, including Sn,” In,’ Ga,” Ge," and Si.’ Unit cell
volumes were given for all alloys at compositions of x=0.4 and for In,,,, In,,,, although
graphs of unit cell volume vs x showed that all solutes expanded the lattice linearly with
composition. The crystal structure of LaNi,Ga has recently been studied.” Alloys Si,,
and Si,, were studied by Meli, et al.," and Achard, et al.’ More recently, Sn has been

extensively studied by Luo, et al.” and Percheron, et al.”

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139

2. Caltech Results

Lattice parameters of dehydrided alloys were calculated from XRD patterns taken on the
Inel CPS-120 diffractometer with Co Ka radiation. Figure IV-5 shows the unit cell
volumes measured at Caltech of Sn, alloys vs x (open squares). Sn-substitution in LaNi,
has been extensively studied, and included in this figure for comparison are unit cell
volume measurements from XRD (Luo, et al.,” diamonds) and NPD data (Joubert, et
al." circles) data. The Caltech measurements agree more closely with the NPD results.
Judging from the linear fits to the data from the literature, equation IV.1 and the Caltech

data, equation I'V.2, the lattice parameter measurements made at Caltech are accurate.

3. Discussion

Figures IV-6 - 10 show the unit cell volume data extracted from the literature
(diamonds) with measurements made at Caltech (open squares) for Ge, Si, In, Ga, and Al,
respectively. The only measurements which seem inconsistent with the rest of the data
are the Si, alloys tested by Meli, et al." (Figure IV-7, circles). While the scatter in all of
the unit cell volume measurements is 0.1 - 0.15%, the volumes for these alloys deviate by
~1% from the other measurements. These data points were not included in further
analysis. For all other alloys, the Caltech measurements are consistent with those found
in the literature, and all measurements are approximately linear with solute composition,
x. The dependencies of unit cell volumes on solute atom and composition are shown in

equations IV.1 - IV.12.

Group IVA V..” = 86.73 +10.18 *x, (TV.1]°
V,. = 86.95 + 9.97 *x,, [IV.2]
Vee = 86.85 + 2.27 * xg, {Iv.3] °
V, = 86.91 - 0.28 *x, [Iv.4] *"°

Group IIIA V, = 8685 +10.35 *x, [IV.5] *

140

V., = 86.76 + 3.84 *x, [IV.6] *
V, = 86.72 + 3.77 *x, {1V.7] *°
Transition V, = 87.01 + 150 *x, (IV.8] “
Metals Ven = 86.92 + 4.87 *x,, [Iv.9] *"°
V., = 86.83 + 2.02 *x,, [TV.10]
V.. = 86.71 + 0.59 *x, (tv.11]
Vo. = 86.55 + 1.70 *x,, [Iv.12] ”

The lattice parameters of Ni-substituted LaNi, are frequently found to follow
Vegard’s law, but this has not been related to the solute metallic radius. Although the
alloy lattice parameters may depend on several effects (magnetic, bond strength, site
occupation), the predominant controlling factor is the solute metallic radius. The effect
of bond strength should be small in transition metal-substituted alloys, because they do
not have large heats of formation with La. Fe and Co are magnetic, and order on the Ni
sublattice. As all of the p-shell metals used in this work are non-magnetic and are
expected to occupy the crystallographic 3g site, their effects on the alloy lattice
parameters should be comparable. When the lattice parameters and unit cell volumes are
plotted vs solute composition, the slopes of the lines characterize the effectiveness of the
solute in expanding the lattice. These values should be proportional to the metallic radius
of the solute, or more accurately, the difference R,-R,, (included in the table of solute
properties found in Appendix A). Figure IV-11 displays the slopes found in equations

IV.1 - [V.12 vs the difference between the solute metallic radius and that of nickel.

cuthbert

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0.1

0.0

o - Caltech measurements }

Figure IV-6 Unit cell volumes of LaNi,_,Ge, alloys vs Ge composition, x.
{ @ - literature measurements,

141

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Oo - Caltech measurements }

Figure IV-7 Unit cell volumes of LaNi._,Si, alloys vs Si composition, x.
{ @ , @ - literature measurements,

142

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146

147

C. Hydride Volume Expansion

1. Literature Survey

The lattice expanding behavior of LaNi, and its Ni-substituted alloys has been
discussed in the introduction, sections §11.C.3 and §II.D.3, respectively. There are
numerous measurements of total lattice expansions between dehydrided and fully
hydrided alloys to be found in the literature. Figures IV-12 - IV-14 show total a-axis,
c-axis, and volume expansions of ternary LaNi, alloys substituted with transition metals
vs solute composition x. Equations IV.13 - IV.16 are linear fits to the lattice expansion
data available in the literature. Lattice expansions for sp-shell metal-substituted LaNi, are

included with the results in the next section.

Aala* = 7.73 -0.09 x; Acie“ =7.47-1.00x; AV/V“ =24.71- 1.01 x [IV.13]"
Aala “= 7.59-1.28x; Acic™ =7.64-1.57x; AV/V"=24.58-4.66x [IV.14]°*
Aala* = 7.57-0.55x; Acie ™ =7.08 - 3.25 x; AV/V “= 23.88 - 5.04.x [IV.15]""
Aala™ =7.39-1.05x; Acic™ =7.01-2.38x; AV/V™ = 24.08 - 5.61 x [IV.16]}°”

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151

2. Caltech Results

As discussed in Methods (III.C.5), the lattice parameters of hydrided alloys were
measured by ‘poisoning’ the surface of the alloys with CO after they had been fully
charged with hydrogen. Evolution of the diffraction patterns of several metal hydrides as
they desorbed hydrogen are exhibited in Figures IV-15 - IV-23. Total lattice expansion
was determined between lattice parameters obtained from patterns taken immediately
after hydriding and those of dehydrided (10~ torr @ 300° C) alloys. Discontinuous lattice
expansion was determined from patterns containing diffraction peaks from both the a-
and B-phase hydrides.

Figures IV-24 - IV-26 show the a-axis, c-axis, and volume expansions of sp-shell
metal substituted LaNi,. Data from the literature are denoted by closed symbols and
Caltech measurements (both total and discontinuous) by open symbols. Total lattice
expansion values measured at Caltech are consistent with those taken from the literature.
The total and discontinuous lattice expansion values measured at Caltech in Figures I'V-
24 - IV-27 are connected by vertical lines, and total expansion is in all cases greater than
discontinuous.

The lattice expansions of Ge, alloys measured at Caltech are shown in Figure IV-27.
There is some experimental scatter in the total lattice expansion values of the Ge, alloys
(closed symbols). This can be attributed to a variation in the alloy hydrogen composition
after poisoning resulting from inconsistencies in the poisoning procedure. This is
confirmed by examining the diffraction patterns during hydrogen desorption, Figures IV-
19 - IV-23. There is no diffraction pattern of the Ge,, alloy exclusively in the B-phase
(Figure IV-22). The §-phase lattice parameters of the first Ge,,H, diffraction pattern
after hydriding were used to calculate total lattice expansion, as they are greater than
those obtained from subsequent patterns. The presence of diffraction peaks from the a

phase while the B-phase is in the continuous regime of volume expansion implies that

152

some of the alloy was passivated, and did not absorb hydrogen. The total volume
expansion measured for this alloy is smaller than that measured for Ge,., implying that
the Ge, , alloy was not at its maximum hydrogen composition when the diffraction pattern
was measured. Similarly, the total volume expansion measured for the Ge,, alloy is less
than that of the Ge,, alloy. The XRD patterns in Figure IV-20 show that this alloy
desorbed hydrogen very quickly. The diffraction pattern taken 1 day after hydriding
shows only a few small peaks from the $-phase (<5%). It is probable that this alloy lost
hydrogen before diffraction patterns were measured. Because the total lattice expansion
values measured at Caltech for M, alloys are consistent with those values found in the
literature, the results for Ge, alloys will be considered accurate except when obvious

inconsistencies are present, as in Ge, and Ge, ,.

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166

3. Discussion

In general, the reduction in total lattice expansions of hydriding are approximately
linear with substituted composition. Equations IV.17 - IV.21 are linear fits of the total
lattice expansion data to substituted composition. The somewhat non-linear behavior of
Ac/c and AV/V in Al. alloys results from the inclusion of measurements by Gruen, et al.
(x,, = 0.4 & 0.5). Only data from Ge,,, Ge,,, and Ge,, were used in the Ge, fits of total

expansion data. Because the intercepts of the Ge, fits are consistent with the total lattice

expansion of unsubstituted LaNi,, we have confidence in these fits.

AV/V* = 24.54 - 20.53x; Aa/a* = 7.61 - 2.78x; Ac/c* = 7.56 - 12.55x. [IV.17]*"
AV/V" = 24.61 - 18.46x; Aa/a” =7.61-5.61x; Ac/c” =7.60- 4.85x. [IV.18]
AV/V® = 24.52 - 14.90x; Aa/a® = 7.68 - 1.67x; Ac/e* =7.42- 9.81x. [IV.19]
AV/V“ = 24.76 - 13.59x; Aa/a® = 7.70 - 3.58x; Ac/c“ = 7.61 - 5.05x. [IV.20]*
AV/V™ = 24.18 - 12.32x; Aa/a® = 7.64 - 4.09x; Ac/c = 7.68- 3.31x. [IV.21] 7”
AV/V™ = 23.40- 9.08x; Aa/a” = 7.54 -2.72x; Acie =6.74- 2.66x. [IV.22]°

AV/V = 26.85 - 61.52x; Aa/a® = 7.73 - 6.31x; Ac/c* = 9.00 - 38.50x. [IV.23]
AV/V“= 13.61 - 2.44x; Aa/a® = 6.02 - 1.44x; Acie =1.90- 2.61x. [IV.24]

The discontinuous lattice expansion values. of Ge, alloys, however, are not at all linear
with x... The c-axis and volume expansions are substantially reduced for x = 0.2, but
subsequent substitution has a much smaller effect on the lattice expansion. The reduction
in a-axis expansion is more linear, but it also tapers off for higher compositions.
Equations IV.23 are fit to lattice expansions for x = 0.4 (Aa/a) and x = 0.2 (Ac/c and
AV/V). The slopes of these lines are much greater than those of equations IV.24, fit to x

= 0.3 (Aa/a) and x = 0.2 (Ac/e and AV/V). After some “critical” substituted composition,

167

k the lattice expansion levels off and shows only a slight reduction with increasing Ge
composition. In addition Ge-substitution reduces the lattice expansion anisotropically.

In all measurements of discontinuous lattice expansion (above and §II.D.3), the
discontinuous lattice expansion is more anisotropic than the total lattice expansion. That
is, the c-axis continues to expand much more that does the a-axis as more hydrogen is
absorbed. The intercepts of Ac/c“ (Figure IV-25) and AV/V™ (Figure IV-26) are less
than the expansion values of LaNi,, while the intercept of Aa/a™ is approximately equal to
that of Aa/a‘“”s (Figure IV-24). This implies that the measurements by Gruen, et al. were
not taken on alloys at their maximum hydrogen compositions.

A solute’s effectiveness in suppressing lattice expansion of hydriding is quantified by
the slopes of equations IV-17 - IV.22. Figure IV-28 plots the slopes of AV/V from total
lattice expansion measurements vs the AH,,,. The two parameters are certainly

correlated, but the scarcity of data limits further analysis.

891

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169

D. Microstrain of Dehydrided Alloys

1. Literature Survey

The pronounced anisotropic peak broadening present in XRD and neutron diffraction
patterns of hydrogen activated LaNi, and its substituted alloys was first noted by Achard,
et al.,” and by Fischer, et al.” Percheron, et al. later performed strain analyses of
LaNi, M. alloys using the Williamson-Hall method, as explained in §III.C.1.b.ii.” They
found prominent anisotropic strain broadening in XRD patterns of LaNi,,.Mn, and LaNi,
(Al, alloys, as seen in Figures [V-29 and I[V-30, reprinted from reference 29. Large
microstrains were found in the (hhO) and (OkO) directions and almost none in the (00/2)
direction. The microstrains and their anisotropy were found to diminish with alloy
substitutions. Mn substituted LaNi, was found to have an isotropic microstrain
distribution for compositions x, > 1, and Al relieved the anisotropy for x, > 0.2.” In
later work, this microstrain anisotropy was refined from XRD patterns with modified
Rietveld refinement codes.” MH microstrain refinement is now routine, and code for

this has been implemented in the publicly available program FullProf.”

2. Caltech Results

Our original work on lattice microstrain analysis studied the microstrain developed in
LaNi,,M,, alloys after a single gas-phase hydrogen absorption/desorption cycle.” We
found a pronounced variation in lattice microstrain with solute chemistry. The XRD
patterns taken from these alloys are displayed in Figure [V-3a. These measurements
were performed with the Inel CPS-120 in a low-resolution configuration and the Ka,/Ka,
radiation splitting was not accounted for in the analysis. The dAk vs Ak relationship is
shown in Figure IV-31 for each M,, alloy. Only reflections perpendicular to the basal
plane {(110), (200), and (220)} were used in the Williamson-Hall plot. This is consistent

170

of pcose
LaNiggn Mry og
20
23
Bcos@ LaNiggMnog
155 SP
ee fe]
sob . .2sin@
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broadening in LaNi, Mn.. Data points “1”
represent hki families of diffraction peaks
(reprinted from ref. 29).

peose LaNt.. Al
8 : .

171

t Bcos® LaNizg Aloo

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Figure [V-30

Williamson-Hall plots of anisotropic XRD line
broadening in LaNi, Al. Data points “1”
represent hki families of diffraction peaks.

(reprinted from A. Percheron-Guégan. ..)

172

with the findings by Percheron-Guégan, et al. that microstrain occurs predominantly in

27,29

that direction upon hydrogen absorption.“ When a linear fit is applied to these points,
the microstrain distribution in each sample can be correlated with the slope of the line,

and the grain size with the y-intercept.

3. Discussion

The diffraction patterns in Figure [V-3a are displayed in order of increasing heat or
formation of binary La-M compounds, AH,,,,. We can see in the insert that the widths of
the diffraction peaks seem to decrease in the same sequence. Figure IV-32 displays the
microstrains extracted from the diffraction patterns by the Williamson-Hall method. The
correlation observed in the peak breadths is confirmed by Figure [V-32, which displays
the microstrains obtained from Figure IV-31 vs AH,,,. This figure shows that the
microstrain decreases with increasing heat of formation of a compound of the solute with
La.

The significance of this measurement of the microstrain in the alloy is that it can be
used to rank the resistance of the lattice to defect generation. Alloys with low values of
microstrain resulting from hydrogen activation would either have a smail driving force
for defect generation (small lattice dilatation), or would be more resistant to this type of
defect generation. In either case, such alloys are candidates for long-lived anodes in Ni-
MH cells. The hypothesis that such a stabilization of the lattice can be produced by
alloying LaNi, with large AH,,, elements is supported by the X-ray line broadening
results. As seen in Figure IV-32, the abundance of lattice defects found in the crystalline
structures of hydrogen activated LaNi,,M, alloys is negatively correlated with IAH,,,|. If
the Haucke phase lattice is made more stable with respect to crystalline defect generation,
it is reasonable to assume that it will be more stable during hydrogen absorption-
desorption cycling. This will also be true because, as mentioned in the introduction,

lattice defects promote metal atom diffusion.

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heat of formation with La.

174

175

M.H. Mendelsohn, D.M. Gruen, and A.E. Dwight, Nature, 269 (1977): 45.

Z. Blazina and A. Drasner, J. Phys.: Condens. Matter, 10 (1998): 4777.

M.A. Fremy, D. Gignoux, J.M. Moreau, D. Paccard, and L. Paccard, J. Less-Common
Met., 206 (1985): 251.

M.H. Mendelsohn, D.M. Gruen, and A.E. Dwight, Mat. Res. Bull., 13 (1978): 1221.
J.C. Achard, A.J. Dianoux, C. Lartigue, A. Percheron-Guégan, and F. Tasset, in The
Rare Earths in Modern Science and Technology, vol. 3, eds. G.J. McCarthy and J.J.
Rhyne (Plenum Press, New York, 1982), p. 481.

C. Witham, R.C. Bowman, Jr., and B. Fultz, J. Alloys Comp., 253-254 (1997): 574.
S. Luo, W. Luo, J.D. Clewley, Ted B. Flanagan, and L.A. Wade, J. Alloys Comp., 231
(1995): 467.

D.M. Gruen, M.H. Mendelsohn, and A.E. Dwight, Advances in Chemistry Series 167
Transition Metal Hydrides, R. Bau, ed., American Chemical Society, Wash. DC,
(1978): 327.

H. Diaz, A. Percheron-Guégan, and J. C. Achard, Int. J. Hydrogen Energy, 4 (1979):
445.

M.H. Mendelsohn, D.M. Gruen, and A.E. Dwight, Inorg. Chem., 18(12) (1979):
3443.

F. Meli, A. Zuettel, and L. Schlapbach, J. Alloys Comp., 190 (1992): 17.

S. Luo, J.D. Clewley, T.B. Flanagan, R.C. Bowman, and L.A. Wade, J. Alloys Comp.,
267 (1998): 171.

J.-M. Joubert, M. Latroche, R. Cerby, R.C. Bowman, Jr., A. Percheron-Guégan, and
K. Yvon, J. Alloys Comp., in press.

T. Sakai, K. Oguro, H. Miyamura, N. Kuriyama, A. Kato, and H. Ishikawa, J. Less-
Common Met., 161, 193 (1990).

176

C. Lartigue, A. Percheron-Guégan, J.C. Achard, and F. Tasset, J. Less-Common Met.,
75 (1980): 23.

R. C. Bowman, Jr., J. S. Cantrell, T. W. Ellis, T. B. Flanagan, J. D. Clewley, and S.
Luo, in Proc. 10th World Hydrogen Energy Conf., Cocoa Beach, FL, June, 1994, pp.
1199-1207.

J. Lamloumi, A. Percheron-Guégan, J.C. Achard, G. Jehanno, and D. Givord, J.
Physique, 45 (1984): 1643.

H.H. van Mal, K.H.J. Buschow, and F.A. Kuijpers, J. Less-Common Met., 32 (1973):
289.

C. Colinet, A. Pasturel, A. Percheron-Guégan, and J.C. Achard, J. Less-Common
Met., 134 (1987): 109.

A. Pasturel, F. Liautaud, C. Colinet, C. Allibert, A. Percheron-Guégan, and J.C.
Achard, J. Less-Common Met., 96 (1984): 93.

M. Latroche, J. Rodriguez-Carvajal, A. Percheron-Guégan, and F. Bourée-Vigneron,
J. Alloys Comp., 218 (1995): 64; M. Latroche, A. Percheron-Guégan, and F. Bourée-
Vigneron, J. Alloys Comp., 265 (1998): 209.

S. Ono, K. Nomura, E. Akiba, and H. Uruno, J. Less-Common Met., 113 (1985): 113.
E.MacA. Gray, E.H. Kisi, and R.I. Smith, J. Alloys Comp., in press.

S. Bagchi, D. Chandra, W.N. Cathey, R.C. Bowman, Jr., and F.E. Lynch, subm. 1997
TMS Meeting.

J. Lamloumi, A. Percheron-Guégan, C. Lartigue, J.C. Achard, and G. Jehanno, J.
Less-Common Met., 130 (1987): 111.

W.L. Zhang, M.P.S. Kumar, A. Visintin, S. Srinivasan, and H.J. Ploehn, J. Alloys
Comp., 242 (1996): 143.

A. Percheron-Guégan, C. Lartigue, J.C. Achard, P. Germi, and F. Tasset, J. Less-
Common Met., 74 (1980): 1.

177

M.H. Mendelsohn, D.M. Gruen, and A.E. Dwight, J. Less-Common Met., 63 (1979):
193.

J.C. Achard, F. Givord, A. Percheron-Guégan, J.L. Soubeyroux, and F. Tasset, J.
Phys. (Paris), Collog. 5, 40 (1979): 218.

P. Fischer, A. Furrer, G. Busch, and L. Schlapbach, Helv. Phys. Acta, 50 (1977): 421.
C. Lartigue, A. Le Bail, and A. Percheron-Guégan, J. Less-Common Met., 129
(1987): 65.

P. Thompson, J.J. Reilly, and J. M. Hastings, J. Less-Common Met., 129 (1987): 105.
J. Rodriguez-Carvajal, "FULLPROF: A Program for Rietveld Refinement and Pattern
Matching Analysis," Abstracts of the Satellite Meeting on Powder Diffraction of the
XV Congress of the IUCr, p. 127, Toulouse, France (1990).

C.K. Witham, R.C. Bowman, Jr., B.V. Ratnakumar, B. Fultz, and S. Surampudi, in
Proc. 11th Annual Battery Conf. on Apps. and Adv., January 9-12, 1996, Long Beach,
CA, p. 129.

178

V. Isotherms
A. Gas-phase

1. Literature Survey

There are a variety of hydrogen pressure-composition isotherms of Ni-substituted
LaNi, that have been reported in the literature. Isotherm measurements have been

reported for ternary alloys of transition metals Cr,’ Mn,” Fe,'** Co,’ and Cu '**; and p-
shell metals Al,?°“ Si,?""*"* Ga,” Ge,'*"® In,°” and Sn.”* Almost all elemental
substitutions for Ni have the effects of decreasing the plateau pressure, maximum
hydrogen capacity, and pressure hysteresis of the isotherms. Notable exceptions to this
are Mn, which has almost no effect on hydrogen capacity for x < 1.0; and Co, which
exhibits ordering on the LaNi, lattice as discussed in the introduction, changing the
stability of the hydrogen sites.

Extensive work has been done on alloys substituted with Sn ”” and with Al,°"™ and
isotherm characteristics for these alloys are included in Tables V-1 and 2. The isotherms
measured by Meli, et al. of Si-substituted alloys were measured dynamically,’’ and these
results can only be compared to the current work sparingly. Isotherm characteristics for
the alloys Ga,,, Si,,, Ge,, and In, will be quoted here (Table V-3) but the
thermodynamic parameters were obtained from isotherms at only two temperatures, and

: 5,16
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182

2. Caltech Results

a. LaNi,,Ge, |

Room temperature isotherms of all Ge substituted alloys are shown in Figure V-1.
The reversible hydrogen capacity is reduced with Ge substitution, and the capacities at 2
atm. are shown in Table V-4. Both the absorption and desorption plateau pressures and
the hydrogen absorption capacities decrease with Ge substitution. Figure V-1 also shows
that the hysteresis between the absorption and desorption isotherms is decreased with
increasing Ge composition. The hysteresis ratios of the hydride formation and
decomposition plateau pressures at mid-plateau have been calculated using equation
111.2 and are shown for room temperature isotherms in Table V-4.
Elevated temperature isotherms were measured for all Ge substituted alloys, and are
presented in Figures V-2 through V-6 for alloys Ge,,, Ge,,, Ge,,, Ge,,, and Ge,,,
respectively. The pressures at mid-plateau are used in van’t Hoff plots (Figure V-7) to
calculate the enthalpies and entropies of hydride formation and decomposition for each
alloy, which are presented in Table V-4. The isotherm and thermodynamic data for x =
0.4 agree with the values reported by Mendelsohn, et al (Table V-3)."° We see that in all
cases the enthalpy of hydride decomposition is higher in magnitude than that of
formation. The absolute enthalpy also decreases with increasing Ge composition. Both
of these points are consistent with the decrease in isotherm plateau pressure. The entropy
calculations show a somewhat less consistent trend of an increase in absolute entropy
with increasing Ge composition. The enthalpies of hydride formation and decomposition

are approximately linear with substituted composition, and can be expressed by:

AH™ (x) =-14.54 - 68x, kJ, and [V.1]

AH™,.(x) = 15.26 + 5.5x,, KJ. [V.2]

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By examining Figures V-2 through V-6, we can determine the boundaries of the
dilute (a) and ordered (B) hydride phases. The room temperature plateau of
LaNi,,Ge, ,H, ‘is particularly flat, and extends to ~6 H/f.u. Elevated temperature
isotherms exhibit a decrease in capacity. The isotherm resolution is not high enough to
distinguish a splitting at y > 4, but we can see the 75° C plateau beginning to slope
upwards. Alloys with higher Ge-composition have progressively lower capacities. The
anomalous decrease in entropy for the alloy with xGe = 0.5 may be caused by the presence
of the second phase that was found in this material. The compositional trend for the
entropy of hydrogen absorption correlates with the widths of the plateau shown in Figure
V-1. The increase in plateau slopes of the isotherms as x,, increases is indicative of a

reduction in the critical temperature for hydride formation.

b. LaNi,M,

The compositions of the alloys in this set were chosen to give the alloys plateau
pressures of ~0.4 atm. This would insure that the results from electrochemical cycling
tests in our slightly pressurized cells would not be confused by hydrogen evolution, self-
discharge, etc. The compositions were chosen by reviewing the isotherm plateau pressure
data reported in the literature.

Room temperature isotherms of all alloys M, are shown in Figures V-8a and b. The
alloy Sn,, was not prepared at Caltech, it was obtained from Hydrogen Consultants, Inc.,
and its isotherm was measured at the University of Vermont.” The absorption plateau
pressures of most of these alloys are ~0.5 atm. as planned. Si,, has an absorption plateau
pressure of 0.7 atm., implying either that the solute composition of this alloy is somewhat
less than 0.4 or the values quoted in the literature are in error. The widths of the isotherm
plateaus are shown in Table V-5. Absorption and desorption pressures are taken at mid-

plateau and the hysteresis ratios calculated from these values using equation III.2 are

192

. ° * . . .
also shown in Table V-5. Hysteresis values are a maximum for Al,,, and minimum for

0.4°

3. Discussion

To note the differences between these sets of alloys, we must examine their properties
more closely. The maximum capacities of the alloys in these sets follow approximately
the same linear relationship with solute composition. These values are shown in Figure
V-9. Also included in this figure are the capacities of the M, alloys. For the most part,
the Ge, alloys have capacities that are slightly lower than the Sn, and M, alloys. With the
exception of In, , and Si,,, the M, alloys have capacities that are approximately equal to
the Sn, alloys.

In the introduction, an overview was given of the deuterium site occupation of a
variety of deuterided alloys of LaNi, and LaNi,,M,. Recall that a loss in overall capacity
was predominantly accompanied by the blocking of the 4h (1 site/cell) and 120 (2
sites/cell) sites in the 6d cluster. This corresponds with Ni-substitution at the 3g site
preventing hydrogen (deuterium) occupation in these sites. Assuming all Ni-substitution
occurs at the 3g site, an alloy M, will have an average of x 3g Ni sites per unit cell
substituted by M. The approximately linear relationship between maximum hydride
capacity and substituted composition seen in Figure V-9 implies that there is a simple
blocking of hydrogen occupation sites by the substituting metal atom. The above
considerations imply that a substituted composition of x will block 3x sites from being
occupied by hydrogen (deuterium). This is consistent with the reduction in capacity from

y = 6.0 at x = Oto y =4.5 at x =0.5 that is seen in Figure V-9.

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It is instructive to compare the isotherms of LaNi,,Ge, alloys to those of LaNi, Sn,
alloys, taken from ref. 22. Figure V-10 presents these 2 sets of data side by side. It
should be noted that the ordinate of the Ge graph covers 1 order of magnitude in pressure,
while that of the Sn graph covers 2 orders of magnitude. The alloys Sn,, and Sn,,, have
sloping plateaus, indicating some inhomogeneity in Sn distribution. The diffraction
peaks from these alloys are quite sharp, implying that this inhomogeneity does not exist
on a macroscopic scale, and likely results from non-equilibrium site occupation on the
crystalline lattice. The hysteresis between absorption and desorption curves seems quite
different, but this is an artefact of the different scaling. Otherwise, these 2 graphs are
almost identical. Each isotherm is spaced evenly throughout the pressure range, the
hydrogen capacities of correspondingly substituted alloys are almost equal, and the
sloping character at the beginnings and ends of the isotherms are markedly similar. This
leads us to believe that Sn and Ge act almost identically as alloying solutes, with the main
distinction between them being the expansion of the LaNi, lattice, dependant upon the
solute metallic radius.

When the metal hydride plateau pressures are plotted vs the alloy unit cell volumes on
a semi-log graph (Figure V-11a), we see the linear relationship noted by a number of
researchers. However, we also notice that Ge seems more effective than Sn in
suppressing the plateau pressures of the alloys. Although it takes more Ge than Sn to
expand the LaNi, unit cell a given amount, a hypothetical Ge-substituted alloy will have a
lower plateau pressure than a Sn-substituted alloy with the same unit cell volume. This
discrepancy becomes more pronounced when we realize that the logarithm of the plateau
pressure is proportional to the Gibbs’ free energy of hydride formation, and subtract the
effect of entropy. Because our hypothetical Ge alloy has a larger solute composition than
the Sn alloy, it has a lower hydrogen composition. Hydrogen absorbed into the Ge-

substituted lattice has a more restricted configuration, leading to a greater loss of entropy

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199

than hydrogen absorbed into the Sn-substituted lattice. To eliminate the effect of
differences in entropy, the enthalpies of hydride formation and decomposition should be
plotted vs alloy unit cell volume, as in Figure V-11b. We can examine the composition
weighted chemical effect of the solute as the difference between the two curves.

In the literature of ternary alloys of LaNi,, we found the following relationships

between solute composition and enthalpy of hydride decomposition:

AH" (x) = 15.48 + 9.08 x,, +0.2 kJ/(mol H) [V.3]°

AH™ (x) = 1488 + 1.14x, kJ/(mol H) [Vv.4]°

Fitting the Sn, and Al, data to linear relationships, we obtain:

AH™ (x) = -15.21 - 9.79 Xy, kJ/(mol H), [V.5] 22,23

AH™.(x) = 15.42 + 9.58 x,, kJ/(mol H), and [V.6]*”

AH™,..(X) 15.438 + 8.31 Xa kJ/(mol H). [V.7] 11,24

This relationship was not available for all solutes owing to scatter in the data or a non-
linear relationship (Co).
From these equations and the unit cell volume relationships determined in §IV.B.3,

we obtain the following linear relationships between AH™ and unit cell volume:

Absorption: AH™,(V) =-229 + 2.80V,, kJ/(mol H), [V.8]
AH" (V) = -77.1 + 1.06 V,, kJ/mol H), [v.9]°"

Desorption: AH" (V) = 193 - 241V, kJ/(mol H). [V.10]"°~
AH“ (V) = 180 - 2.25 V,, kJ/(mol H), [V.11]

200

AH (V) = 130 - 1.67V,, kJ/mol H), v.11}?
AH", (V) = 74.5 - 1.03V,. kJ/mol H), v.12)22
AH™,(V) = 34.0 - 0.56V,, _kJ/(mol H). [v.13]

The intercepts of these lines are not equal to the enthalpies of hydriding because the
volumes used are not differential volumes, (V,,— Vuni,)- The slope of the corresponding
equation V.8 — V.13, characterizes the change in hydriding enthalpy per A’ of unit cell
volume for each solute. This enthalpy vs unit cell volume parameter has been plotted vs
solute atom metallic radius and the heat of formation with La in Figures V-12a and b,
respectively. If the unit cell volume were the only factor affecting the enthalpy of
hydride formation, this parameter would not vary with solute element. From the spread
in these values (0.56 — 2.8) is clear than there are second order effects controlling the
alloy plateau pressure, but the available data are too sparse to discern any trends.

The sloping nature at the ends of the plateaus is more pronounced for solutes with
larger metallic radii, even though these alloys have smaller substituted compositions.
This is indicative of larger solutes being more effective in reducing the miscibility gap
between the o- and B-phases. It is possible that the disorder retained after the annealing
treatment used to homogenize the alloys is greater for large solutes. The interaction
energy between solute atoms would drive them to form inhomogeneous concentration
distributions, leading to a sloping plateau. It is also true that the larger strain field
resulting from larger solutes could create a distribution of H sites with a broader range of

energies.

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202

H.H. van Mal, K.H.J. Buschow, and A.R. Miedema, J. Less-Common Met., 35
(1974): 65.

A. Percheron-Guégan, C. Lartigue, and J.C. Achard, J. Less-Common Met., 109
(1985): 287.

C. Lartigue, A. Percheron-Guégan, J.C. Achard, and F. Tasset, J. Less-Common Met.,
75 (1980): 23.

W. Luo, S. Luo, J.D. Clewley, T.B. Flanagan, R.C. Bowman, Jr., and J.S. Cantrell, J.
Alloys Comp., 75 (1993): 147.

S. Bagchi, D. Chandra, W.N. Cathey, R.C. Bowman, Jr., and F.E. Lynch, subm 1997
TMS Meeting.

J. Lamloumi, A. Percheron-Guégan, C. Lartigue, J.C. Achard, and G. Jehanno, J.
Less-Common Met., 130 (1987): 111.

H.H. van Mal, K.H.J. Buschow, and F.A. Kuijpers, J. Less-Common Met., 32 (1973):
289.

J. Shinar, D. Shaltiel, D. Davidov, and A. Grayevsky, J. Less-Common Met., 60
(1978): 209.

M.H. Mendelsohn, D.M. Gruen, and A.E. Dwight, Nature, 269 (1977): 45.

M.H. Mendelsohn, D.M. Gruen, and A.E. Dwight, RE In Mod Sci & Tech, 1 (1977):
25.

J.-M. Park and J.-Y. Lee, Mat. Res. Bull., 22 (1987): 455.

T. Sakai, H. Miyamura, N. Kuriyama, A. Kato, K. Oguro, H. Ishikawa, and C.
Iwakura, J Less-Com Met., 159 (1990): 127.

203

M. Latroche, A. Percheron-Guégan, Y. Chabre, C. Poinsignon, and J. Pannetier, J.
Alloys Comp., 189 (1992): 59.

M.H. Mendelsohn and D.M. Gruen, Rare Earths In Mod Sci & Tech, 2 (1980): 593.
M.H. Mendelsohn, D.M. Gruen, and A.E. Dwight, Inorg. Chem., 18(12) (1979):
3443.

F. Meli, A. Ziittel, and L. Schlapbach, J. Alloys Comp., 190 (1992): 17.

S. Srivastava and O.N. Srivastava, J. Alloys Comp., 267 (1998): 240.

Z. Blazina and A. Drasner, J. Phys.: Condens. Matter, 10 (1998): 4777.

M.H. Mendelsohn, D.M. Gruen, and A.E. Dwight, Mat. Res. Bull., 13 (1978): 1221.
J.S. Cantrell, T.A. Beiter, and R.C. Bowman, Jr., J. Alloys Comp., 207/8 (1994): 372.
S. Luo, W. Luo, J.D. Clewley, T.B. Flanagan, and R.C. Bowman, Jr., J. Alloys
Comp., 231 (1995): 467.

S. Luo, J.D. Clewley, Ted B. Flanagan, R.C. Bowman, Jr., and L.A. Wade, J. Alloys
Comp., 267 (1998): 171

H. Diaz, A. Percheron-Guégan, and J. C. Achard, Int. J. Hydrogen Energy, 4 (1979):
445.

J. Lamloumi, A. Percheron-Guégan, J.C. Achard, G. Jehanno, and D. Givord, J.
Physique, 45 (1984): 1643.

204

Vi. Electrochemical Kinetics

As explained in Methods (§III-E), the electrochemical cell is a unique, though
intricate, environment to investigate chemical reactions. It provides a direct method to
measure chemical potentials, but it is difficult to isolate individual species and reactions
from the testing apparatus and environment. Therefore, a single electrochemical
measurement should not be used exclusively to characterize a material. We are more
likely to obtain accurate information about a material by interpreting the results of a
series of electrochemical tests in conjunction with other techniques that characterize the
microstructure of a sample. In light of this recommendation, I will present results from
three different electrochemical tests and will make use of information gained from
diffraction experiments to interpret these results. It should be noted, however, that the
point of testing the alloys’ electrochemical kinetics was to insure their viability for use in
consumer Ni-MH rechargeable batteries. Gaining a deeper understanding about the
electrochemical processes taking place at the MH anode is beyond the scope of this

thesis.

A. DC Micropolarization

1. LaNi,,Sn,

Figure VJ-1 shows the micropolarization curves of the LaNi, Sn, alloys. The values
of the exchange currents estimated from the slopes of micropolarization curves of
different MH alloys show an interesting trend (Figure VI-11 and Table VI-1). The
exchange current (i,) increases initially upon Sn substitution from 0.77 mA for the binary
alloy to 1.35 mA for the alloy with x = 0.1. Further addition of Sn decreases the
exchange current. Nevertheless, the kinetics of Sn substituted alloys are superior to those

of the binary alloy for Sn compositions of x < 0.3 in unsealed cells. With Sn

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205

206

compositions of x > 0.4, the kinetics of hydrogen absorption and desorption are slowed

considerably, and are even slower than in the binary alloy.

2. LaNi,,Ge,

Figure VI-2 shows the approximately linear micropolarization curves of LaNi, Ge,
alloys. The exchange currents estimated from the slopes of micropolarization curves
show an improvement in the kinetics of hydrogen absorption and desorption upon Ge
substitution (Figure VI-11 and Table VI-2). There is a small decrease at low Ge
concentrations compared to the binary alloy, i.e., from 0.77 mA for the binary alloy to
0.63 and 0.71 for Ge compositions of 0.1 and 0.2, respectively. At higher Ge
compositions, i.e., x > 0.3, the exchange current is higher than that of the binary alloy and
continues to increase with Ge composition. Unlike the case for Sn solutes, which cause
marginal slowing of kinetics at x > 0.4, high Ge concentrations have no adverse effects

on the kinetics of the alloy.

3. LaNi,,M,

Micropolarization measurements of the LaNi,,M, alloys can be seen in Figure VI-3.
The exchange current densities calculated from these measurements, shown in Figure
VI-11 and Table VI-3, show a trend with composition different than that seen with the
Sn, and Ge, alloys. The exchange currents increase slightly for substituted compounds
with x < 0.2. Al,,, has the maximum charge transfer exchange current measured for any
alloy. The exchange currents of alloys with higher nickel substitution decrease

monotonically with substituted composition.

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208

209

B. AC Impedance

Electrochemical impedance spectroscopic measurements were made on mixed anode
powders in the charged state. The powders were activated by gas-phase absorption but
were not electrochemically cycled. The AC impedance data were obtained in the
frequency range of 100 kHz to 5 mHz at a low AC amplitude of 2mV. The charge
transfer resistance of the electrode can be obtained by fitting the Cole-Cole representation
of the impedance spectra to the equivalent circuit shown in Figure III-17.' The observed
impedance patterns of the MH electrodes are simplified by the absence of a diffusional
component, and this element was excluded from the equivalent circuit. The parameters in
the equivalent circuit were calculated by a non-linear least squares fit using the Boukamp
method.’ To be consistent with the other electrochemical tests shown here, the charge
transfer resistance, R.,, is converted to an exchange current using equation IV.E.11. All

charge transfer exchange currents are shown in Figure VI-12 vs substituted composition.

1. LaNi,,Sn,

The impedance plots of LaNi, Sn, alloy electrodes are shown in the Nyquist or Cole-
Cole form in Figure VI-4. The observed impedance patterns of the MH electrodes are
simplified by the absence of a diffusional component. As may be seen in this figure, the
impedance decreases noticeably upon initial substitution of Sn, then increases with x.
The trend is similar to that observed in the DC polarization experiments. It is thus clear
that the kinetics of hydriding improves markedly upon Sn substitution in LaNi, Sn_,, at
least for x < 0.3. Higher amounts of Sn seem to cause sluggish kinetics for hydrogen

absorption and desorption.

210

2. LaNi,,Ge,

The impedance plots of LaNi, Ge, alloy electrodes are shown in the Nyquist form in
Figure VI-5. The figure shows that the impedance increases slightly upon initial
substitution of Ge, but decreases for x = 0.3. This trend is similar to that observed in the

dc polarization experiments.

3. LaNi, M,

The impedance plots of LaNi,.M, alloy electrodes are shown in the Nyquist form in
Figure VI-6. The figure shows that the impedance of all substituted are less than that of
LaNi,. The impedance follows a trend similar to that observed in the dc polarization

experiments. All substituted alloys have lower impedance than LaNi,. There is a

minimum vs substituted composition for Al,,,.

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C. Tafel Polarization

As mentioned in Methods (§IIL.E.1.b.iii), the Tafel polarization measurements of MH
electrodes show interference from mass transfer effects. Desorption limiting currents are
used here to correct both anodic (desorption) and cathodic (absorption) Tafel plots for the
mass transfer interference. This is in contrast to the approach of Zheng, et al.” who used
different limiting currents for charge and discharge. The latter approach might introduce
uncertainties in the cathodic limiting currents, due to simultaneous hydrogen evolution at
high negative potentials.’ Anodic Tafel exchange currents for all alloys are shown in
Figure VI-13 vs substituted composition, and cathodic Tafel exchange currents appear in
Figure VI-14. Transfer coefficients calculated from the inverse slopes of anodic and
_ cathodic Tafel polarization are plotted vs substituted composition in Figures VI-15 and

VI-16, respectively.

1. LaNi, Sn,

Figure VI-7 illustrates the Tafel behavior of LaNi,,Sn, alloys during charge and
discharge. The overpotentials at any current density can be seen to decrease upon the
initial substitution of Sn, but increase for the highest Sn concentration. The limiting
currents were measured in separate potentiodynamic measurements at a potential 400 mV
more positive than the equilibrium potential, and are listed in Table VI-1. The corrected
Tafel plots (Figure V1-8) are more linear.

The absorption exchange current (Table VI-1) increases upon Sn substitution and
shows a maximum at a Sn composition of x = 0.2. The desorption exchange current
improves more significantly upon Sn substitution, but has a maximum near x = 0.1.
Nevertheless, the kinetics of desorption continue to be better than the binary alloy for
x <0.4. The Tafel slopes show an interesting trend: the slope for the absorption process

decreases with increasing Sn concentration, whereas the slope for the desorption process

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increases (Table VI-1). The change in both the parameters is less marked for x > 0.2.
The transfer coefficients during absorption (Figures VI-15) calculated from the Tafel
_ Slopes increase with increasing Sn composition. The corresponding transfer coefficients

during desorption(Figures VI-16), however, decrease with increasing Sn composition.

2. LaNi, Ge,

Figure VI-9 illustrates the Tafel behavior of LaNi,,Ge, alloys during charge and
discharge. The limiting currents were measured in a separate potentiodynamic
experiment at higher positive potentials, i.e., 400 mV away from the reversible potential
(Table VI-2), and used to correct for the above mass-transfer interference. The diffusion
limiting current on discharge is highest for a Ge composition of 0.1 < x < 0.2 (in the
range of 500 mA/g) and is reduced at high Ge compositions.

The absorption exchange current density increases upon Ge substitution and shows a
maximum at a Ge composition of x = 0.3 (Figure VI-13 and Table VI-2). The
desorption exchange current (Figure VI-14 and Table VI-2) improves more significantly
upon Ge substitution, and also has a maximum near x = 0.3. The transfer coefficient for
the absorption process is not monotonic with substituted composition, although all Ge,
alloys have absorption transfer coefficients greater than that of LaNi,. The transfer

coefficient for the desorption process has a minimum at x = 0.3.

3. LaNi,,M,

“Measurements of the Tafel behavior of LaNi,.M, alloys are shown in Figure VI-10.
The limiting currents were estimated from the Tafel curves using trends noted with the
previous measurements. The anodic and cathodic exchange currents show the same
trend. The exchange currents increase with substituted composition except for Ga,,,,
which has exchange currents slightly lower than the other alloys. The transfer

coefficients are somewhat less regular. The desorption transfer coefficient of In,,, (0.59)

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is ‘slightly greater than that of LaNi, (0.55) and subsequent alloys have transfer
coefficients decreasing slightly with substituted composition. The absorption transfer
coefficients are ‘not monotonic with substituted composition. All transfer coefficients of
_ the M_ alloys are approximately 0.5 except for the absorption transfer coefficient of Ga,,,,

which is ~ 0.33.

D. Discussion

Because the potentials used in DC micro-polarization and AC impedance experiments
are perturbed only slightly from equilibrium (8 and 2 mV, respectively), they should
reflect the kinetics of charge transfer at the interface between the electrolyte and the MH
surface. The exchange currents measured by these two techniques (Figures VI-11 and
12) are very similar, indicating that they are accurate. For all alloys but Sn,, and Sn,,,
i,“ > i°°. The only measurements that seem anomalous are from Ge,,. As noted in
§IV.D.2, the diffraction pattern of this alloy showed substantial peak broadening and the
presence of a second phase. Although there is no obvious reason for a second phase to
make the reaction resistance different when measured by an AC technique, it will change
the resistive and capacitive behavior of the alloy. This could easily change the effective
reaction resistance measured by AC impedance.

The broad diffraction peaks of the Ge,, alloy are indicative of a grain size much
smaller than those of the other Ge-substituted alloys, resulting in a higher specific surface
area and more available sites for charge (hydrogen) transfer. This explains why the
exchange currents of the Ge,, alloy is substantially higher than other Ge-substituted
alloys. The XRD measurements of the M, alloys also showed that the grain size
decreases in the order Ge > Si > Ga > Al, resulting in a corresponding increase in surface
area. This can be used to rationalize a corresponding increase in electrode kinetics. It is

certain that other phenomena are important in charge transfer, and this trend with alloy

surface area is not consistent with all of the measurements.

221

The exchange currents calculated from Tafel polarization measurements are different
from those described above. These exchange currents should reflect the total current
from all electrochemical reactions taking place during charge (absorption) and discharge
(desorption) at high overpotentials, E-E,,,,. The reactions at least include:
charge/discharge, hydrogen and oxygen evolution, and corrosion. For Sn, and Ge, alloys
having low substituted-compositions, x, the Tafel exchange currents are higher than the
charge transfer exchange currents. These alloys have plateau pressures greater than 1
atm, and hydrogen evolution would be favored at the high over-potentials used in Tafel
measurements. The Tafel exchange currents of the M, alloys are approximately equal,
implying that the reactions occurring are equivalent. The Ge, alloys kinetics are
depressed, rather than enhanced, for x > 0.3.

The Tafel transfer coefficients characterize the change in activation energy for the
forward and backward processes during charge and discharge. It is known’ that lower
plateau pressures facilitate absorption, whereas higher plateau pressures are desirable for
desorption. This can be used to rationalize an increase in the absorption transfer
coefficient, and a decrease in the desorption transfer coefficient, with increasing
substituted composition. The plateau pressures of the M, alloys are ~0.5 atm. This
plateau should be favorable for both absorption and desorption. The transfer coefficients
of the M, alloys for both absorption and desorption are ~0.5, indicating activation barriers
changing symmetrically with electrode overpotential.

It is thus clear that the electrochemical kinetics of hydrogen absorption and desorption
improve markedly upon substitution of Ni with sp-shell metals. Because substituted
alloys have smaller surface areas than the binary alloy, substitution must affect the charge
transfer process more directly. One obvious explanation for this change is the increase in
the lattice constants of the alloys and the corresponding decrease in plateau pressure.
Alloys with lower plateau pressures and larger heats of formation for hydrogen

absorption should have higher kinetics. It is possible that on an atomic scale, a solute

222

cansing a large lattice mismatch will disrupt the atomic configuration at the surface of the
alloy, facilitating hydrogen absorption.

Unlike LaNi, Sn, alloys, the kinetics of hydrogen absorption in LaNi, Ge, alloys do
not seem to be suppressed at high solute concentrations. Instead, the kinetics appear to be
more facile in alloys with high Ge concentrations. The Tafel kinetics of the M, alloys are
similar to each other, implying that the kinetics of hydrogen evolution is similar, as
would be expected from their similar plateau pressures. The charge transfer kinetics of
the M., alloys, however, are almost all correlated with alloy grain size. As mentioned
several times, the grain size, substituted composition, solute metallic radius, and AH, ,,
are all correlated, making determination of the rate determining phenomena difficult at

best.

N. Kuriyama, T. Sakai, H. Miyamura, I. Uehara, and H. Ishikawa, J. Alloys Comp.,

192, 161 (1993).

> B.A. Boukamp, Solid State Ionics, 20, 31 (1986).

* TT. Sakai, M. Matsuoka, and C. Iwakura, in Handbook on the Physics and Chemistry
of Rare Earths, K. A. Gschneidner, Jr., and L. Eyring, eds., Vol. 21, Elsevier Science
B. V., Amsterdam (1995), p. 133.

“ G Zheng, B.N. Popov, and R.E. White, J. Electrochem. Soc., 143 (3), 834 (1996).

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232

Vil. Electrochemical Cycling

A. Cycling conditions

Before lifetime testing of the alloys began, the cycling conditions were varied using
the alloy LaNi,,Sn,,- These tests were performed both to find the optimal conditions for
cycling, to insure that our testing methods were not appreciably different in effect than
those of other researchers, and to explore the nature of the corrosion reaction and its
products. The final cycling conditions used for the cells with ~1 g MH alloy were:

i) ~C/6 charge (60 mA) to 110% of the previous discharge capacity,

ii) 15 minute open-circuit stand,

iii) ~C/2 (150 mA) discharge to -0.5V vs Hg/HgO,

iv) 15 minute open-circuit stand.

1. Effect of Amount of Hydrogen Absorption

The cyclic lifetime of LaNi,,Sn,, (# of cycles to 50% capacity) observed in our
experimental test cell is on the order of 200 cycles (Figure VII-1), whereas the
commercial Ni-MH cells are known to have a lifetime of ~1,000 cycles. The shorter
cyclic lifetimes in our studies are attributed to the ratio of the electrode capacities and the
test conditions. The commercial cells are fabricated in a positive-limited design, with
about 1.4 - 1.5 times excess capacity in the negative electrode, to provide tolerance to
overcharge reactions and to avoid hydrogen evolution at the negative electrode.
Accordingly, approximately 65 - 70% of the active MH material experiences hydrogen
absorption-desorption cycling during each cycle. The balance is used as reserve to
replace the material that degrades during cycling. This unused alloy is presumably held

at the same potentials as the utilized material with respect to the alkaline electrolyte, but

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234

would not experience the ot- to B-phase change described in §11.C.3. The present cells, on
the other hand, are negative-limited, and the MH electrode is overcharged to 110% right
from the beginning to insure maximum hydrogen absorption during each cycle.

In order to simulate the conditions experienced by alloys in positive-limited cells
using the present test cell, the MH electrode is charged only to about 70% of its
maximum capacity (measured after activation) in each cycle and discharged to a preset
cut-off voltage. This partial charge can be represented by a new variable, the state of
charge (SOC). The limits for the SOC could be either 0 to 70% or 30 to 100%. The cells
cycled from 100% capacity might experience a larger lattice volume at high hydrogen
composition, possibly enhancing La diffusion. In the former case, the degradation
resulting from volume dilatation could be less, but that due to corrosion would be more
prominent as compared to the latter case. In this experiment, therefore, three values of
the cut-off voltage were chosen for the discharge, i.e., -0.7, -0.75, and -0.8 V, while
keeping the SOC the same (0 to 70%).

An additional factor to consider when interpreting the results of this experiment is the
nature of the hydrogen absorption and resulting lattice expansion. Presumably, the first
step during charge would be to hydride all of the available (not degraded or passivated)
MH material to the limit of the o-phase,

M + YoH — MH,,. [VIt.1]

In the case of LaNi,,Sn,,, Yo = 0.3 H/LaNi, ' = 16.7 mAh/g = 5.1% of the initial capacity

of 325 mAh/g. When all of the active material has been hydrided to the limit of the dilute
phase, individual grains or particles of MH,,, will experience the transformation to MH,,,

MH,,, + (yp - Yo)H — MH,g. [Vi.2]

In the case of LaNi,,Sn,,, Yo -yp = An,,, = 5.3 H/LaNi, = 294 mAh/g = 90.5% of the

initial capacity. This means that in a hypothetical anode containing activated MH powder

that has experienced no degradation, (70-5.1)/90.5 = 71.7% of the active material will

experience the discontinuous volume expansion during each cycle. In addition, the

235

discharge cut-off voltages of -0.7, -0.75, and -0.8 V imply that the cells have some
residual hydrogen capacity at the beginning of each cycle. If the hydrogen composition
_ of the discharged MH powders is yo, then 70/90.5 = 77.3% of the active material will
experience the discontinuous volume expansion during each cycle. However, dynamic
electrochemical isotherms (Figure VII-2) show that at the current densities we use for
cycling, only 30 mAh/g or 10% of the initial capacity to 0.5V is extracted between
potentials of 0.8 V and 0.5 V, implying the alloy has reached < 10% SOC by the time it
reaches 0.8V.

Figure VII-3 illustrates the cyclic lifetime of the alloy LaNi,,Sn,, under the above
conditions, with the discharge cut-off voltages of -0.7, -0.75, and -0.8 V, respectively.
Because the initial charge state and amount of hydrogen absorption are the same in all
these cases, the amount of MH material experiencing lattice dilatation is the same. As
may be seen from the figure, the capacity decline is the same for all cut-off potentials
tested. The flat region of the cyclic lifetime curves represents the cycles in which the MH
material which has degraded is “replaced” by the reserve material. The cyclic lifetime of
cells with a partial degree of hydrogen absorption is longer than those with complete
hydrogen absorption in each cycle (Figure VII-3). Operating at low range of SOC, as in
a positive-limited cell design, would contribute to a longer cycle life, as evident from
Figure VII-3. The cyclic lifetime appears to be longer with a low end of charge SOC;
the capacity of those cells with only 70% of charge exceeds that of the cell with 100%
charge after ~90 cycles. The capacity fade of all the cells seem identical, and is therefore

not dependent upon cut-off voltage in the range —0.7 V to —0.8 V.

2. Effect of Anodic Potentials

Although varying the cut-off voltage as described above did not have much effect on
the MH alloy capacity fade, we were not convinced that anodic potentials are unimportant

with respect to the alloy degradation. The presence of the corrosion product Ni(OH), (see

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236

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section VII.B below) in electrochemically cycled electrodes implies that the local
potential can drop below the macroscopic measurement of the anode potential.

In order to examine the effect of the corrosion occurring in the discharged state, we
varied step iii, the conditions between discharge and the subsequent charge. Figure
VII-4 shows the cyclic lifetime curves of cells with 0 and 15 minutes of open-circuit
stand time, as well as a cell held for 15 minutes at the discharge cut-off voltage, -0.5 V.
It is clear from the figure that the cycle life improves upon a decrease in the open-circuit
stand time after discharge. The enhanced corrosion in the discharged state might be
responsible for this behavior.

Figure VII-5 shows the cyclic lifetime curves of cells tested under even more severe
conditions. To simulate conditions that may arise during deep discharge and over-
discharge of cells, the electrode potential was held potentiostatically at -0.3 and -0.1 V for
15 min. after discharge. Figure VII-5 shows the effect of driven oxidation on the cyclic
lifetime, compared to the standard condition of a 15 min. open-circuit stand between
discharge and charge. It is clear from Figure VII-5 that the cycle life is rather strongly
affected by the overdischarge. At the -0.3 V holding potential, serious degradation sets in
after about 30 cycles, whereas at a potential of -0.1 V, the degradation sets in right from
the beginning.

In addition to a faster degradation rate, there are large fluctuations in the capacity as
shown in Figure VII-5. These may be attributed to the surface films formed on the MH
electrode under these conditions. Due to the increased interfacial resistance, the cells
quickly attain the voltage limit set for charge. Increasing the voltage limit for charge
improved the capacity, but the fluctuations were still present. Returning to the standard
cycling conditions, however, eliminated the fluctuations, which suggests that these

fluctuations are a result of film formation and dissolution during cycling.

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3. Effect of Temperature

In order to study the effect of temperature on the corrosion of the MH alloy and thus
on its cycle life, cycling studies were carried out at four different temperatures. The
strategy for these tests was to cycle cells with alloys Sn, at temperatures such that the
alloy plateau pressure was ~0.5 atm. This ensured that the slightly pressurized character
of our test cells would not affect the results. The affect of temperature was then
determined by noting the change in cyclic lifetime from that of room temperature cycled
MH alloys. Differences in the cyclic life can therefore be attributed to different rates of
corrosion at these temperatures. The MH alloys thus chosen include LaNi,,Sn,, at 10° C,
LaNi,,5n,, at 25°C, LaNi, 50, ,, at 44° C, and LaNi, Sn, ,at 59° C.

The cycle life of these alloys at the specified temperatures is illustrated in Figure
Vil-6. As may be seen from the figure, the cycle life does deteriorate more quickly at

high temperatures, probably due to enhanced corrosion.

4. Effect of Electrode Preparation

The capacity retention that was realized in our electrodes is superior to that reported
for similar formulations with electrodes made with Teflonized carbon.’ Nevertheless, in
order to verify if higher Teflon content or lower electrode density can improve the cycle
life any further, cells were fabricated with 50 wt.% Teflonized (33 wt.%) Vulcan-X
carbon and are being cycled under identical conditions. The results thus far (Figure
VII-7) indicate that the cyclic lifetime of a cell with this alternative construction is not
any better than our standard cells. This implies that the binder concentration in our
standard electrode is enough to provide good mechanical integrity to the electrode and

that the charge densities are not too high to affect the cyclic lifetime.

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5. Conclusions |

The cyclic lifetime of LaNi,,M, alloys with selected ternary solutes is governed by
the nature of the ternary solute, as well as the test conditions. The corrosion of the MH
alloy occurs more prominently in the discharged state and in deep and/or over-discharge.
This can result in a rapid deterioration of the cell capacity during cycling. The surface
films formed therein also cause large fluctuations in the measured capacity during
cycling. The electrode fabrication methods used here have been found to be as good as, if

not better than, other reported methods.

B. Effects of Electrochemical Cycling

The next step in our research was to learn more about the effects of electrochemical
cycling on LaNi,-based alloys. Comparing cycling tests with microstructural
characterization of MH material in cycled electrodes would shed light on the processes
involved in the alloys’ capacity degradation. Therefore, cell kinetics were measured
periodically during cycling, and several cells were cycled to varying levels of degradation
and disassembled so that the phase composition and microstructure of the AB, anode
material could be determined. These tests were done with LaNi,,Sn,,, our baseline alloy,

. 3
as the anode material.

1. RESULTS

The first set of cells (‘d' - 'g") were cycled to give information about the effects of
cycling on charge transfer kinetics. The MH powder used in these cells came from a
single ingot, but the mixture used in cells ‘d' and 'f was made at an earlier time than that
used in cells 'e' and 'g'. Representative linear polarizations of cell 'd’ are given in Figure
VII-8. It can be seen that there is a precipitous decrease in charge transfer resistance after

the first 25 cycles and relatively minor decreases after subsequent cycling. There is also a

245

small increase in the equilibrium potential of the cells. Impedance spectroscopy for cell
f? (Figure VII-9) shows a similar change in electrode charge transfer resistance with
cycling. ‘After the first 25 cycles, the charge transfer resistance decreases by an order of
magnitude, and subsequent cycling produces a minor reduction in resistance.

Cycling of cells ‘d’ - ‘g’ is shown in Figure VII-10. The capacity loss the cells
experienced from cycling was appreciable, ranging from 38% to 55%. Electrodes 'e' and
‘g' were stored under argon in the discharged state for 1 month before XRD patterns were
measured. After cycling, 'd' and 'f' electrodes were over-discharged to ~20 times their
current capacities and studied by XRD. The diffraction peaks were indexed. Each X-ray
peak was fit to a gaussian line shape function superimposed on a linear background to
determine its integrated intensity. The XRD patterns, seen in Figure VII-11, of these
electrodes show the presence of LaNi,,Sn,,, La(OH),, Ni, and Ni(OH), in varying
degrees. The initial presence of Ni powder in electrodes 'd' - 'g’ makes determining the
source of the Ni and Ni(OH), phases uncertain.

To help determine the origin of the corrosion products in cycled electrodes, cells 'a' -
'c' were made without Ni powder in the mixture. These cells were then cycled under
identical conditions to different levels of capacity degradation (Figure VII-12). When
linear polarization and impedance spectroscopy were performed on new and cycled cells,
it was found that the addition of Ni powder to the anode mixture did not affect the cell
kinetics at these different stages of cycling. After cycling, each electrode was removed
from its cell and analyzed by XRD, shown in Figure VII-13. The average intensity of
the La(OH), (110), (101), (200), and (201) peaks, each normalized to the LaNi, (111)
peak intensity, was plotted against the capacity degradation of each cell. This analysis
was also performed for the Ni(OH), (100) and (101) peaks. These graphs, shown in
Figure VII-14, indicate that there is a linear increase in both the La(OH), and Ni(OH),

phases with the loss of capacity in these cells.

246

2. DISCUSSION

"The decrease in charge transfer resistance as measured by linear polarization and by
impedance spectroscopy can be explained by an increase in the surface area of the alloys.
As hydrogen is electrochemically cycled in and out of each alloy, the discontinuous
~ volume change associated with the o-/B-phase transformation causes the alloy to
decrepitate. The alloy breaks into smaller powders (~1pm) and its surface fractures. The
. increase in surface area increases the number of reaction sites at which charge transfer
can occur. This increase in surface area is supported by the increase in line width of the
MH diffraction peaks. The broadening results from the small size of coherent scattering
domains, caused by small grain size and microstrain in the alloys. The smaller grain size
would also mean a shorter diffusion length for hydrogen that was entering and leaving the
alloy, corresponding to increased electrochemical kinetics. The presence of Ni metal
from the MH corrosion will also create a catalytic effect on water reduction at the alloy
particle surface.

The XRD results show a strong correlation between capacity degradation and phase
fractions. The fact that all the electrodes yielded similar results despite the differing
cycling rates used implies that the phases identified are the products created during

electrochemical cycling.

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247

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Cycle Number

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249

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C. Hydrogen Absorption Capacity

This section will revisit results from isotherm measurements of gas-phase capacities

and compare them to capacities measured in electrochemical cells.

1. Literature Survey

It is well known that alloying of LaNi, reduces the hydrogen absorption capacity. It is
also known that the measured electrochemical capacity of a MH electrode depends on the
discharge current density used.*”

It did not prove possible to use binary LaNi, in Ni-MH cells because its high plateau
pressure (P=~2.0 atm.) meant that the alloy was not easily charged and did not retain
charge well in the slightly pressurized cells we used to make electrochemical
measurements. Two methods to charge binary LaNi, were found in the literature.
Percheron, et al. were somewhat vague in explaining how they achieved capacities of 320
mAh/g.° They do describe activating their alloys with four hydrogen absorption-
desorption cycles before electrode construction and write, “This operation is feasible at
temperatures below 25-°C or with electrodes which have been treated to decrease the
velocity of recombination of hydrogen at the interface.”” Sakai, et al. did not hydrogen
activate their crushed samples. Instead, they achieve electrochemical capacity by cycling
of a LaNi, anode at -20° C, where the plateau pressure is below 1 atm. After
electrochemical activation at -20° C, the metal hydride will charge at room temperature,
but retains a very high rate of self-discharge.’ “When the alloys have equilibrium
pressures in the range of 10-1 atm, the corresponding electrodes were activated easily at
20° C.”” Tests by van Rijswick confirm that alloys used in pressurized cells need no

special treatment to reach close to theoretical capacities.”

255

2. Caltech Results

Figure VII-15 compares the reversible hydrogen absorption capacities of the ternary
substituted alloys of LaNi, as measured in gas-phase reactions to those measured in
electrochemical cells. The electrochemical capacity of the binary alloy is particularly
low. During its charging, significant hydrogen evolution is observed to occur on its
surface, which seems to be favored over hydrogen absorption.

Figure VII-15 shows that the maximum capacities as measured in slightly
pressurized prismatic cells are slightly lower than the corresponding gas-phase. The
capacity suppression at low solute compositions is notable, and can be attributed to the
high plateau (absorption) pressures of the anode materials, which prevents their complete
utilization in our electrochemical test cells. The charge potential is higher for alloys with
higher plateau pressures, and under these conditions the competing hydrogen evolution
reaction is favored over hydrogen absorption by the metal hydride, and hydrogen gas is
formed at the electrode surface. At higher solute compositions, where plateau pressure is
not a problem, the charge process becomes efficient and the difference between the
electrochemical capacities and gas-phase capacities is approximately constant. We
believe this residual discrepancy can be attributed to several sources. Our electrodes are
hot pressed at 300° C in air. During this step a surface oxidation layer will be formed
from active material, thus decreasing the material’s maximum capacity.’ In addition, the
stronger surface (hydroxide) films formed in the electrochemical environment may
impose larger polarization losses, both in the charge transfer as well as the diffusion
processes, especially at the high discharge rates and low electrode dispersion in our tests.
The formation of such films is likely to be aided by the high current densities and low
discharge cut-off potentials used in cycling. All the above factors can result in an
incomplete utilization of the metal hydride material, making the material’s measured

capacities lower than would be obtained under quasi-equilibrium conditions.”

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257

Of the Ni-substituted LaNi, alloys tested here, Sn compositions of 0.2 < x < 0.3

~provide the highest maximum electrochemical capacities of approximately 325 mAh/g.

D. Cyclic Lifetime

1. Literature Survey
See § IL.F.3.a.

The capacity vs cycle number curve of MH electrodes has been modeled by several

authors.*”*“*

As mentioned above (§VII.C), the maximum capacity of an alloy is never
achieved in the first cycle. The capacity increases during the first few cycles, while the
alloy undergoes an activation process usually attributed to decrepitation. The increase in
surface and decrease in the path length for hydrogen diffusion that takes place during
activation decreases the over-potential that causes the electrode to reach its cut-off
potential before all its hydrogen is discharged. After activation is complete, the capacity
begins to decay. This capacity loss is overwhelmingly dominated by the oxidation of
active MH alloy, but has also been attributed to the electrical isolation of particles and an
over-potential resulting from the build-up of corrosion products. The decay resulting
from oxidation of active material is exponential in nature implying that degradation takes
place as a first-order reaction with the amount of degraded material proportional to the
amount of material involved in the charge transfer reaction. The capacity vs cycle
number curve was first modeled by Willems, et al. as decaying exponentially with the
number of cycles after the alloy had experienced activation.” Later studies confirmed the
exponential decay and attempted to include alloy activation * and discharge current
density effects “ in their models. For some long lived alloys (e.g., Si,,), researchers have
postulated a reduction process by which corroded material is slowly converted back to
LaNi,, giving the alloy a constant capacity after several hundred cycles.” The slow

capacity decay of state-of-the-art alloys has prompted some authors to characterize their

258

decay by a linear function.” This follows from a Taylor’s expansion of the exponential
decay:
: -§n
C(n) = limg_. C, e°” = C, (1 -6n) [V1.3]
5 =capacity decay parameter, n=cycle number,

C, = hypothetical initial capacity

2. Caltech Results

It may be noted that the alloys shown here all experience a shorter cyclic lifetime than

16,17

those achieved by other experimenters,’ and this can be attributed to the strenuous
cycling conditions used to test our cells. The deep discharge to a cut-off potential of -0.5
V vs Hg/HgO accelerates the loss in capacity,’ compared to the -0.7 V often used.'*"” The
high applied current density and the low dispersion of active material used in the
electrode powder mixture mean that the overpotentials experienced at a particle surface
will be high, resulting in faster corrosion locally and earlier cut-off during discharge. In
addition, MH alloy powder particles that do not activate during the first few cycles are
more likely to become passivated by electrical isolation from the current collector.
Several cells were made using each alloy, and only the cells with the best lifetimes were
used in subsequent analysis. It was assumed that the vagaries of battery assembly, rather
than the properties of the anode alloy, were responsible for the variability in lifetime
measurements.

To characterize the lifetimes of LaNi,,M, alloys during electrochemical cycling, I will
use the exponential decay parameter 6, fit to capacity-cycle curves of alloys after
activation. This should be a better characterization of the alloy lifetime than a linear
decay. This model is easily applicable to our alloys because the gas-phase activation and
thermal cycling of all but the Sn, alloys insured fast electrochemical activation. In
addition, the fast discharge rate and low cut-off potential used insures that the capacities

of even stable alloys such as MmB, will decay exponentially.

259

a. LaNi,,Sn,
The cyclic lifetime of cells containing LaNi, Sn, metal hydride alloys are presented in
Figure VII-16. The Sn, alloys were the first family of alloys to be melted, and our
- material preparation technique was unfortunately not yet perfected when they were
cycled. These alloys were activated with a single gas-phase absorption. Evidence of this
can be seen in the particularly slow activation of the Sn,, and Sn,, alloys. An additional
cell was made with the Sn,, alloy, denoted by “Sn,,a” in the figure, that had been
activated by 5 thermally driven gas-phase absorption-desorption cycles, but this alloy was
subsequently used in the temperature study (§VII.A.4) instead of the lifetime study.

Each cell was subjected to 200 charge-discharge cycles. The capacity retention was
found to improve with increasing Sn composition. After 100 charge-discharge cycles,
alloys with x = 0.25 or 0.3 exhibit capacities in excess of 200 mAh/g, an impressive
number when compared to the state of art MmNi, ,Co,,Mn,,Al,, MH alloys evaluated at
JPL” (Mm in Figure VII-16). The cyclic lifetime in cells with x = 0.4 exceeds that of

cells with 0.2 < x < 0.3, as evidenced by the capacity curves after 150 cycles.

The cyclic lifetimes obtained from the LaNi, ,M,, alloys M € {Si, Ge, Al, Sn, In} are
shown in Figure VII-17 with that of LaNi,. These alloys were activated with one gas-
phase hydrogen absorption before electrochemical cycling, and the anode mixture did not
include Ni powder. The Sn alloy has the best initial capacity and cyclic lifetime of all the
materials in this test. The binary, Si, Ge, and Al alloys have plateau pressures above 1
atm., and because our cells are only slightly pressurized, the electrochemical capacities of
these materials could not be measured accurately. The fall in capacity experienced by the
LaNi, anode after cycle #4 implies that the cell’s o-ring seal failed after this cycle,
limiting the cell pressure to 1 atm. During subsequent cycles, it is likely that the LaNi,

alloy did not experience the discontinuous volume expansion associated with the a.- to B-

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262

phase transition. This cell maintained an extremely stable capacity corresponding to
hydrogen cycled in the a-phase regime. The seal on the Al,, cell failed during the first
cycle, but the subsequent capacity corresponds to ~1.1 H/f.u., which is greater than the
solubility limit in the o-phase (~0.3 H/f.u.). The slight increase in capacity of the Al,,
cell could result from continued activation.

The Si,, and Ge,, alloys also have plateau pressures greater than 1 atm, but the o-ring
seals of these cells seem to have been sufficient to contain the increase in pressure. The
Si alloy, which has been shown to perform relatively well in tests by Meli, et al.,”
experienced some passivation of the electrode above cycle #15 which is removed later in
the cycling. The Ge,,, In,,, and Si,, alloys have similar capacities after cycle #80.

The cyclic lifetimes of these alloys were characterized with linear decay constants.
The linear decay constants for these alloys are shown in Table VII-1. The relative
stability of these alloys (In > Si > Ge > Sn) during electrochemical cycling is somewhat
different than the relative stabilities determined from further testing of these solutes. The
addition of Ni (or Cu) powder to the anode mixture has been shown to increase the
lifetimes of MH alloys during cycling.” Although these cells did not have Ni powder
present, it is possible that the production of Ni during the corrosion of LaNi, (equation
111.6) in cells with faster initial corrosion stabilized the cells in subsequent cycling.
Because it is not known how the lack of Ni powder or adequate activation affect cyclic

lifetimes, these results are not included in the final analysis.

c. LaNi, Ge,

Figure VII-18 shows the capacities of the cells containing LaNi,,Ge, during 500
electrochemical cycles. By the time these cells were tested, our sample preparation
procedures were mature, as evidenced by the nice exponential decay observed in these
alloys. The capacity loss rate of the LaNi, Ge, alloys is rather unusual. All of the Ge,

alloys experience a rapid loss of 5 to 20% of the capacities in the first ~20 cycles,

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265

followed by an exponential decay. The mechanism underlying this rapid initial fall is not
clear. Equivalent cycle life studies completed with Sn-substituted alloys did not show
this affect, sO it should not result from the battery fabrication procedure. Even with this
rapid initial fall, the capacity retention of the LaNi, Ge, alloys during long term cycling is
impressive compared to the binary or Sn-substituted alloys. After 400 deep (100%)
discharge cycles, the capacities are still in the range of 125 and 150 mAh/g for alloys
with Ge compositions of 0.4 and 0.5, respectively. The capacity decay parameters of the
Ge,, and Ge,, alloys are almost comparable to that of the commercial
MmNi, ,Co,,Mn,,Al,,, although the capacity in the initial stages is lower. The cycling
stability of the Ge,,, alloy is spectacular, and its capacity surpasses that of the MmB,
alloy at cycle #350. We believe that the increased performance of this alloy over the Ge,,

alloy results from the second phase present in Ge, ,.

d. LaNi,,.M.

Figure VII-19 shows the cyclic lifetimes of LaNi, .M, alloys during electrochemical
cycling. Many of the cells experience the quick capacity loss after activation seen with
the Ge, alloys. Each cell was cycled ~400 times, until the cell capacity was
approximately 50 mAh/g, or until the cell experienced catastrophic failure. Both of the
cells containing the LaNi,,,[n,,, alloy cracked and burst. This could result either from a
pressure build-up during cycling or from the volume expansion associated with hydride

formation and the phase changes that accompany corrosion. The behavior of the Si,,

alloy is rather odd, and could be related to the quick capacity loss after activation.

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267

3. Discussion

The goal of this thesis work is to determine the chemical effects of Ni-substitution on
the reversible hydriding properties of LaNi,, in particular its lifetime during
electrochemical charge-discharge cycling. To separate the chemical from mechanical and
microstructure effects, it was sometimes necessary to exclude data from the final analysis.
For example, the alloy Ge,, was not single phase. This would not have presented a
problem in itself, but the small precipitates created effects such as smaller particle size
and possibly increased lattice strain in the alloys, which would increase the surface area
of the alloy and impact the cycling stability. As seen in Fig. 1, the Ge,, alloy is less
stable than the Ge, ,, alloy during electrochemical cycling.

As mentioned above, several alloys (Si,,, Al,,, Ge,,) had plateau pressures greater
than | atm, and this prevented accurate measurement of their capacities in the slightly
pressurized cells used in these tests. An example of this can be seen in the Ge,, curve in
Figure VII-18. This alloy could not retain a charge higher than 30 mAh/g after the 20th
cycle, when the cell O-ring seal was broken. Some solutes (Ge, Si) have large solubilities
in the KOH electrolyte, thereby making the composition of the resultant alloys change
during their cycling. It is possible that the faster decay of the Ge,, trace of Figure VII-18
is enhanced by a greater dissolution of Ge from the smaller alloy particles.

The cyclic lifetimes of the Sn,, Ge,, and M, alloys were characterized with the
exponential decay parameter 5 of equation VII.3. To explore the relationship between the
alloy capacity fade and alloy composition, solute chemistry, and alloy physical
properties, 5 was plotted vs a number of variables characterizing the alloy. Figure
VII-20 shows 6 plotted vs solute composition x. We see that for the Sn, and Ge, alloys, 6
has an almost linear dependence on composition, but the slopes of these correlations are
different. Furthermore, the capacity fade parameter is lower for Ge-substituted alloys

than for alloys with the same amount of Ni replaced by Sn. The characteristic

268

degradation constant for a state-of-the-art commercial MmB, alloy from Rhone-Poulenc,
Inc., is ~0.0022, which could be achieved with solute compositions of x,, = 0.36 or x,, =
0.5.

The discontinuous volume expansion of the Ge, alloys, seen in Figure V-26,
increases only slightly for x > 0.2, but the capacity fade parameter continues to decrease
linearly with composition. This implies that there are additional lifetime controlling
effects present. Sn-substituted alloys are not included in the volume expansion data, but
a plot of 5 vs AV (Figure VII-21) supports this conclusion. These data seem rather
scattered and there is not a clear correlation between the capacity fade parameter and
alloy volume expansion.

Finally, Figure VII-22 displays 5 vs the average bond strength (heat of formation,
AH) of La with the B (= Ni,,.M,,) elements. As seen in Figure VII-22b, the relationship
between 5 and AH,,, is linear for almost all of the M, alloys. The alloys with faster
capacity fade than would be expected from their AH,,, have been labeled in Figures
VII-21 and VII-22 for comparison.

Although it is true that MH lifetimes during charge/discharge cycling of alkaline
rechargeable cells are affected by the alloy volume expansion upon hydrogen absorption,
this alloy property is not adequate in predicting the exponential capacity fade of the
alloys. Another important factor is the resistance of the lattice to the formation of
microstructural defects, such as can be measured by diffraction peak broadening. There
are two important aspects of the crystalline defects that form after hydrogen activation.
First, the defects are evidence that the lattice is unstable and is beginning to decompose.
Second, the presence of lattice defects enhances the diffusion of metal atoms necessary
for phase transformation from the haucke phase structure. The bond strength of La to an
average B ( = Ni,,,,.M,,) element serves as a more accurate predictor of the alloy capacity

fade.

* 0.016

0.014
0.012
0.010

= 0.004
= 0.002
= 0.

S 0.01

OQ.

0.001

MON

269

BAAAS AAAS RAAEA RAS GSDRE LARC SORRELL:

Loobiibitibiiel

a | TTT] ae

pisrtt Lil

Potty ptt | PE

Lett

1 0.2 0.3 0.4 0.5

Solute Composition [x in LaNi,_,.M, |

Figure VII-20 Linear and semi-log plots of 6 vs substituted

composition, x. { @ -Ge,; O -Sn,; @ -M,}

Capacity fade parameter, 6

270

"0.016 FEET TTT TTT TTT TTT TT TTT PTT PTT pT
0.014 F- Slo.42
0.012 —
0.010 E!Mo.17 Gay a4 =
0.008 E- ; S
0.006 — Gey > =
0.004 E- . —
0.002 F-- .
ml RETOUR FHREREREHRGUTOURE CHO GTORRORT ES

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20 18 16 14 12 10

Discontinuous volume expansion, AV/V [%]|

Figure VII-21 Linear and semi-log plots of 5 vs
discontinuous volume expansion of alloy.

{ @ -Ge,;

e -M,}

| 0.016, PITTI PITT TTT TTT TTT pe ITT TTT TE
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0.012 |
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rT! Prerry erry
ee et | teil Ls

0.001

} See O EERO Oe eR Re Se Liiplitsi
16 -18 -20 -22 -24
AH, 3 [kJ / mol atom]

Figure VII-22 Linear and semi-log plots of 5 vs the heat

of formation of La with an average B element.
{ @ -Ge,; O -Sn,; @ -M,}

272

S. Luo, W. Luo, J.D. Clewley, Ted B. Flanagan, and L.A. Wade, J. Alloys Comp., 231
(1995): 467.
K. Petrov, A.A. Rostami, A. Visintin, and S. Srinivasan, J. Electrochem. Soc., 141
(1994): 1747.
C.K. Witham, B. Fultz B.V. Ratnakumar, and R.C. Bowman, Jr., in Hydrogen and
Metal Hydride Batteries, P.D. Bennett and T. Sakai, Ed., PV 94-27, p. 68, The

Electrochem. Soc. Proceedings Series, Pennington, NJ (1994).

A. Ziittel, F. Meli, and L. Schlapbach, J. Alloys Comp., 200 (1993): 157.

Z.-P. Li, W.-Q. Lei, C.-P. Chen, J. Wu, and Qi-Dong Wang, J. Less-Common Met.,
172-174 (1991): 1260.

G. Bronoél, J. Sarradin, M. Bonnemay, A. Percheron, J.C. Achard, and L.
Schlapbach, Int. J. Hydrogen Energy, 1 (1976): 251.

T. Sakai, K. Oguro, H. Miyamura, N. Kuriyama, A. Kato, H. Ishikawa, and C.
Iwakura, J. Less-Common Met., 159 (1990): 193.

M.H.J. van Rijswick, in Hydrides for Energy Storage, eds. A.F, Andresen and A.J.
Maeland (Pergamon, Oxford, 1978), p. 261.

T. Sakai, M. Matsuoka, and C. Iwakura, in Handbook on the Physics and Chemistry
of Rare Earths, K. A. Gschneidner, Jr., and L. Eyring, eds., Vol. 21, Elsevier Science
B. V., Amsterdam (1995), p. 133.

J.J.G. Willems, Philips J. of Research, 39 (1) (1984): 1.

S. Wakao and Y. Yonemura, J. Less-Common Met., 131 (1987): 311.

H. Sawa, M. Ohta, H. Jakao, and S. Wakao, Z. Phys. Chem. NF, 164 (1989): 1527.
Z.P. Li, Y.Q. Lei, B.H. Liu, J. Wu, and Q.D. Wang, Phys. Chem., 183 (1994): 287.
Y.Q. Lei, C.S. Wang, X.G. Yang, H. G. Pan, J. Wu, and Q.D. Wang, J. Alloys Comp.,
231 (1994): 611.

M.P. Sridhar Kumar, K. Petrov. W. Zhang, A.A. Rostami, S. Srinivasan, G.D. Adzic,
J.R. Hounson, J.J. Reilly, and H.S. Lim, J. Electrochem. Soc., 142 (1995): 3424.

273

A. Anani, A. Visintin, K. Petrov, and S. Srinivasan, J. Power Sources, 47 (1994);
261. |

M. Wasz, R.B. Schwarz, S. Srinivasan, and M.P.S. Kumar, in Materials for
Electrochemical Energy Storage and Conversion-Batteries, Capacitors, and Fuel
Ceils, Symposium Proceedings of the Materials Research Society, MRS, Pittsburgh,
PA (1996).

F. Meli, A. Zuettel, and A. Schlapbach, J. Alloys Comp., 1992, 190, 17.

H. Ishikawa, K. Oguro, A. Kato, H. Suzuki, and E. Ishii, J. Less-Common Met., 107
(1985): 105; 120 (1986): 123.

T. Sakai, H. Ishikawa, K. Oguro, C. Iwakura, and H. Yoneyama, J. Electrochem. Soc.,
134 (3) (1987): 558.

A.H. Boonstra and T.N.M. Bernards, J. Less-Common Met., 161 (1990): 355.

274

Vill. Conclusions

By characterizing solutes from a relatively small region of the periodic table, I was
able to find some trends in the physical properties of LaNi, M, alloys and the substituting
element. The solubilities of sp-metals for Ni in LaNi, is strongly dependent upon the
valence state of the substituting element. The expansion of the LaNi, unit cell by solute
substitution is governed by the solute’s metallic radius. The total lattice expansion
associated with the metal-to-hydride phase transition is suppressed by solute substitution,
and the magnitude of that suppression is related to the heat of formation of the solute with
La. There seems to be a critical substituted composition above which substituting Ge for
Ni does not reduce the discrete volume expansion experienced upon hydride formation.
In all cases, the discrete c-axis expansion is more suppressed than the discrete a-axis
expansion. The total lattice expansion of the alloys is more isotropic.

The logarithmic relationship between the alloy plateau pressure and the unit cell
volume of the metal is confirmed for Sn- and Ge-substituted alloys. However, Ge-
substituted alloys have lower plateau pressures than Sn-substituted alloys with the same
unit cell volume. This implies that there is a second order effect controlling the alloy
plateau pressures.

The electrode charge transfer kinetics measured by DC micropolarization and AC
impedance give similar results. The conflicting trends seen in G-substituted and Sn-
substituted alloys make it difficult to ascribe changes in charge transfer resistance to any
single phenomenon. A correlation was seen between surface area measured by XRD and
the charge transfer resistance.

There is a definite trend between alloy stability during electrochemical cycling and

the heat of formation of the solute element with La.

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