NASA Conference Publication 3228

Space

Electrochemical

Research and

Technology

N94-23345

--THRU--

N94-23367

Unclas

(NASA—CP-3223) SPACE

ELECTROCHEMICAL RESEARCH AND

TECHNOLOGY (NASA) 249 p

y31viv H1/44 0200510

Proceedings of a conference held at

NASA Lewis Research Center

Cleveland, Ohio

April 14-15, 1993

NASA
NASA Conference Publication 3228

Space

Electrochemical

Research and

Technology

Proceedings of a conference held at

NASA Lewis Research Center

Cleveland, Ohio

April 14-15, 1993

National Aeronautics and

NASA

Space Administration

Office of Management

Scientific and Technical

Information Program

1993
PREFACE

This document contains the proceedings of NASA's fourth Space Electrochemical Research

and Technology (SERT) Conference, held at the NASA Lewis Research Center on April 14

and 15, 1993. The objective of the conference was to assess the present status and general

thrust of research and development in those areas of electrochemical technology required to

enable NASA missions into the next century. The conference provided a forum for the

exchange of ideas and opinions of those actively involved in the field, in order to define new

opportunities for the application of electrochemical processes in future NASA missions. Papers

were presented in three technical areas: advanced secondary batteries, fuel cells, and advanced

concepts for space power.

iii

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SPACE ELECTROCHEMICAL RESEARCH AND TECHNOLOGY

April 14-15, 1993

Contents

Section One

Advanced Secondary Batteries

Aerospace Applications of Batteries

Shahid Habib, NASA Headquarters ................................................ 3

Applications of X-ray Absorption Fine Structure to the IN SITU Study of the Effect of Cobalt

in Nickel Hydrous Oxide Electrodes for Fuel Cells and Rechargeable Batteries

Sunghyun Kim, Donald A. Tryk, and Daniel A. Scherson, Case Western Reserve University,

and Mark R. Antonio, Argonne National Laboratory .................................... 13

Review of Super Ni/Cd Cell Designs and Performance

Bruce Abram s-Blakemore, Eagle-Picher, Inc . ........................................ 21

Thermodynamics of Nickel-Cadmium and Nickel-Hydrogen Batteries

Digby D. Macdonald and Mark L. Challingsworth, The Pennsylvania State University ............ 27

High Specific Energy, High Capacity Nickel-Hydrogen Cell Design

James R. Wheeler, Eagle Picher Industries ........................................... 41

High Surface Area, Low Weight Composite Nickel Fiber Electrodes

Bradley A. Johnson, Richard E. Ferro, Greg M. Swain, and Bruce J. Tatarchuk, Auburn

University................................................................. 49

Ovonic Nickel Metal Hydride Batteries for Space Applications

S. Venkatesan, D.A. Corrigan, M.A. Fetcenko, P.R. Gifford, S.K. Dhar, and S.R. Ovshinsky,

Ovonic Battery Company ....................................................... 61

Alloys for Hydrogen Storage in Nickel/Hydrogen and Nickel/Metal Hydride Batteries

Anaba Anani, Arnaldo Visintin, Konstantin Petrov, and Supramaniam Srinivasan, Texas

A&M University System; James J. Reilly and John R. Johnson, Brookhaven National

Laboratory; Ricardo B. Schwarz and Paul B. Desch, Center for Materials Science ............... 67

Status of the Development of Rechargeable Lithium Cells

G. Halpert, S. Surampudi, D. Shen, C-K Huang, S. Narayanan, E. Vamos, and D. Perrone,

Jet Propulsion Laboratory ...................................................... 85

Impedance Measurements on a Sprial-Wound Nickel/Metal Hydride Cell Cycled in a Simulated

LEO Orbit

Margaret A. Reid, NASA Lewis Research Center ...................................... 97

NASA Lewis Advanced IPV Nickel-Hydrogen Technology

John J. Smithrick and Doris L. Britton, NASA Lewis Research Center ....................... 109

v

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Section Two

Fuel Cells

Primary and Secondary Electrical Space Power Based on Advanced PEM Systems

N.E. Vanderborgh, J.C. Hedstrom, and K.R. Stroh, Los Alamos National Laboratory, and

J.R. Huff, Ballard Power Corporation .............................................. 129

PEM Regenerative Fuel Cells

Larry L. Swette, Anthony B. LaConti, and Stephen A. McCatty, Giner, Inc . ................... 139

High Power Density Proton Exchange Membrane Fuel Cells

Oliver J. Murphy, G. Duncan Hitchens, and David J. Manko, Lynntech, Inc .................... 149

The Electrolyte Challenge for a Direct Methanol-Air Polymer Electrolyte Fuel Cell Operating

at Temperatures Up to 200 °C

Robert Savinell, Ernest Yeager, Donald Tryk, Uziel Landau, Jesse Wainright, Dominic

Gervasio, Boris Cahan, Morton Litt, Charles Rogers, and Daniel Scherson, Case Western

Reserve University ...........................................................167

Advances in Direct Oxidation Methanol Fuel Cells

S. Surampudi, S.R. Narayanan, E. Vamos, H. Frank, and G. Halpert, Jet Propulsion Laboratory;

A. LaConti and J. Kosek, Giner, Inc.; G.K. Surya Prakash and G.A. Olah, University of Southern

California..................................................................181

Recent Advances in SPED Water Electrolyzer

James F. McElroy, United Technologies Corporation ................................... 193

Section Three

Advanced Concepts for Space Power

High Energy Density Aluminum-Oxygen Cell

E.J. Rudd and D.W. Gibbons, ELTECH Research Corporation ............................. 205

Composite Fiber Structures for Catalysts and Electrodes

Christopher J. Marrion, Donald R. Cahela, Soonho Alin, and Bruce J. Tatarchuk, Auburn

University.................................................................219

Limiting Factors for Carbon Based Chemical Double Layer Capacitors

M. Frank Rose, C. Johnson, T. Owens, and B. Stevens, Space Power Institute .................. 231

A High Performance H2-C12 Fuel Cell for Space Power Applications

Everett B. Anderson, E. Jennings Taylor, Gerald Wilemski, and Alan Gelb, PSI Technology

Company..................................................................243

Hairy Carbon Electrodes Studied by Cyclic Voltammetry and Battery Discharge Testing

Deborah D.L. Chung and Xiaoping Shui, State University of New York at Buffalo, and

Christine A. Frysz, Wilson Greatbatch, Ltd ........................................... 251

vi
Section

One

Advanced Secondary

Batteries
y3goi ► PyS1^

AEROSPACE APPLICATIONS OF BATTERIES

Shahid Habib

NASA Headquarters

Washington, DC

ABSTRACT

NASA has developed battery technology to meet the demanding requirements for aerospace

applications; specifically, the space vacuum, launch loads, and high duty cycles. Because of

unique requirements and operating environments associated with space applications, NASA has

written its own standards and specifications for batteries.

INTRODUCTION

Since 1958, over 300 NASA spacecraft have used primary or secondary batteries. Primary

batteries provide energy to critical systems on launch vehicles and portable equipment used by

humans in space. Secondary battery systems supply power, usually 28 V DC, when spacecraft are

in darkness. When a spacecraft is in sunlight, photovoltaic cells power the satellite and darkness.

When a spacecraft is in sunlight, photovoltaic cells power the satellite and charge the secondary

battery.

The orbit determines the duty cycle for a secondary battery. Low Earth orbit applications

are the most demanding because the spacecraft passes through the Earth's shadow on almost every

orbit. Since a typical orbit period is 95 min, the battery must be recharged rapidly. Satellites in

higher orbits need batteries less frequently--these orbits have longer periods and are more often

totally in sunlight. A typical spacecraft has a design lifetime of three years. (Table 1)

NASA has used nickel-cadmium batteries for over 160 satellites to low Earth orbit,

geostationary orbit, and the nearby planets. Some recent NASA applications are Topex/Poseidon,

the Tracking and Data Relay Satellite, and Mars Observer. The Hubble Space Telescope uses

nickel-hydrogen batteries, and the Extravehicular Mobility Unit used on Shuttle missions has a

silver-zinc primary battery. Descriptions of these applications follow. NASA also uses many

other battery chemistries.

TOPEX/POSEIDON SPACECRAFT

To increase our understanding of global ocean circulation, the Topex/Poseidon spacecraft

(Figure 1) carries two radar altimeters for making very precise measurements of sea level. The

French space agency, Centre National d'Etudes Spatiales (CNES), and NASA built the altimeters.

The designs permit correction of measurement errors induced by the atmosphere and ionosphere.

Uncertainty in the orbital position limits the accuracy to about 14cm, even though the satellite is at a

high altitude for low drag and uses both laser and microwave Doppler tracking methods. On

August 10, 1992, an Ariane rocket launched Topex/Poseidon from Kourou, French Guiana into a

66 degree inclination orbit.

Topex/Poseidon uses NASA's Multimission Modular Spacecraft (MMS), a design first

used in 1980 for the Solar Maximum Mission. Because the Instrument Module always faces the

Earth, the single panel of solar cells that provides at least 2260 W of power has gimbals to track the

PWG"NG PA _£ Rt_ANK NOT Ffl-ME i
sun. The Modular Power System (MPS) contains the power conditioning electronics and three

batteries with 22 NASA 50 Ali cells (Figure 2). Topex/Poseidon is the sixth mission to use

this battery; the first was Landsat 4 in 1982. NASA initiated the design effort for this battery in

1975. From a total production of over 1400 cells, industry has made 23 flight batteries plus 20 test

or spare units. References 1 and 2 define the design. The only qualified cell supplier is Gates

Aerospace Batteries.

TRACKING AND DATA RELAY SATELLITE

From its geostationary orbit the Tracking and Data Relay Satellite (TDRS) links NASA

controllers on the ground with satellites in low Earth orbits (Figure 3). For example, TDRS

provides voice, television, and data communications to and from the Space Shuttle. Two TDRS

satellites are stationed over the Atlantic and the other two over the Pacific. Except for a small zone

over the Indian Ocean, TDRS allows continuous communications between a satellite and the

ground. A ground station at White Sands, New Mexico controls all four satellites. TDRS receives

high rate data at up to 300 Mbps through two large K-band antennas dedicated to (pointed at)

specific satellites. There is also a low rate S-band service that can be shared by up to 20 user

satellites.

Each TDRS satellite has two solar arrays supplying 1700 W while tracking the sun. Three

Ni-Cd batteries can supply 1400 W while TDRS is in darkness for a 1.3 hr period per day. The

batteries are reconditioned by completely discharging the cells just before the two 45-day periods

per year when they are needed.

MARS OBSERVER

Mars Observer began its trip to Mars in October 1992 and will arrive in August 1993

(Figure 4). After a four month period for checkout and orbit adjustments, the spacecraft will map

the surface and profile the atmosphere using a complementary instrument set covering spectral

regions from radio waves through gamma rays. The mapping phase will continue for at least a full

Mars year or 22 months. The 117 minute mapping orbit is nearly circular and polar. Darkness

periods vary seasonally between 36 and 42 minutes.

The solar array power is 1130 W when initially deployed and 930 W when fully deployed

at Mars. The batteries provide approximately 500 W for the spacecraft during launch, maneuvers,

orbit insertion, and occultations. There are two 42 Ali Ni-Cd batteries; the cells are similar to the

NASA 50 Ali cell. Because there are only 17 cells per battery, an electronic circuit produces the

desired 28 V output to the spacecraft. The backups were 37 Ah Ni-Cd batteries with an advanced

design.

HUBBLE SPACE TELESCOPE

The Hubble Space Telescope (HST) (Figure 5), launched three years ago, makes

observations at ultraviolet wavelengths and with higher angular resolution than any telescope that

must look through the Earth's atmosphere. HST's major accomplishments include observations of

the early stages of star formation, and the development of a major storm in Saturn's atmosphere.

HST has a very accurate attitude control system that points the telescope toward a star, galaxy, or

planet while the articulated solar arrays follow the sun. After a Shuttle servicing mission in

4
December 1993, HST will be more productive and reliable with its new solar arrays and

gyroscopes, a new Wide Field and Planetary Camera with corrective optics for improved

resolution, and separate corrective optics for the other instruments.

To meet a 15-year lifetime requirement, the original HST design had Ni-Cd batteries in

units that could be replaced on orbit. Given problems with Ni-Cd batteries and a 4-year launch

delay for other reasons, NASA decided to substitute Ni-H 2 batteries (Figure 6). Each of HST's 6

parallel batteries has 22 cylindrical cells. Actual capacity is 75 to 80 Ah per battery, and in normal

operation the depth of discharge is less than 10 percent. NASA plans to use Ni-H 2 batteries on

other large spacecraft.

EXTRAVEHICULAR MOBILITY UNIT

Astronauts performing space walks wear a special suit known as an Extravehicular Mobility Unit

(EMU). This system provides essential life support--a pure oxygen environment at a differential

pressure of 4.2 pounds per square inch, water vapor and carbon dioxide removal, and temperature

control while the astronaut is outside the Shuttle (Figure 7). The major components of the EMU

are the hard upper torso with the life support systems, lower torso (legs), arms, helmet, and liquid

cooling and ventilation garment. Each component comes in several sizes that are custom fitted to

astronauts. The EMU, which can be reused during a Shuttle mission, is designed for a 15-year

lifetime.

A silver zinc battery provides power for monitoring and controlling the life support

systems, circulating fluids, and communicating with Shuttle. The EMU has a 16.5 V (11 cell)

battery rated at 27 Ali. Each EMU has a new battery that is filled and charged just before launch.

The battery capacity allows continuous operation of the EMU for up to 7 hours and can be

recharged in orbit with Shuttle power. A recent series of tests (5) allowed the wet life of these

batteries to be extended from 135 to 166 days or about 3 Shuttle missions.

AEROSPACE FLIGHT BATTERY SYSTEMS PROGRAM

The origin of the NASA Aerospace Flight Battery Systems Program was the

Administrator's concern over battery-related launch delays and even the shorting of a battery prior

to a launch. NASA formed a Steering Committee in 1985 to review battery requirements for future

missions and to investigate ways to increase the safety, reliability, and performance of space power

systems. The Steering Committee participated in the development of a program plan for a NASAwide effort with objectives, tasks, a budget, and a schedule.

The program objectives are to enhance cell/battery safety and reliability, maintain current

battery technology, increase fundamental understanding of primary and secondary cells, support

development of advanced technology for flight use, reduce battery technology related risk, and

ensure that safe and reliable batteries are available for NASA's future missions. The technical

approach employed in the program establishes technical guidelines for procurement, design,

production, testing, and operation; opens and maintains communications lines among NASA field

installations and the aerospace battery community; qualifies advanced technologies; and implements

independent cell verification. The program recently redefined some tasks in response to continuing

problems with the production of Ni-Cd cells.

The program's sponsor is the Engineering Division in the Office of Safety and Mission
Quality at NASA Headquarters. NASA's Lewis Research Center (LeRC) has overall program

management responsibility. Specific projects (Figure 8) are assigned to various NASA field

installations. The program has tasks addressing battery systems technology, secondary batteries,

primary batteries, and program management. Reference 4 has a general history and overview of

this program.

The Battery Systems Technology Task addresses handbooks, training, a data base, and the

NASA Battery Workshops. Under this task, Goddard Space Flight Center (GSFC) is preparing a

Handbook for Nickel-Hydrogen Batteries and a Handbook for the Handling and Storage of NickelCadmium Batteries that will become the basis for the training effort by Kennedy Space Center

(KSC). LeRC is developing the data base that will compile knowledge gained from

manufacturing, ground testing, and flight experience. This effort includes a computer bulletin

board for problem reporting, technical notes, test reports, and policy documents. The Marshall

Space Flight Center (MSFC) conducts the NASA Aerospace Battery Workshop, (see NASA

Conference Publications 3119 and 3140).

To provide independent checks and balances, NASA invested in expanded and improved

battery test facilities at GSFC, the Jet Propulsion Laboratory (JPL), and the Naval Weapons

Support Center in Crane, Indiana. Crane currently places its test data into a computer data base.

In addition, LeRC is developing a method for non-destructive evaluation by impedance

spectroscopy of secondary cells, and MSFC is establishing procedures for destructive physical

analysis.

The Secondary Battery Technology Task includes both Ni-Cd and Ni-H 2 battery activities.

The Ni-Cd efforts specifically address the cell quality and reliability problems that surfaced in the

late 1980's. In the near term, GSFC will convert the existing NASA Ni-Cd specification,

Reference 5, into a guidelines document and where necessary strengthen the technical and quality

assurance requirements. Gates will manufacture cells with modified plates and separators for a

statistically designed experiment that evaluates nickel attack level, positive plate loading, and

negative plate loading. GSFC and Crane began testing cells developed by Hughes and SAFT

about two years ago. Additional cells from SAFT and Sanyo will enter testing at Crane in 1993.

The far-term solution is advanced Ni-Cd cells with electrochemically impregnated plates

and non-nylon separators (similar to the Topex backups and the Hughes cells under test at Crane)

or Ni-H2 cells. LeRC leads the effort for Ni-H2 evaluation that demonstrated the advantages of a

26 percent KOH electrolyte concentration and a catalyzed wall wick for aerospace applications.

GSFC will prepare a guideline document for Ni-H2 cells.

JPL leads an effort to develop computer models of the Ni-Cd and N1-H 2 electrochemical

systems. The initial Ni-Cd model, available now, uses the table look-up approach to predict

battery performance, specifically voltage and efficiency vs. temperature and state-of-charge.

Predictions made by the second version, with a one dimensional electrochemical model, generally

match test data, although the model needs some modifications and additional verifications. The

goal is models that can be used to develop accelerated test methods for determining battery quality

and reliability.

The objective of the Primary Battery Task is improved performance, reliability, and safety,

particularly for cells used on Shuttle missions. As the leader for this task, Johnson Space Center

(JSC) prepared a Primary Battery Design and Safety Handbook. JSC investigated Yardney's

patented thermal short circuit protection method for Lithium K-cells but found the performance

6
degradation unacceptably high. Both JSC and JPL obtained promising test results on some

LiSOC12 cells. JSC supported development on two Zr-0 2 cells with capacities of 9-12 and 150-

200 Ali exceeded performance goals for capacity and specific energy at discharge rates of 1 to

6 A.

SUMMARY

Batteries used in several current space programs are described. NASA's unique

requirements for aerospace batteries have necessitated the writing of their own standards and the

establishment of a Steering Committee for oversight of the requirements, safety and reliability of

these systems.

REFERENCES

1. GSFC 74-150000, Specification for the Manufacturing of Aero-Space Nickel-Cadmium

Storage Cells.

2. NASA 50 Ali Battery Manual, NASA CR 166701.

EMU Ag-Zn Battery Wet-life Extension Test, in The 1991 NASA Aerospace Battery

Conference, NASA Conference Publication 3140, 1992, pp. 249-260.

4. Manzo, M. and O'Donnell, P. NASA Flight Battery Program, An Update, in the 1991

NASA Aerospace Battery Conference, NASA Conference Publication 3140, 1992, pp. 7-

18.

NHB 8073.1, NASA Specification for Manufacturing and Performance Requirements of

NASA Standard Aerospace Nickel-Cadmium Cells.

FIGURES

Fig. 1 Topex/Poseidon Spacecraft

Fig. 2 A 50 Ali cell, and two smaller cells

Fig. 3 Space Shuttle deploys a Tracking and Date Relay Satellite

Fig. 4 Mars Observer spacecraft

Fig. 5 Hubble Space Telescope

Fig. 6 Hubble's Ni-H2 cells

Fig. 7 Extravehicular Mobility Unit

Fig. 8 NASA Aerospace Battery Program tasks and management
Table 1. Typical Aerospace Battery Duty Cycles

Aerospace Minimum Maximum Cycles

Application Charge Discharge per

(orbit) Time Time Year

Low Earth 60 min 35 min 6000

Genstationary 22.7 hr 1.3 hr 90

10
11
NASA NASA

LeRC I I MSFC/ARC/KSC

• PROGRAM • POWER SYSTEMS

SUPPORT & BATTERY

APPLICATIONS

• NI-Hi TECH SUPPORT

• PROBLEM

REPORTING

SYSTEM

NASA HEADQUARTERS

S&MQ OFFICE

PROGRAM MANAGER

BATTERY

--------------------- STEERING

COMMITTEE

NASA LeRC

• PROGRAM OFFICES

•

PROJECT MANAGER

FIELD CENTERS

• FIE DoD

NASA

GSFC

NASA

JSC

JPL

• BATTERY SYSTEMS • PRIMARY BATTERIES • APPLIED NI-CAD

GUIDELINES & GUIDELINES & TECHNOLOGY

TRAINING QUALIFICATIONS BATTERY SAFETY

• EVALUATION & & HANDLING

RESOLUTION OF GUIDELINES

W-CAD PROBLEMS

• VERIFICATION OF

QUALITY &

RELIABILITY OF

SECONDARY CELLS

12
1,59 N94'2

APPLICATIONS OF X-RAY ABSORPTION FINE STRUCTURE TO THE IN SITU

STUDY OF THE EFFECT OF COBALT IN NICKEL HYDROUS OXIDE ELECTRODES

FOR FUEL CELLS AND RECHARGEABLE BATTERIES

Sunghyun Kim, Donald A. Tryk, Daniel A. Scherson

Department of Chemistry

Case Western Reserve University

Cleveland, OH 44106-7078

Mark R. Antonio

Argonne National Laboratory

Chemistry Division

Argonne, IL 60439

ABSTRACT

Electronic and structural aspects of composite nickel-cobalt hydrous oxides have been examined in

alkaline solutions using in situ X-ray absorption fine structure (XAFS). The results obtained have

indicated that cobalt in this material is present as cobaltic ions regardless of the oxidation state of nickel in

the lattice. Furthermore, careful analysis of the Co K-edge Extended X-ray absorption fine structure data

reveals that the co-electrodeposition procedure generates a single phase, mixed metal hydrous oxide, in

which cobaltic ions occupy nickel sites in the Ni0 2 sheet-like layers and not two intermixed phases each

consisting of a single metal hydrous oxide.

INTRODUCTION

Nickel oxide electrodes have found widespread application in a variety of energy storage and

energy generation devices, including rechargeable batteries and fuel cells. ? The addition of certain metal

cations to the nickel hydrous oxide lattice can profoundly affect its electrochemical characteristics. ? In

particular, cobalt has been found to shift the potential associated with the oxidation of Ni(OH) Z to

Ni00H2 and to improve the charge acceptance of nickel oxide electrodes used in alkaline batteries. ? In

contrast, composite nickel oxide films involving iron as the guest metal exhibit high electrocatalytic activity

for oxygen evolution and, as such, may be of practical value in alkaline water electrolysis. 3,4

The studies presented herein have been aimed at elucidating the structural and electronic properties

of composite nickel-cobalt hydrous oxides using in situ X-ray absorption fine structure (XAFS) as a probe

of the host (Ni) and guest (Co) metal sites in the lattice.

Basic aspects of XAFS as applied to the study of electrochemical interfaces have been amply

reviewed in the literatures and will not be discussed here. This technique can provide structural

information without relying on long range order. This factor is of crucial importance, as high charge

storage capacities can only be achieved by using materials in high area form, which are in many cases

amorphous or consist of particles that are too small to achieve sufficient coherence for standard X-ray

diffraction techniques to be very useful. Furthermore, high intensity X-rays in the energy range of interest

can penetrate through thin electrolyte layers and low-Z window materials, such as organic polymers and

thin aluminum sheets. This provides a means of performing measurements in situ, that is, with the

electrode under potential control in an appropriately designed electrochemical cell.

13
EXPERIMENTAL

Pure metal (either nickel or cobalt) and nickel-cobalt composite hydrous oxides were prepared by

cathodic (galvanostatic) electrodeposition2,6 from solutions of the corresponding metal nitrates on a solid

graphite electrode. After deposition, the films were rinsed with water and characterized in the

spectroelectrochemical cell first by cyclic voltammetry and subsequently by XAFS in deaerated 1.0 M

KOH. The cell for in situ fluorescence XAFS experiments employed in this study has been described

elsewhere.?

In situ XAFS measurements for the pure Ni and Ni/Co hydrous oxide films were conducted in the

discharged (i.e., reduced) and partially and nominally fully charged (i.e., oxidized) states. For the latter

experiments, the potential was first scanned to a value more positive than the onset of oxygen evolution and

then reversed to a value sufficiently negative for the current to drop essentially to zero but still positive with

respect to the onset of NiOOH reduction. This strategy made it possible to eliminate problems associated

with oxygen bubble formation during spectral acquisition. For measurements involving partially oxidized

films, the potential was scanned up to a value on the rising part of the Ni(OH)2 oxidation peak and then

reversed to a no-current voltage condition.

All experiments were performed at beam-line IV-2 at the Stanford Synchrotron Research

Laboratory. Details regarding these measurements as well as the method employed in the analysis of the

EXAFS data have been given in previous work.8

RESULTS AND DISCUSSION

I. ELECTROCHEMISTRY

The cyclic voltammogram of a composite 9:1 Ni/Co hydrous oxide in 1.0 M KOH, shown in curve

a, Figure 1, displayed characteristic oxidation and reduction peaks associated with the Ni(OH)2/Ni0OH

redox process. In agreement with the observations of Corrigan, 2 but at variance with those of Cordoba et

a1. 4, the overpotential for oxygen evolution was larger for this composite oxide than for a pure nickel

hydrous oxide prepared using the same procedure (see curve b, Figure 1). Also in harmony with earlier

data4 are the shifts in the Ni(OH) 2/NiOOH redox features in the negative direction (ca. 50 mV) induced by

the presence of Co in the Ni hydrous oxide lattice. As pointed out by Corrigan and Bendert, 2 no

voltammetry features associated with the Co(OH)2/CoOOH redox couple can be identified for this

composite Ni/Co film. This is particularly interesting, as the redox peaks of pure Co hydrous oxide in this

same electrolyte occur at a potential of 0.09 V and thus more negative than those observed for a pure Ni

film and therefore should be clearly discernable in the voltammogram.

II. XANES

A. Nickel K-edge

The Ni K-edge XANES for the composite Ni/Co hydrous oxide film in the reduced (fully

discharged) and oxidized (nominally charged) state recorded at -0.3 and +0.3 V vs. SCE, in 1.0 M KOH

are shown in Curves a and b, Fig. 2, respectively. These curves are nearly identical to those reported by

Pandya et a1. 9 for pure Ni hydrous oxide in 1.0 M KOH. Particularly noticeable is the shift in the overall

absorption edge region toward higher X-ray energies for the oxidized (N1 3+) film compared to the reduced

(N12+) film, including the peak at about 8363 eV. As discussed by Pandya et al. 9 the pre-edge feature at

about 8332 eV is ascribed to the is —^ 3d electronic transition,for which the intensity is found to be larger

14
for the oxidized than for the reduced state. It must be stressed that in both cases this spectral feature is

relatively small and therefore consistent with a slightly distorted octahedral environment.

B. Cobalt K-edge

The Co K-edge XANES for the Ni/Co composite film in the oxidized and reduced states are very

similar (see Curves a and b, Fig. 3) displaying a very small pre-edge peak. This observation provides

strong evidence that the nature of the cobalt sites is not modified by the structural and electronic changes in

the lattice associated with redox processes involving the nickel sites. A comparison between these curves

and those obtained for pure cobalt hydrous oxide films (prepared using the same electrodeposition method:

see curve c in this figure) 8 clearly shows that the cobalt sites in the latter composite hydrous oxide can be

assigned to Co3+. This is somewhat surprising since Co2+ in a pure cobalt hydrous oxide film undergoes

oxidation at a potential only slightly more negative than that observed for the Ni(O1-1)2/Ni00H couple.4

As is well known, however, the nature of the ligands can profoundly alter the redox potential of the

CO3+/Co2+ couple, e.g., Eo[Co(NH3 )6] 2+/3+ = 0.108 V, while E°[Co(H20)6]2+/3+ = 1.83 V. 10 It can

therefore be concluded that the environment of the cobaltic ions in the hydrated nickel oxide lattice is

such that the potential required for their reduction is shifted to much more negative values. Such a claim

has been been made earlier in the literature, although without much substantiation.l 1 Additional evidence

for the presence of Co 3+ sites in these composite hydrous oxides was obtained from the in situ Co Kedge

EXAFS data (vide infra).

III. EXAFS

A. Nickel Kedge

The Fourier transforms (FTs) (without phase shift correction) of the k 3x(k) Ni K-edge EXAFS for

the Ni/Co hydrous oxide films in the reduced (430 V vs. SCE) and oxidized (+0.30 V) states (see Curves

a and b, Fig. 4, respectively) are very similar to those reported by Pandya et al.9,12,13 for the oxidized and

reduced forms of battery-type nickel hydrous oxide electrodes. In accordance with their observations, the

Ni-O and Ni-Ni distances within the sheet-like Ni0 2 layers (1st and 2nd major shells, respectively) were

found to be somewhat smaller for the nominally oxidized film, d(Ni-0) = 1.92 ± 0.02 A, d(Ni-NO I =

2.82 ± 0.02 A, shown in Curve b, Fig. 4, compared to its reduced counterpart, d(Ni-0) = 2.09 ± 0.02 A,

d(Ni- NO, = 3.11 ± 0.02 A. The close similarity between the results obtained for the pure and composite

Ni/Co hydrous oxide films indicates that within the sensitivity of this technique, the presence of cobalt at a

nominally 10 metal-atom percent does not appear to significantly affect the structure of the nickel hydrous

oxide.

Unlike the behavior observed for composite Ni/Fe hydrous oxide films in the nominally fully

oxidized state, for which contributions due to N1 2+ were invariably observed, the Ni-Ni shell for a fully

oxidized 9:1 composite Ni/Co film displayed a single well-defined shell attributed to N1 3+ sites. This

affords unambiguous evidence that under these conditions the nickel in Ni/Co films can be completely

oxidized, leaving no residual Ni2+ ions in the lattice.

C. Cobalt K-Edge

A number of interesting conclusions can be drawn from the analysis of the FT of the k 3 x(k) Co Kedge EXAFS spectra of these composite Ni/Co films (see Figure 5). In particular, the best fit to the Co Kedge EXAFS data yielded a value of d(Co-0) = 1.90 ± 0.02 A and thus, within experimental error, is

15
identical to d(Co-O) in crystalline COOOH, i.e., 1.90 A.14 This observation is consistent with the presence

of cobaltic sites in the lattice, as proposed on the basis of the XANES data. More noteworthy, however, is

the fact that composite Ni/Co films purposely charged only partially (obtained at +032 V: see curve c,

Figure 5) displayed clearly resolvable Co-Ni shells associated with Co 3+-Ni2+ and CO3+-Ni3+ interactions.

These results clearly indicate that the co-electrodeposition procedure generates a single phase, mixed

metal hydrous oxide, in which cobaltic ions occupy nickel sites in the Ni0 2 sheet-like layers and not two

intermixed phases each consisting of a single metal hydrous oxide.

LITERATURE

1. McBreen, J. In Modern Aspects of Electrochemistry, Conway, B.; White, R.; Bockris, J. 0'M,

Eds.; Plenum Press: New York, 1991, Vol. 23.

2. Corrigan, D. A.; Bendert, R. M. J. Electrochem. Soc. 1989, 136, 723.

3. Corrigan, D. A. J. Electrochem. Soc. 1987, 134, 377.

4. Cordoba, S. I.; Carbonio, R. E.; Lopez Teijelo, M.; Macagno, V. A. Electrochem. Acta 1986, 31,

1321.

5. Abruna, H. D. In Electrochemical Interfaces, Abruna, H. D. Ed.; VCH: New York, 1991, Chapter

1.

6. Burke, L. D.; Twomey, T. A. M. J. Power Sources 1984, 12, 203.

7. Kim, S.; Bae, I. T.; Sandifer, M.; Ross, P. N.; Carr, R.; Woicik, J.; Antonio, M. R.; Scherson, D.

A. J. Am. Chem. Soc. 1991, 113, 9063.

8. Kim S.; Tryk, D. A.; Antonio, M. R.; Scherson, D. A. J. Phys. Chem. submitted, March 1993.

9. Pandya, K. I.; Hoffman, R. W.; McBreen, J.; O'Grady, W. E. J. Electrochem. Soc. 1990, 137,

383.

10. Milazzo, G.; Caroli, S. Tables of Standard Electrode Potentials, Wiley: New York, 1978, pp.

337.

11. Weininger, J. L. In The Nickel Electrode, Gunther, R. G.; Gross, S.; Eds., Proceedings of the

Symposium, The Electrochemical Society, 1982.

12. McBreen, J.; O'Grady, W. E.; Pandya, K. I.; Hoffman, R. W.; Sayers, D. E. Langmuir 1987, 3,

428.

13. Pandya, K. I.; O'Grady, W. E.; Corrigan, D. A.; McBreen, J.; Hoffman, R. W. J. Phys. Chem,

1990, 94, 21.

14. Delaplane, R. G.; Ibers, J. A.; Ferraro, J. R.; Rush, J. J. J. Chem. Phys. 1969, 50, 1920.

ACKNOWLEDGEMENTS

This work was supported in part by the Department of Energy through a subcontract from

Lawrence Berkeley Laboratory. Additional funding was provided by Eveready Battery Co., Westlake,

Ohio.

16
1.5

1.0

_o

.D

4)

N

0 0.s

Z

0.0

Q

E

C

Q)

L

U

-U.4 -V.L U.V U.L U.4 U.O

Potential (V vs SCE)

a

b

8320 8340 e350 8380 840

Energy (eV)

Fig. 1 Cyclic voltammetry curves for composite 9:1 Ni/Co Fig. 2 Ni K-edge XANES for a composite 9:1 Ni/Co hydrous

(curve a) and a pure Ni (curve b) hydrous oxide oxide film in the fully discharged (Curve a) and nominally

electrode in deaerated 1.0 M KOH. fully oxidized (Curve b) recorded at -03 V and +03 V vs

Scan rate: 5mV s t . Cross sectional electrode area: 0.74 cm2. SCE, respectively in 1.0 M KOH

17
i

a ^>

C

,

,

,

,

i'

1.6

1.4

1.2

1.0

0

0.6

N

9 0.6

O

Z

0.4

0.2

0.0

N

U a

C

Q1

LL

b

7700 7720 7140 7760

Energy (eV)

Fig. 3 Co K-edge XANES for a composite 9:1 Ni/Co hydrous

oxide film in the oxidized (0.3V) and reduced (43V)

states (Curves a and b, respectively). Also shown in

in this figure are the corresponding spectra for a pure

Co hydrous film prepared by the same electrodeposition

procedure described in the Experimental Section in the

reduced state (Carve c).

0 2 4 6

r (A)

Fig. 4 Phase-uncorrected FT for the k3X(k) Ni K-edge EXAFS

for 9:1 Ni/Co hydrous oxide films in the reduced (43 V vs

SCE, Curve a), and oxidized (+0.30 V, Curve b) states.

18
a

to

v^ C

ctt

U_

b

r (A)

Fig. S Phase-uncorrected FT for the k 3X(k) Co Kedge EXAFS

for 9:1 Ni/Co hydrous oxide films in the reduced (Curve a),

fully oxidized (Curve b) and partially oxidized (obtained

at +432 V, Curve c) states. Other conditions are given in

the Caption Figure 2.

19

L
Pn_S (•

N 4 2 3.3 4

REVIEW OF SUPER Ni/Cd CELL DESIGNS AND PERFORMANCE

Bruce Abrams-Blakemore

Eagle-Picher. Inc.

Colorado Springs, CO 80911

ABSTRACT

Eagle-Picher Industries, Inc., in cooperation with Hughes Aircraft Company, began production of the

Super Nickel-Cadmium cell in 1989. Since that time the Super Nickel-Cadmium cell has been

deployed in a wide variety of satellites. This paper will review one of those programs and provide a

performance update. We will discuss storage requirements and capacity histories for the various

Super NiCad Cell designs.

INTRODUCTION

The Super NiCad was first introduced in 1989 by Hughes Aircraft Company for use in the LESAT

satellite. The cell is manufactured at Eagle-Picher Industries in Colorado Springs, Colorado. The

Super NiCad differs from standard nickel cadmium cells in several ways:

The positive plate is made by an electrochemical-chemical deposition process. This

process is based on the Air Force patented buffered deposition process.

2. The positive plate loading target is in the range of 1.65 g/cc void. This is lower than

the standard NiCad positive loading which is in the range of 2.0 g/cc void.

3. The negative electrode in the Super NiCad has similar loading levels to the standard

NiCad but is produced by an aqueous electrochemical chemical deposition method.

4. The negative to positive ratio is much higher in the Super NiCad normally exceeding a

theoretical negative to positive of 1.8. This allows for precharge in the range of 40%

of excess negative capacity.

5. Finally, the Super NiCad incorporates an impregnated zirconia separator. This separator is as wettable as the nylon used in standard NiCads, however it is much more

resistant to the alkaline environment.

Because of the unique design there are some special storage considerations with Super NiCad:

1. Unlike the standard NiCad, the Super NiCad should never be stored under short-circuit.

2. The Super NiCad is best stored under trickle charge.

3. When constant charging during storage is impractical, the Super NiCad can be periodically "topped-off."

21

PRfiCAMIZ PAGE PLANK NOT FfLoVED
4. When any type of charging is impractical, the Super NiCad is best stored open-circuited

and completely discharged at low temperature.

The Super NiCad was developed to correct observed failure mechanisms of the standard nickel

cadmium cells. It was anticipated that the changes discussed above would provide some performance

advantages:

The deposition method and loading limits were to provide better active material

utilization, better ability to withstand deep discharge and less plate expansion over cycle

life.

2. The negative manufacturing process was designed to improve active material utilization

and decrease plate deterioration, thereby minimizing cadmium migration.

3. The higher precharge had better gas management as its goal, additionally there was the

expected benefit of better overcharge protection.

4. The zirconia separator was chosen in order to avoid the oxidation to carbonate that was

seen with the nylon separator.

The purpose of this paper is to explore functioning Super NiCad programs and determine how well

the cells have met performance goals. Specific data will be drawn from the SMEX 9 AH program.

We will first look at the beginning of life (BOL) performance characteristics. Next we will discuss

cell performance during accelerated life cycle testing. Finally, we will briefly discuss cell performance in flight.

ACCEPTANCE AND STRESS TESTING

Beginning of Life

All Super NiCad cells that are destined for flight or accelerated life cycling are sent through an

acceptance test regime (see Figure 1). The test is based on the NASA approved flight cell acceptance

test. It involves testing at 0° C, 10° C and 22° C. Charge currents are varied from trickle charge

(C/20) to rapid charge (C/2). The cells are normally discharged at a C/2 current, but there are tests

to demonstrate charge retention (self-discharge rate) and high current discharge performance (at 5C).

In addition to the electrical testing, representative plates are subjected to a destructive stress test (see

Figure 2).

Since the normal satellite operating temperature is around 10° C, all capacity qualification testing is

performed at this temperature. The Super NiCads have demonstrated from 85 to 95% utilization at

this early stage of cell life.

22
ACCEPTANCE TESTING PROCEDURE

Capacity Test at 10° C

- Cells are shorted down to 0.010 volt maximum, 16 hour stand.

- Cells are soaked at 10° C for 4 hours.

- C/10 charge for 20 hours, 1 hour open-circuit stand, C/2 discharge to 1.0 volt.

Overcharge and Capacity Test at 0° C

- Cells are shorted down to 0.010 volt maximum, 12 hour minimum.

- Cells are soaked at 0° C for 4 hours.

- C/20 charge for 72 hours, 1 hour open-circuit stand,

C/2 discharge to 1.0 volt.

High Current Discharge Test at 22° C

- Cells are shorted down to 0.010 volt maximum, 12 hours minimum.

- Cells are soaked at 22° C for 4 hours.

- C110 charge for 20 hours, 5C pulse discharge for 10 seconds, C/2 discharge to 1.0

volt.

Charge Retention (Discharged) Test at 22° C

- Cells are shorted down for 12 hours.

- C110 charge for 10 minutes, open-circuit for 24 hours, minimum voltage 1.15 volts.

Figure 1. Description of EPI Acceptance Test.

23
EPI POSITIVE ELECTRODE STRESS TEST

Hot Formation: 5C rate, 70° C, 20% KOH, 2 cycles.

Pre-Stress: Weight, thickness and visual appearance characterization.

Initial Capacity: Charge at 5C for 12 min., C/2 for 1 hour. Discharge at C rate to 1.0 V

vs. Ni sheet.

200 Cycles: Charge at 10C for 12 minutes. Discharge at 10C for 8 min. Diode clamped

to -1.0 V.

Ending Capacity: 5 Cycles, same as initial capacity.

Visual appearance.

Scrub, rinse, and dry.

Weight loss and swelling.

Figure 2. Stress Test Regime.

After 200 stress cycles (charge and discharge at 10°C) the plates show less than 3% thickness growth

and no blistering. By comparison, electrodes for standard NiCads seldom survive the stress test

intact.

Cell voltages at 0°C do not exceed 1.50 volts on overcharge at C/20 (see Figure 3).

1.530

1.510

"0° C.

LU 1.490 C/20 RATE,

ti 1.470--

1.450--

1.430--

1.410 1 1 i

0 10 20 30

TEMPERATURE 'C

Figure 3. Typical Super NiCad End-of-Charge Voltage vs. Temperature, C/10 Rate.

24
Typically all cells of a given lot perform within a 20 millivolt range.

Figure 4 depicts average cell performance values for the SMEX acceptance test.

SMEX AVERAGE ACCEPTANCE TEST RESULTS

0°C Test 10°C Test

Average Capacity 11.55 AH 11.76 AH

Average EOC Voltage 1.491 1.4824

Average EOC Pressure 2.5 psi 11.4 psi

Average Final Voltage after 10 Second 5C Discharge: 1.1429V

Figure 4. Results of SMEX acceptance testing.

ACCELERATED LIFE CYCLING

The super NiCad has been tested in both simulated LEO and GEO orbits and has passed all NASA

requirements to date.

FLIGHT PERFORMANCE

The super NiCad has demonstrated end-of-charge voltages in the ranges of 1.40 to 1.48 and end-ofdischarge voltages in the ranges of 1.18 to 1.21. These end-of-charge voltages are approximately 20

millivolts lower than those seen in standard nickel cadmium cells and the discharge voltages are

similarly higher than those seen in standard nickel cadmium cells.

Initial flight data for the SMEX program agree with the trends that were suggested by the accelerated

testing.

CONCLUSION

Beginning-of-life test data show increased active material utilization in the super NiCad design, as

anticipated. Cell voltages and pressures are well within acceptable limits. Stress testing indicates that

the Super NiCad positive has a longer theoretical life than the standard Nickel cadmium positive.

Accelerated life cycling has provided a positive picture of the long term Super NiCad performance.

Super NiCad cells have surpassed the life normally expected in a standard NiCad. All super NiCads

currently in flight are performing well and within all specified limits.

25
911q0/1/

p_,lq N94-23a

THERMODYNAMICS OF NICKEL-CADMIUM AND NICKEL-HYDROGEN BATTERIES

Digby D. Macdonald and Mark L. Challingsworth

Center for Advanced Materials

The Pennsylvania State University

517 Deike Building

University Park, PA 16802

ABSTRACT

Thermodynamic parameters for Nickel-Cadmium (NiCad) and Nickel-Hydrogen (NiH2) batteries are calculated for

temperatures ranging from 273.15K (0°C) to 373.15K (100°C). For both systems, we list equilibrium and

thermoneutral voltages for the cells, and in the case of the NiH2 battery, these data are provide for hydrogen

fugacities ranging from 0.01 to 100 (atm) to simulate the full discharged and charged states. The quality of the input

thermodynamic data are assessed and the effect of assuming different cell reactions is analyzed.

INTRODUCTION

Nickel-cadmium (NiCad) and nickel-hydrogen (NiH2) batteries are used extensively in aerospace systems, because

of their relatively high energy and power densities and their excellent cycling capabilities [1-4]. However, given the

severe demands placed on cycling performance for satellites in low earth orbit (LEO) and recognizing the need to

keep battery weight to an absolute minimum, efforts continue to improve the cycle life particularly to high depths of

discharge. The current standard in this regard for NiH2 batteries is a cycle life of 30,000 cycles to 50% loss in

capacity (of the nickel electrode) at 50% depth of discharge. The cycle life of NiCad batteries is not as favorable

(generally a few thousand cycles), so that these systems are commonly not employed where cycle fife is the key

operational characteristic.

It is well recognized that the cycle lives of NiCad and NiH2 batteries are limited by irreversible structural and

physicochemical changes that occur in the porous nickel electrodes as the batteries are cycled between the charged

and discharged states. Although the mechanistic details of the degradation processes are not well-understood,

various models have been developed that provide a good phenomenological simulation of the loss of capacity [5,6].

However, a complete understanding of the degradation mechanisms requires a careful analysis of the cell

thermodynamics so that the voltage and energy efficiencies can be placed on a rational basis. Thermodynamic data

for a variety of alkaline batteries has been published previously by Pound, Singh, and Macdonald [7] and for sealed

NiCad batteries by Hodge et al [8]. The present work recalculates and expands the thermodynamic database for the

27

PRrEG613tNG PAGE BLANK NOT MMUD
NiCad and NiH2 systems by deriving values for relevant parameters (equilibrium cell and thermoneutral potentials)

over a wide range of temperature (0'C — 100'C) and, in the case of NiH2, hydrogen pressure (0.01-100 atm).

BATTERY THERMODYNAMICS

Following conventional formalism, the cell reactions for NiCad and NiH2 batteries can be written as

Cd + 2 NiOOH + 21­12O = Cd(OH)2 + 2Ni(OH)2 (1)

and

2 H2 + NiOOH = Ni(OH) 2

(2)

respectively, assuming that the charged nickel active material is NiOOH, or as

Cd + 2NiOOH • H2O = Cd(OH)2 + 2Ni(OH)2 W)

and

H2 + NiOOH • H2O = Ni(OH) 2 + H2O

(2')

if NiOOH • H2O is assumed to be the active material at the nickel positive electrode [8]. The equilibrium voltages

for these four reactions are given by the Nernst equation as

2.303RT

cell = AE°cell + F

logaH2O (3)

eo 2.303RT 1/2

cell = AEO F

log f H2 (4)

AEcell = AEcell (3')

and

eo 2.303RT 1/2 2.303RT ,

cell _ cell + F lOgfH2 — F l°gaH2O (4 )

where fH2 is the fugacity of hydrogen, aH 2O is the activity of water, and the standard cell voltages are given by

0 DEcell = -AG0cell / nF (5)

28
AGco,ll , n, and F are the change in standard Gibbs energy for the cell, the number of electrons involved in the cell

reaction, and Faraday's constant, respectively. The change in standard Gibbs energy for the cell is

AGOccell - l vpAfGP — ^,uR^fGR

Products Reactants

where A fGI is the standard Gibbs energy of formation of component I, n is the stoichiometric coefficient for the

reaction, and subscripts P and R refer to "products" and "reactants," respectively, in the cell reactions (i.e. Reactions

( 1 ), (2), (1'), and (2'))•

A full assessment of the thermodynamics of a battery requires the estimation of the thermoneutral potential of the

cell, defined as

,Tcell — —AHcell / nF (7)

and hence

AE T aAEc

e

ell

cell :- AEc

e 

ell — T aT

(g)

Using standard thermodynamic identities, we derive expressions for AE ell for the four cells of interest as

T AHi RT2 a1naH2p

^cell,l - — 2F F aT (9)

OHo

AEccell,2 = — F 2 (10)

^Ho

AEell 1 2F (11)

T OH2, RT2 a In aH2p

AEcell,2' - — F + F aT

(12)

where OHk is the change in standard enthalpy for Reaction k. Note that a term involving the fugacity of hydrogen

does not appear in Equations (10) and (11) [corresponding to Reactions (2) and (2')], because we have assumed that

fH2 is a constant, independent of temperature.

(6)

and

29
THERMODYNAMIC DATA

The thermodynamic data employed in this study are summarized in Tables 1 and 2, together with the sources from

which they were taken. Heat capacities for Ni(OH)2, Cd(OH)2, Ni00H • H2O, and Ni00H had to be estimated

using CWS° ratios, which are found to be constant for given classes of compounds [7]. The value adopted for the

Gibbs energy of formation for Cd(OH)2 was adjusted slightly from that given by Naumov et al (Ref. (a), Table 1),

474.5 U/mol, so as to yield the observed equilibrium voltage for the NiCad battery [ 1.45V (2)]. The accuracies of

the various thermodynamic quantities summarized in Table 1 are largely unknown and are difficult to estimate

without detailed information on the original experiments from which they are derived.

Noting that heat capacities are generally expressed in the form

C°=A+BT+CT'2

P

(13)

we derived expressions for the Gibbs energy and enthalpy of formation of any given compound at temperature T as

follows.

0 G° 0 G°o — (T — T )4 S O + A T — T T In T

f T = f T o f T f o—

TO

2 2

+Bf T —To _T(T—TO ) +Cf 1 — 1 ][T(l + 1 )_,]

2 IT To 2 T To

AfHT= OfHo +Af(T—To )+ Bf [T2 —To] —C 1 — 1

2 T To

and

OfST = `^fHT -ofGT)/T

where O fSo

 and AfHO are the entropy and enthalpy of formation at the reference temperature (To= 298.15K)

and Af, B f, and Cf are the coefficients in Equation (13) for the formation of the compound of interest.

Thermodynamic parameters of formation calculated using Equations (14) to (16) for various active components in

NiCad and NiH2 batteries are summarized in Table 3 as a function of temperature from 273.15K (0'C) to 373.15K

(14)

(15)

(16)

30
(100°C). Note that no values are given for Cd and H2 because, by definition, the parameters of formation are zero at

all temperatures.

Calculation of the equilibrium and thermoneutral cell voltages requires a knowledge of the activity of water, which

for 8 molal KOH can be expressed as (7)

In aH2O = —6.345750 + 6.125771 x 10 3 T — 2.093874 x 106 / T2

(17)

+2.324516 x 108 T3

The activity of water for this electrolyte ranges from 0.514 at 263.15K to 0.618 at 393.15K.

The data summarized in Table 3 were used to estimate equilibrium and thermoneutral cell voltages for NiCad and

NiH2 batteries as a function of temperature, as given in Tables 4-7. We should note, at this point, that the

thermodynamic data for NiOOH • H2O are poorly known, so that lower confidence should be placed in the

potentials calculated from Reactions (Y) and (2') than is those calculated from Reactions (1) and (2). Accordingly,

in this work, we will emphasize the thermodynamic calculations that involve NiOOH as the oxidized, active material

at the positive electrodes of NiCad and NiH2 batteries.

DISCUSSION

As seen from the data summarized in Tables 4 and 5, the equilibrium potentials for NiCad and NiH2 batteries with

8m KOH electrolyte decrease with increasing temperature and, for the latter system at any given temperature,

increase with increasing hydrogen partial pressure. The dependence of AEcell e on temperature results directly from

the entropy and heat capacity changes for the cell reactions as well as from the change in water activity in the

electrolyte, whereas the dependence on hydrogen pressure (in the case of NiH2 batteries) results directly from the

thermodynamics of the negative electrode. We also note that much larger changes in DEcell with temperature are

predicted if the oxidized nickel phase is assumed to be NiOOH • H2O rather than NiOOH, but the poor quality of the

data for the former precludes any in-depth analysis of this difference.

Thermoneutral potentials for NiCad and NiH2 batteries are listed in Tables 6 and 7. The thermoneutral voltage

corresponds to that voltage that the cell would have to operate at so that the entropic dissipation of energy is zero.

Of course, the data calculated in this work do not include the various irreversible sources of energy dissipation,

associated with heat generation due to the internal cell resistance and arising from the reactions occurring at the

31
electrodes on charging and discharging. These irreversible effects are best treated using irreversible thermodynamic

methods of the type proposed by Ratjke et al [9].

For both the NiCad and NiH2 batteries, the thermoneutral voltage is greater than the equilibrium cell voltage at

equivalent temperatures and hydrogen pressure (for the NiH2 system). This relationship arises from the fact that the

entropy changes for the cell reactions on discharge are negative, corresponding to exothermic processes. Likewise,

on charging both cells are endothermic provided that the voltage lies between AEcell and cell'

Finally, we note that the equilibrium cell potentials calculated in this work for the NiH2 battery are in good

agreement with our previous calculations and with experimental data [7]. A similar comparison cannot be made

with experimental data for NiCad batteries because we have used these data to adjust the Gibbs energy of formation

of Cd(OH)2 used in the calculations. However, it is worth noting that our equilibrium cell potential at 25°C (1.45V)

is significantly more positive than that calculated by Hodge et al [8] (1.29V). This difference may be attributed to

the large uncertainty in the Gibbs energy of formation for NiOOH • H2O [8] and to the fact that Hodge et al [8] did

not apply a correction for the activity of water. Clearly, the calculated value of Hodge et al [8] is much lower than

the open circuit voltage of 1.45V quoted by Halpert [2]. Because irreversible processes within the cell are expected

to reduce the measured open circuit voltage below the equilibrium cell voltage, it is evident that the calculations of

Hodge et al (8) are at odds with experiment.

REFERENCES

1. P.C. Milner and U.B. Thomas, in "Advances in Electrochemistry and Electrochemical Engineering," Vol. 5, p. 1

(Interscience, N.Y., 1967).

2. G. Halpert, J. Power Sources, 1^, 119 (1985).

3. J.K. McDermott, J. Power Sources, 1$, 265 (1986).

4. L. Miller, J. Brill, and G. Dodson, J. Power Sources, 29, 533 (1990).

5. D.D. Macdonald, B.G. Pound, and S.J. Lenhart, J. Power Sources, 22, 447 (1990).

6. D.D. Macdonald, M. Urquidi-Macdonald, S.D. Bhatka, and B.G. Pound, J. Electrochem. Soc., 13a, 1359

(1991).

7. B.G. Pound, R.P. Singh, and D.D. Macdonald, J. Power Sources, .La, 1 (1986).

8. B.J.R. Hodge, R. Bonnaterre, and F. Putois, J. Power Sources, 2, 211 (1975).

9. S.K. Ratkje, T. Ikeshoji, and K. Syverud, J. Electrochem. Soc., 137, 2088 (1990).

32
TABLE 1

Thermodynamic Properties of Components at 298.15K

Component OfGo

(kJ/mol)

A fHo

(kJ/mol)

0 fSo (f)

(J/K.mol)

So

(J/K.mol)

AfCP (a,b,f)

(kJ/mol)

Ni 0 0 0 29.87 16.987 + 0.02946T

Cd 0 0 0 51.756 a 22.217 + 0.012301T

02 0 0 0 205.03 29.957 + 0.00414T - 1.674 x 10 5 T-2

H2 0 0 0 130.574 a 27.2797 + 0.0032635T + 0.50208 x 10 5 T-2

H2O -237.178 a -285.828 d -163.174 69.915 a 10.669 + 0.042284T-6.90308 x 105T-2

Ni(OH)2 1159.070 -541.799 d -277.474 88.0 18.106 - 0.03686T + 1 .

 17 19 x 10-5T-2

Ni00H -321.7 -391.231 d -233.207 66.98 30.186 - 0.03523T + 1.423 x 105T-2

Ni00H•H20 1 -564.422 c -629.859 d -219.478 150.624 c 87.388 - 0.082844T + 8.661 x 105T-2

Cd(OH)2 -482.347 -571.642 d -299.496 87.864 a 12.726 - 0.0197T + 1.17192 x 105T-2

(a)G.B. Naumov, et al, Handbook of Thermodynamic Data, USGS-WRD-74-001, US Geol. Survey, 1974.

(b)P.G. Pound, R.P. Singh, & D.D. Macdonald, J. Power Sources,1$, 1 (1986).

(c)B.J.R. Hodge et al, J. Power Sources, Z, 211 (1975).

(d)Estimated from 0 fHo = 0 fGo + TA fS°.

(e)Calculated for the formation reaction from the elements.

(f) D.D. Wagman et al, J. Phys. Chem. Ref. Data, 11, Suppl. No. 2 (1982).

TABLE 2

Heat Capacities of Active Compounds (a,b,f, Table 1)

Compound C (kJ/mol)

H2O 52.928 + 47.614 x 10 -3 T - 7.238 x 105 T-2

Ni(OH)2 92.33

NiOOH 90.77

NiOOH•H20 200.09

Cd(OH)2 92.18

33
TABLE 3

Thermodynamic Parameters of Formation for Active Components in NiCad

and NiH2 Batteries as a Function of Temperature

Temperature Gibbs Energy Enthalpy Entropy

(K) (J/mol) (J/mol) (J/K-mol)

Ni(OH)2

273.15 -466014.9 -541988.4 -278.1382

283.15 -463234.9 -541909.9 -277.8557

293.15 -460457.7 -541835.1 -277.5959

298.15 -459070 -541799 -277.474

303.15 -457682.9 -541763.9 -277.3572

313.15 -454910.5 -541696.4 -277.1382

323.15 -452140.1 -541632.6 -276.9376

333.15 -449371.7 -541572.5 -276.7544

343.15 -446605 -541516.1 -276.5875

353.15 -443839.8 -541463.4 -276.436

363.15 -441076.2 -541414.3 -276.299

373.15 -438313.8 -541368.9 -276.1758

NiOOH

273.15 -327593.5 -391776.8 -234.9745

283.15 -325230.3 -391554.5 -234.237

293.15 -322874.9 -391337 -233.5396

298.15 -321700 -391230 -233.205

303.15 -320526.8 -391124.2 -232.8793

313.15 -318185.7 -390915.9 -232.2536

323.15 -315851.2 -390712 -231.66

333.15 -313522.9 -390512.6 -231.0964

343.15 -311200.5 -390317.4 -230.5608

353.15 -308883.8 -390126.4 -230.0517

363.15 -306572.3 -389939.7 -229.5673

373.15 -304266 -389757 -229.1063

NiOOH•H2O

273.15 -570232.1 -631718.9 -225.1038

283.15 -567888.4 -630963.5 -222.763

293.15 -565571.2 -630224 -220.5461

298.15 -564422 -629860 -219.481

303.15 -563278.9 -629499.7 -218.4432

313.15 -561010.3 -628789.9 -216.4453

323.15 -558764.1 -628093.9 -214.5448

333.15 -556539.2 -627411.5 -212.7346

343.15 -554334.6 -626742 -211.0085

353.15 -552149.4 -626085.1 -209.3611

363.15 -549982.6 -625440.3 -207.7873

373.15 -547833.4 -624807.5 -206.2827

34
TABLE 3 (continued)

Temperature

(K)

Gibbs Energy

(J/mol)

Enthalpy

(J/mol)

Entropy

(J/K-mol)

Cd(OH)2

273.15 -489876.4 -571855.4 -300.1235

283.15 -486862.4 -571767.8 -299.8593

293.15 -483851.4 -571683.2 -299.6129

298.15 -482347 -571642 -299.496

303.15 -480843.3 -571601.5 -299.383

313.15 477837.8 -571522.6 -299.1684

323.15 -474834.8 -571446.4 -298.9677

333.15 -471834 -571372.9 -298.7802

343.15 -468835.5 -571302.1 -298.6049

353.15 465839 -571233.7 -298.441

363.15 -462844.4 -571167.9 -298.2878

373.15 459851.5 -571104.5 -298.1446

H2O

273.15 -241078.8 -286186.8 -165.1239

283.15 -239514.8 -286051.7 -164.3391

293.15 -237955.6 -285906.4 -163.5557

298.15 -237178 -285830 -163.165

303.15 -236401.7 -285751.3 -162.7751

313.15 -234852.9 -285587 -161.9981

323.15 -233309.6 -285414 -161.2256

333.15 -231771.8 -285232.7 -160.458

343.15 -230239.5 -285043.4 -159.6957

353.15 -228712.8 -284846.5 -158.939

36115 -227191.8 -284642.2 -158.188

373.15 -225676.6 -284430.8 -157.4429

35
TABLE 4

Calculated Equilibrium Cell Voltages for NiCad Batteries as a Function of Temperature

e

°Eeell (V)

T/K Reaction (1) Reaction (1')

273.15 1.461 1.459

283.15 1.457 1.438

293.15 1.453 1.418

298.15 1.450 1.408

303.15 1.448 1.397

313.15 1.444 1.377

323.15 1.439 1.356

333.15 1.435 1.334

343.15 1.430 1.313

353.15 1.425 1.292

363.15 1.420 1.270

373.15 1.415 1.748

36
TABLE 5

Calculated Equilibrium Cell Voltages for Nil-12 Batteries as a Function

of Temperature and Hydrogen Pressure

T/K 0.01 0.1

PH2/atm

1.0 10.0 100.0

Reaction 2

273.15 1.380 1.408 1.435 1.462 1.489

283.15 1.374 1.402 1.430 1.458 1.487

293.15 1.368 1.397 1.426 1.455 1.484

298.15 1.365 1.394 1.424 1.453 1.483

303.15 1.361 1.391 1.422 1.452 1.482

313.15 1.355 1.386 1.417 1.448 1.479

323.15 1.348 1.381 1.413 1.445 1.477

333.15 1.342 1.375 1.408 1.441 1.474

343.15 1.335 1.369 1.403 1.437 1.471

353.15 1.329 1.364 1.399 1.434 1.469

363.15 1.322 1.358 1.394 1.430 1.466

373.15 1.315 1.352 1.389 1.426 1.463

Reaction (2')

273.15 1.459 1.432 1.405 1.378 1.351

283.15 1.440 1.412 1.384 1.356 1.327

293.15 1.420 1.391 1.362 1.333 1.304

298.15 1.410 1.381 1.351 1.322 1.292

303.15 1.401 1.371 1.341 1.311 1.281

313.15 1.382 1.351 1.320 1.289 1.258

323.15 1.362 1.330 1.298 1.266 1.234

333.15 1.342 1.309 1.276 1.243 1.210

343.15 1.322 1.288 1.254 1.220 1.186

353.15 1.302 1.267 1.232 1.197 1.162

363.15 1.282 1.246 1.210 1.174 1.138

373.15 1.262 1.225 1.188 1.151 1.114

37
TABLE 6

Calculated Thermoneutral Cell Voltages for NiCad Batteries as a

Function of Temperature

AE Tcell M

T/K Reaction (1) Reaction (1')

273.15 1.580 2.033

283.15 1.574 2.040

293.15 1.570 2.046

298.15 1.569 2.050

303.15 1.568 2.053

313.15 1.568 2.059

323.15 1.569 2.065

333.15 1.572 2.071

343.15 1.575 2.077

353.15 1.580 2.083

363.15 1.585 2.089

373.15 1.591 2.095

38
TABLE 7

Calculated Thermoneutral Cell Voltages for NiH2 Batteries as a Function of

Temperature and Hydrogen Pressure

pH2/atrn

T/K 0.01 0.1 1.0 10.0 100.0

273.15 1.611 1.584 1.557 1.530 1.503

283.15 1.615 1.586 1.558 1.530 1.502

293.15 1.618 1.589 1.560 1.531 1.502

298.15 1.620 1.590 1.561 1.531 1.501

303.15 1.621 1.591 1.561 1.531 1.501

313.15 1.625 1.594 1.563 1.532 1.501

323.15 1.628 1.596 1.564 1.532 1.500

333.15 1.63 2 1.599 1.%6 1.533 1.500

343.15 1.635 1.601 1.567 1.533 1.499

353.15 1.639 1.604 1.569 1.533 1.498

363.15 1.642 1.606 1.570 1.534 1.498

373.15 1.645 1.608 1.571 1.534 1.497

Reaction (2')

273.15 2.065 2.038 2.011 1.983 1.956

283.15 2.081 2.053 2.025 1.997 1.969

293.15 2.094 2.065 2.036 2.007 1.978

298.15 2.100 2.071 2.041 2.012 1.982

303.15 2.106 2.076 2.046 2.016 1.986

313.15 2.116 2.085 2.054 2.023 1.992

323.15 2.124 2.092 2.060 2.028 1.996

333.15 2.131 2.098 2.065 2.032 1.999

343.15 2.137 2.103 2.069 2.035 2.001

353.15 2.142 2.107 2.072 2.037 2.002

363.15 2.156 2.110 2.074 2.038 2.002

373.15 2.149 2.112 2.075 2.038 2.001

39
N 9 4 2 1

HIGH SPECIFIC ENERGY, HIGH CAPACITY NICKEL-HYDROGEN CELL DESIGN

James R. Wheeler

Eagle Picher Industries

Joplin, Missouri 64801

ABSTRACT

A 3.5 inch rabbit-ear-terminal nickel-hydrogen cell has been designed and tested to deliver high capacity at steady

discharge rates upto and including a Crate. Its specific energy yield of 60.6 wh/kg is believed to be the highest yet achieved

in a slurry-process nickel-hydrogen cell, and its 10°C capacity of 113.9 AH the highest capacity yet of any type in a 3.5

inch diameter size. The cell also demonstrated a pulse capability of 180 amps for 20 seconds. Specific cell parameters

and performance are described. Also covered is an episode of capacity fading due to electrode swelling and its successful

recovery by means of additional activation procedures.

CELL DESCRIPTION

A development program was initiated with the objective of maximizing power and specific energy for 3.5 inch

diameter nickel-hydrogen cells while still retaining the long-cycle life and ruggedness of the positive slurry electrode.

Eagle-Picher designed cells to meet this objective as part of a joint project with a major aerospace company. We have

now completed and tested a 100 ampere-hour cell design in two separator versions. One version has a single layer of Zircar

(zirconia-fiber cloth) separator for each positive electrode (hereafter termed -3 cells), and the other version has a single

asbestos separator (-5 cells), but they are otherwise identical and were made from the same lots of electrodes and other

components.

The electrolyte quantity for single layer Zircar is just slightly larger than the single layer asbestos version, and this

should ensure at least equal cycle life. At Eagle Picher, single-layer Zircar cells in an on-going life test have exceeded

20,000 cycles at a depth of discharge of 15%. End of discharge and charge voltages have thus far been stable.

A photograph of a sample cell (both versions are externally the same) and a table of weights and dimensions are

provided by Figure 1 and Table 1. All of the cells were equipped with strain-gage pressure monitors.

The internal cell design is a dual stack with a back-to-back electrode configuration and continuous leads to rabbitear terminals. The rabbit-ear terminal is a feature which permits a shorter battery height and therefore a shorter thermal

path when the cells are vertically mounted. Thus, the high specific energy can be further improved at the battery level

by reducing the length of cell sleeve mountings and cell inter-connections. If the cells are mounted parallel to abase plate,

the rabbit-ear terminals help there as well because the cells can be rotated to minimize the length of the interconnects.

The positive electrodes are slightly thicker and more porous than Eagle-Picher's standard high-bend-strength slurry

plaque. They are approximately double the bend-strength of dry sinter, but only about 3/4 the strength of the standard

slurry product. Bend strength may be considered a measure of the resistance of the electrode to plate swelling.

Four Zircar cells and three asbestos cells were built. The electrolyte concentration was set at about 39% in order

to maximize the specific energy of the cell. It was hoped that this high concentration would not be so great as to cause

the positive electrodes to swell, since it has proved compatible with the standard electrodes in the past. As we were to

learn later, this was not the case.

A catalyzed wall wick is incorporated for improved thermal operation and gas management, making the cell suitable

for either LEO or GEO applications. The final electrolyte levels for the two versions were about the same, with the single

Zircar version holding, on average, only 1% more per cell than the single asbestos version.

The pressure vessel is Inconel 718. MEOP translates to a minimum burst safety factor of 2.5. The actual maximum

pressure reached by the highest Zircar version cell under conditions of severe overcharge was 3% greater than the value

for the asbestos.

41

PrRECIDiIVQ PAGE BLANK NOT E4i..MED
INITIAL CELL TESTING

Four Zircar-separator cells and three asbestos cells were built. After activation, the cells were tested to the customer's

performance specification, which conservatively assumed a rated capacity of 90 AH. For the initial acceptance testing

therefore, standard capacity charges were 9 amps for 16 hours. The discharge rate used was 68 amps. Performance data

are summarized in tables 2 and 3. Average discharge voltage performance for each type of cell is shown by the curves

in Figure 2. Mid-discharge values (45 minutes) are comparable to those achieved in shorter 3.5 inch cells, indicating that

the internal bussing has an adequate cross-section. The voltage advantage of Zircar over asbestos is apparent, and is

essentially the same for the single Zircar configuration as would be expected from a double Zircar design.

Testing also included a 180 amp pulse test (See Figure 2) . The pulse was conducted after a 16 hour, 9 amp charge

and 15 minutes of the 68 amp discharge. Pulse duration was 20 seconds. Minimum terminal voltages reached at the end

of the pulse were 1.17 V for the Zircar version cells, and 1.13 V for the asbestos cells.

The Zircar version achieved 60.1 watt-hours per kilogram at 10°C to 1.0 V at a discharge rate of 68 amps (C/1.47

for a 100AH nominal rating). (Watt-hour values were not merely estimated on mid-point voltages, but were actually

measured by the automated data collection software.) In the same test the asbestos version achieved 59.0 watt-hours per

kilogram. Capacities would have been greater and discharge voltage higher at a more-normal C/2 rate. They would also

likely have been higher had the cells been charged closer to a true C/10 rate, i.e. at 10 amps rather than 9.

Without the strain gages, the values would be 60.6 wh/kg and 59.4 wh/kg respectively. These are believed to be the

highest energy densities yet achieved at the cell level for nickel-hydrogen cells with slurry-process positive electrodes.

The performance exceeds other large cells recently reported, including some 83AH cells [ I ] and the 90AH HST cells [2].

The evolution of size and power of slurry-type cells at Eagle-Picher is shown in Figure 3. The progression in efficiency

has paralleled a rather steady increase in cell size and increasing operating pressures over the last 10 years.

Capacity retention at 10°C was measured by charging the cells for 16 hours at 9.0 amps, and, after an open circuit

stand of 72 hours, discharging at 68 amps to 1.0 V. The percentages of capacity retained, when compared to the standard

10°C test, were 84.6% for Zircar, and 85.8% for asbestos. These are in the expected range for nickel-precharged cells.

SUBSEQUENT TESTING

After initial acceptance tests 6 of the 7 cells were shipped to the customer, with one asbestos cell being retained in

cold storage at Eagle Picker. At the customer's facility the cells were subjected to characterization tests with a nominal

rating of 100AH. The testing included discharges at 74.5 amps and 100 amps. Counting the initial acceptance and

conditioning testing, 51 full-depth cycles had been performed on the cells when, within a few cycles of the end of this

period, the cells declined in capacity from an initial value of about 114 AH to about 102 AH at 10°C The cells were returned

to Eagle Picher at this point, where the low capacity condition was confirmed. There was no evidence of a second plateau

in the discharge curve.

A destructive physical analysis (DPA) was performed on one of the single Zircar cells. The results showed that the

positive electrodes had expanded in thickness about 12% above their initial installed thickness. This condition is consistent

with the literature [3] [4] indicating a correlation of positive electrode swelling with high electrolyte concentration.

Analysis of the electrolyte in the cell components confirmed the high concentration at an average of 39.3 1 %. It showed

that the electrolyte distribution was uniform among the four quadrants of the stack, but that each separator had given up

about a third of its electrolyte to its corresponding positive electrode. It was therefore decided to remove the fill-tube

welds, open all of the cells, and reactivate them, this time fixing concentration at a lower value of 35%.

To replace the cell destroyed in the DPA, the cell with asbestos separatorthat had been stored (S!N 5-3) was added

to the group. Although this cell was never shipped to the customer and did not see the additional cycles, it was opened

and reactivated with the other five in order to achieve the same electrolyte concentration as the others.

Following reactivation the cells were rapid-cycled at 10°C using 10 amp charges for 16 hours and high rate (78 amp)

discharges, followed by a resistor-let-down to less than 0.01 volts. A chart of rapid-cycle performance is shown in figure

4. The initial performance was disappointing (about 92 AH), but with continued cycling the capacity rose steadily. After

24 full-depth cycles the capacities of the Zircar cells were again above 100 AH. Note that the discharge rate of 78 amps

42
was greater than the 68 amps used in the initial acceptance tests, and the charge rate was also greater by I amp. It can

be seen thatthe replacement asbestos cell (S/N 5-3) was initially higherthan the others and remained fairly level in capacity

until about cycle 12, whereupon it appeared to rejoin the other two asbestos cells in performance. This cell however

developed a leak due to the reopening process and had to be repaired. In the process of repairing a damaged fill tube, a

small piece of aluminum from a fixture fell into the cell. This may affect the future performance of this cell.

At cycle 16, due to small leaks in the activation plumbing on some of the cells, it was decided to reset the nickel

precharge, which in effect meant adding a small amount of hydrogen. This had a positive effect on capacities. Following

cycle 28, the cells were removed from the test fixtures and welded off. This involved another setting of nickel precharge

similar to the last. Following pinch-off the cells were subjected to one additional I 0°C rapid cycle, this time at the lower

discharge rate of 74.5 amps. The capacity jumped about 5 AH after pinch-off. It is probable that this is due partly to the

lower discharge rate (74.5 amps versus 78 amps) and possibly because capacities were somewhat hydrogen-limited prior

to that time due to plumbing leakage. (Pinch-off involves some hydrogen addition). The cells were then subjected to an

abbreviated acceptance test using 10 amp charges for 16 hours and 74.5 discharges, a lower discharge rate than for the

28 rapid cycles just completed.

Performance data for the 25 rapid cycles and the final mini-ATP are summarized in Table 4. It is clear that much

of the original high-capacity performance has been restored. Charge retention at 10°C was also confirmed at 86% to 88%.

PLANS

The cells are being returned to the customer and will be cycled to confirm the latest results. Vibration testing is also

planned. Life testing is then planned for up to 15 years and will be to a GEO regime (real-time eclipse cycling to an 80%

depth of discharge with shortened sun-times).

CONCLUSIONS

By using single-layer separator and slightly thicker, slightly more-porous positive electrodes, specific energies of

3.5 inch nickel-hydrogen cells can exceed 60 watt-hours per kilogram and provide good, all-around performance, even

at discharge rates of C/1.28. The use of single layer Zircar separator appears promising, but requires further validation

through life-testing.

The testing demonstrates that a 3.5 inch-diameter cell can be made to yield capacities above 110 AH . This is

important to spacecraft designers who are requiring larger-capacity batteries form any app] ications. If these higher-power

requirements can be satisfied by a 3.5 inch design, thermal characteristics will be better than with a larger-diameter cell.

It was also demonstrated that a swollen-positive-electrode condition can be corrected through application of

activation techniques. The long-term efficacy of this requires confirmation through cycle-testing. Such testing is planned

for these cells.

ACKNOWLEDGMENT

The professional assistance of Paul Daugherty in collecting and assessing data for this paper, and of David Cooke

for cell construction and testing, are gratefully acknowledged.

1. T. M. Yang, C. W. Koehler, A. Z. Applewhite, "An 83AH Nickel Hydrogen Battery for Geosynchronous Satellite

Applications," 24th IECEC Washington, D.C., Vol. 3, p. 1375

2. Jeffrey C. Brewer, Thomas H. Whitt, J. Roy Lanier, Jr., "Hubble Space Telescope Nickel-Hydrogen Batteries

Testing and Flight Performance," 26th IECEC, p. 257, Vol 3

3. J. J. Smithrick, and S. W. Hall, "Validation Test of Advanced Technology for IPV Nickel-Hydrogen Flight Cells

- Update," 27th Intersociety Energy Conversion Engineering Conference Proceedings, IECEC'92, San Diego,

CA, 1992, Vol. 1, p.1.215

4. H. S. Lim, and S. A. Verzwyvelt, "KOH Concentration Effect on Cycle Life ofNickel Hydrogen Cells, III. Cycle

Life Test," Journal of Power Sources, Vol. 22, Mar-Apr 1988, pp. 213-220

43
TYPE SEPARATOR

CELL

DESCRIPTION Single Single

Zircar Asbestos

Weight 2279 2236 (grams)*

Cell Length (in) 10.97 10.97 Dome to Dome

Cell Length (in) 11.56 11.56 Overall

* strain gage weight subtracted (17g)

TABLE 1

CAPACITIES TO

1.OV (AH)

AT 68 AMPS

SEPARATOR TYPE

Zircar Asbestos

30°C 81.3 82.9

20°C 99.1 99.7

10°C 113.9 111.3

0°C 117.6 113.1

-10°C 115.6 108.2

TABLE2

44
PERFORMANCE TYPE SEPARATOR

DATA Zircar Asbestos

180 Amp Pulse, 1.173 1.130 Minimum Voltages

WH/Kg 60.1 59.0

WH/Kg* 60.6 59.4

* Strain gage weight subtracted

TABLE 3

Capacities to 1.OV

(AH) at 74.5 Amps

(Post-Reactivation)

SEPARATOR TYPE

Zircar Asbestos

30°C 77.8 81.0

10°C 107.6 101.8

0°C 117.8 109.1

TABLE 4

45
FIGURE 1

100AH Cells

10°C Discharge Voltage

1.4

Rote: 68A with 180A 20 sec. pulse

1.3

4

------------ Single Asbestos

Single Zircar

V)

0

1.2

a.^

rnQ

Q 1.1

1

O 15 30 45 60 75 90 105

Minutes

FIGURE 2

46
65

60

Pf^ 55

cry

Z

50

Q

4 5

40

SPECIFIC ENERGY OF EP CELLS

WITH SLURRY SINTER POSITIVE PLATES

r^7 rr^ ^ O X17 O O

mc:i__ r ^ ^ a Z Z S S

pC p 1-C7 C7 ^ r- p

A M PERE-H O U RS (PRESS U RE)

FIGURE 3

100AH CYCLES AFTER REACTIVATION

CYCLES AT 10°C 1MTH 16 HR CHARGES, 78A DISCHARGES

113 — -- -- --

DISCHARGE RATE TO 74.5A

AND PINCH-OFF

108 -- -- - - _ - - -

103

O

=

98 a

Q ,.

93 r ----

RESETNIPRECHARGE

88

^ r— c— ^ ^— N N N N N C7

CYCLE NUMBER

FIGURE 4

S/N3-1

S/N3-3

• S/N3-4

o S/N 5-1

• S/N5-2

S/N5-3

47
HIGH SURFACE AREA, LOW WEIGHT COMPOSITE

NICKEL FIBER ELECTRODES

By Bradley A. Johnson, Richard E. Ferro, Greg M. Swain,

and Bruce J. Tatarchuk

Department of Chemical Engineering and

Space Power Institute

Auburn University, Alabama, 36849

ABSTRACT

The energy density and power density of light weight aerospace batteries utilizing

the nickel oxide electrode are often limited by the microstructures of both the collector

and the resulting active deposit in/on the collector. Heretofore, these two microstructures

have been intimately linked to one another by the materials used to prepare the collector

grid as well as the methods and conditions used to deposit the active material. Significant

weight and performance advantages have been demonstrated by Britton , and Re1d 2 at

NASA-LeRC using FIBREX nickel mats of ca. 28-32 μm diameter. Work in our

laboratory has investigated the potential performance advantages offered by nickel fiber

composite electrodes 3.a- 5 containing a mixture of fibers as small as 2 pm diameter

(Available from Memtec America Corporation). These electrode collectors possess in

excess of an order of magnitude more surface area per gram of collector than FIBREX

nickel. The increase in surface area of the collector roughly translates into an order of

magnitude thinner layer of active material. Performance data and advantages of these thin

layer structures will be presented. Attributes and limitations of their electrode

microstructure to independently control void volume, pore structure of the Ni(OH)2

deposition , and resulting electrical properties will be discussed.

INTRODUCTION

The secondary nickel-hydrogen (Ni/H 2 ) battery is a hybrid of the secondary Ni-Cd

battery and the H2 /02 fuel cell; a combination of the two best performing electrodes from

these two systems. The Ni/H2 battery possesses many advantages over other battery

systems which include (I) long cycle life, (2) stability to overcharge and (3) a built-in

state-of-charge indicator (cell pressure). However, the energy density and power density

of light weight aerospace batteries utilizing the nickel hydroxide electrode are often limited

because they must be designed with large void volumes to avoid clogging of the

conductive pathways with precipitant during electrochemical impregnation.

Previous research has concentrated on increasing the porosity of powder sinters

while attempting to maintain acceptable structural integrity and electrical conductivity.

PREC&NNG PAGE BLANK NOT FILMED 49
Other attempts have been to impregnate the powder sinter substrate more efficiently with

active material resulting in an increase of the electrochemical utilization. More recently,

work has involved the impregnation of FIBREX (Tm National Standards) nickel fibers

which has been shown to possess superior performance characteristics, especially in terms

of weight savings compared with plaque-based materials. Our laboratory has investigated

the potential performance advantages offered by nickel fiber composite electrodes 2. ' using

nickel fibers as small as 2 gm. The supposition was that the incorporation of small

diameter nickel fibers (2 gm) into FIBREX, or the use of the stand alone small diameter

fiber composite electrode, would result in a significant increase in the surface area

available for deposition of the active material without a significant reduction in void

volume. Superior performance characteristics of the nickel fiber composite electrode have

been demonstrated. By controlling the microstructure of both the collecting grid and the

active deposit, the performance of the traditional nickel hydroxide electrode has been

enhanced. Furthermore, the electrochemical impregnation at a reduced effective current

density results in a much thinner layer of active material, which is more electrochemically

accessible.

The key advantages of using small nickel fibers versus sintered nickel plaques or

FIBREX include an order of magnitude higher specific surface area per gram of fiber.

This not only provides increased area for the active deposit, but also increases

electrochemical accessibility to the active material. Additionally, there is a low ohmic

resistance within the microstructure due to the sinter bonded fibers and low mass transport

resistance within the microstructure voids resulting in easier transport of electrolytes.

Moreover, this electrode can be prepared with adjustable porosities and void volumes to

enhance inter- and intra-electrode transport while accommodating the required

precipitation of chemical products at the cathode. Finally, its electronic properties are

independent of any mechanical pressing.

EXPERIMENTAL METHODS

A. ELECTRODE STRUCTURE AND PREPARATION

Sintered fiber mats were produced using small nickel fibers and cellulose. The

nickel fibers had diameters of approximately 2 pm and were 2-3 mm in length. The

cellulose, a mixture of soft and hardwoods, was 20-30 pm in diameter and 100 to 1000

gm in length. These materials were then incorporated into a paper preform. This was

done by first thoroughly mixing the metal fibers in an industrial blender. They were then

agitated with the cellulose fibers until all fibers were uniformly dispersed. The solution of

dispersed fibers was then further diluted and formed into a 16 cm diameter paper preform

on a sheet mold. The preform was then pressed and allowed to dry in air.

Once the paper preforms were fabricated, circular discs ca. 19 mm were cut from

the preforms. Four of these discs were pressed together and placed on top of a 1 mil thick

nickel foil. This structure was then sintered at 1323 K for 30 minutes under a hydrogen

atmosphere. The sintering creates many intimate electrical contacts per fiber and provides

the mechanical stability required to withstand stress associated with swelling and

50

I
contraction of the active material during cycling. Additionally, during the sintering

process, virtually all of the cellulose gasifies, resulting in large pores and voids inside the

electrode microstructure.

The electrochemical impregnation of the Ni(OH) 2 active material was carried out

using the methods described by Britton'. The electrode was placed in an acidic solution of

l .5 M Ni(NO3 ) 2 , 0.175 M Co(NO3 ) 2 , and 0.075 M NaNO3 . Using current densities in the

range of 5-10 mA/cm 2 of geometrical surface area for 3-6 hours, a loading of 1.9 ± 0.4 g

Ni(OH)2/cc was achieved. The electrochemical reduction of nitrate produces hydroxide

ions in the vicinity of the bare metal fibers. Nickelous ions in solution combine with the

hydroxide ions to form nickel hydroxide, which precipitates out of solution onto the

surface of the fibrous network. After impregnation, the structure is rinsed with distilled

water, dried in an over at 393K for l hour, and stored under ambient conditions for a

minimum of 12 hours.

B . ELECTRODE EVALUATION

Cycle-life testing was performed in a half cell arrangement as described by

Britton'. A platinum gauze counter electrode was used and all potentials are reported

versus a saturated calomel electrode (SCE). All measurements were made at room

temperature in 26 wt% KOH. Cyclic voltametry and cyclic chronopotentiometry were

performed using a EG&G Princeton Applied Research Model 273

Potent iostat/Galvanostat. During cyclic chronopotentiometry, a current of C/2 was used,

while discharge current varied from C/2 to 3C. The C-Rate is defined as being the

theoretical rate at which all active material will be reduced in a period of 1 hour. From

these measurements, conditioning times, capacities, and efficiencies were determined.

Scanning electron micrographs were obtained with a JEOL 440 Microscope. Both

impregnated and non-impregnated electrodes were examined to determine the overall

effects of impregnation on the electrode microstructure. Furthermore, SEM was utilized

to confirm the sintering of metal contacts and to view the overall fiber distribution.

DICUSSION AND RESULTS

Figures l and 2 show SEM micrographs of the nickel fiber composite electrode

before and after impregnation, respectively. In the micrograph of the bare electrode, the

sintered contacts between the individual fibers along with the sizable void volume can be

seen (Fig. 1). The large number of sintered joints provides high conductivity pathways

and mechanical strength. The impregnated electrode possesses a relatively thin and

conformal coating over the entire electrode surface (Fig. 2). The active material possesses

many apparent defects and grain boundaries and exhibits regions of apparent high

crystallinity. The high surface area and large void volume of the electrode allows for a

considerable loading of active material while maintaining a sizable void volume.

Table l compares the physical properties of the electrode microstructures

examined during the course of this research before and after impregnation. The two

columns of data most noteworthy are the values for FIBREX and those for the nickel fiber

51

I
composite electrode. The composite electrode is thinner, possesses a 10 fold increase in

specific surface area and is slightly more porous than FIBREX. After impregnation, the

nickel fiber composite electrode possesses a higher weight of active material per unit void

volume, a higher volume percent loading and a thinner layer of active material compared

with FIBREX. Differences in surface area and active deposit thickness result from the

smaller diameter fibers used in the composite electrode. It is important to note that while

the electrodes possessing stainless steel fibers have a significantly larger surface area

than FIBREX, the performance data was poor due to the weak adhesion of the active

material to stainless steel. The incorporation of small diameter nickel fibers into FIBREX

improves the physical properties of the electrode, however, the nickel fiber composite

electrode possesses the most desirable properties.

Figure 3 shows typical charge and discharge curves for impregnated FIBREX and

nickel composite electrodes. The loading level per unit void volume is 1.52 g/cc and 1.69

g/cc, respectively. Again, the larger surface area of the composite electrode results in the

ability to deposit a considerable amount of active material while maintaining a sizable void

volume. In both cases, charging was performed at a C/2 rate while discharging was

performed at a 3C/2 rate. It can be seen that the potential during charging (oxidation) at

the two electrodes is ca. 0.8 and 0.5 V, respectively. The more positive potential plateau

at the FIBREX electrode is indicative of a higher ohmic resistance and/or mass transport

resistance and suggests that much of the current being passed is going toward oxygen

evolution, a parasitic side reaction which reduces the overall charge/discharge efficiency.

The more negative potential plateau observed at the nickel fiber composite electrode

reflects less ohmic loss and suggests that a higher percentage and the current passes is

going toward the oxidation of the active material to form the oxyhydroxide rather than

toward oxygen evolution, although it is likely that even at this potential oxygen evolution

is occurring to some degree. Upon reversal of the current to initiate the discharge

process, it can be seen that the electrode potential of FIBREX quickly drops and plateaus

at a more negative potential compared with the nickel fiber electrode. Again, this

difference is likely due to the increased ohmic loss within the FIBREX electrode. The

most noteworthy difference between the two electrodes is the increased capacity present

at the composite electrode. The percent utilization at the composite electrode is greater

than 200%. The maximum utilized capacity was usually reached within 5 to 10 cycles

which compares most favorably with the greater than 200 number of cycles often required

for many commercially available electrodes.

Table 2 contains the specifications used in the charge and discharge cycles. The

FIBREX-containing electrodes contain a larger amount of active material and

consequently, a higher rating (mA-hr). This active deposit is relatively thick compared

with that estimated for the nickel fiber composite electrode. The increased thickness

produces more mass transport impedance and results in a reduced efficiency. Most

importantly, the utilization typically observed at the FIBREX electrodes, containing small

metal fibers, is near 200%. However, a 300% utilization is routinely observed at the

nickel fiber composite electrode. A less efficient utilization is observed at the stainless

steel composite electrode because of the poor adhesion of the active material.. The active

material was observed to detach from off the electrode during cycling.

52
The effect of the discharge rate on the utilization was examined and is shown in

Figure 4. Data for a FIBREX and two nickel fiber composite electrodes are shown, one

with a low loading of 0.66 g/cc and one with a high loading of 1.69 g/cc. All charging

cycles were performed at C/2 and each data point was obtained after 5 cycles. The data

for the FIBREX electrode show the expected trend of decreasing utilization of the active

material with increasing rate. This decrease is due to the deficiency of protons within the

active material near the reaction sites and to the mass transport impedance through the

thick film. On the contrary, the data for the nickel fiber composite electrodes show an

increasing trend of utilization with discharge rate up to a rate near 2C where the utilization

begins to decrease. Higher utilizations are observed at the composite electrode with the

lower loading, as would be expected.

Figure 5 shows a plot of the percent utilization versus the cycle number for a

typical impregnated composite electrode. The plot shows how the conditioning time is

short at the composite electrode as full utilization is reached within the first 10 cycles.

The utilization is approaching 280% after 10 cycles and remains relatively constant for the

first 500 cycles. In each cycle, the charging step was performed at a C/2 rate and the

discharging was performed at a 3C/2 rate. After 500 cycles, the utilization begins to

decrease and by the time 1000 cycles have been performed, the utilization is near 150%.

The percent utilization is defined as the ratio of the charge passed during the reduction of

the active material (discharging), before the electrode potential decreases to -0.2 V vs.

SCE, to the charge passed during oxidation of the active material (charging).

Figure 6 lists some of the possible reactions which might be occurring to give rise

to the anomalous increase in capacity. At present, we have not identified the source of the

increased capacity but we are currently conducting a detailed mechanistic study of the

nickel hydroxide/oxyhydroxide redox process at the composite electrode. One possibility

involves the generation of hydrogen at the nickel electrode for extended periods of time

during the discharge cycle. The discharge cycle tests, described herein, were always

performed for a defined period of time regardless of the electrode potential value. In

many cases, the electrode potential decreased to ca. -1.2 V vs. SCE and remains there for

extended periods of time during the discharge cycle in order to maintain the desired

current flow. Consequently, significant quantities of hydrogen were produced. In this

same potential region, bare nickel substrate is exposed due to the reduction of nickel

hydroxide. The bare substrate provides sites for the dissociative chemisorption and or

oxidation of hydrogen which can then be injected into the active material adjacent to these

reaction sites. The injected protons may increase the ability to discharge residual or

isolated capacity by increasing the proton concentration which supports conductivity4.

The limiting step in the reduction of nickel oxyhydroxide to nickel hydroxide is known to

be the proton transfer. Increasing the availability of protons in the active material will

result in an increased utilization. The hydrogen gas generated can also be occluded within

the pores and defects of the active material and react chemically with isolated nickel

oxyhydroxide to reduce it back to nickel hydroxide as part of a self discharge

mechanisms-6

All of the cycle tests reported herein were conducted during exposure to the

atmosphere. CO2 will readily dissolve in alkaline solution to form CO 3 2 - anions. These

53
anions may affect the overall performance by intercalating and assisting in the

compensation of the excess charge introduced by the presence of Co l- and higher nickel

oxides7 . It is also possible that the CO 3 2 - anions can be reduced to a hydrocarbon species.

Another possible reaction which may contribute to the charge passed during discharge is

the reduction of carbon in the presence of excess protons to form methane. The carbon

forms on the electrode surface because of the incomplete gasification of the cellulose

during the sintering process. This is only speculation at this point as the potentials at

which these two reactions occur at a substantial rate have not been identified. Additional

sources for the extra capacity may also involve the chemical reaction of occluded oxygen

with the lower oxide to form the higher oxide. If this reaction were to occur at an

appreciable rate in competition with the electrochemical reduction during discharge, then

additional capacity could be obtained. This chemical reaction may have more of an

influence on the performance of the composite electrode given the thin and conformal

nature of the active material deposit. Also, oxygen reduction may occur concomitant with

the reduction of the higher oxide8 . Finally, some additional capacity will result from the

oxidation of the bare nickel substrate.

ACKNOWLEDGEMENTS

NASA-LeRC Electrochemical Technology Branch,

Contract # NAG3-1 154

Doris Britton

Patricia O'Donnell

Peggy Reid

54
REFERENCES

1. D.L. Britton, "Lightweight Fibrous Nickel Electrode for Nickel Hydrogen Battery",

NASA-TM-100958,1988.

2. D. Kohler, J. Zabasajja, A. Krishnagopalan, B. Tatarchuk, J. Electrochem. Soc., 137,

136 (1990).

3. D. Kohler, J. Zabasajja, F. Rose, B. Tatarchuk, J. Electrochem. Soc., 137, 1750,

(1990).

4. A.H. Zimmerman, in Proceedings of the Symposium on Nickel Hydroxide Electrodes,

eds. D.A. Corrigan and A.H. Zimmerman, Electrochemical Society, vol 90-4, 1980,

p. 311.

5. Z. Mao and R.E. White, J. Electrochem. Soc., 132, 1282 (1992).

6. A. Visintin, S. Srinivasan, A.J. Appleby and H.S. Lim, J. Electrochem. Soc., 13-2, 985

(1992).

7. C. Delmas, Y. Borthomiell, and C. Fauce, in Nickel Hydroxide Electrodes, eds. D.A.

Corrigan and A.H. Zimmerman, Electrochemical Society, vol. 90-4, 1980, p. 119.

8. D.T. Sawyer and L.V. Interrante, J. Electroanal. Chem., 2, 310 (1961).

55
Table 1. Physical properties of fiber based electrode microstructures

before and after impregnation.

FIBREX FIBREX/SS FIBREX /Ni Steel Nick.

BEFORE II(PREGRATION

Thickness 35 35 35 19 19

(mils)

Surf. Area/

Weight 160.5

(CM2/g)

Density

0.49

(g/cm3)

Porosity 94.4

(t)

292.8 276.7 2535 2297

0.55 0.55 0.33 0.35

93.7 93.7 95.8 95.8

Wt. of

Ni(OH) 2

0.42

/cm3 of

Void

Porosity 80.9

(t)

Loading 14.3

(Vol t)

Effective

Thickness 17.1

of Ni(OH)2

(μm)

AFTER IXPREGNATION

0.21 0.67 1.12 0.66

86.7 74.1 66.0 75.9

7.53 21.0 30.8 20.5

6.49 2.64 3.51 2.47

56
Table 2. Specifications used for the different electrode microstructures

during cycle testing.

FIBRE FIBREX+SS STEEL FJBREX+NI NICKEL

Weight 124.6 mg 138.7 mg 28.70 mg 185.0 mg 31 mg

Ni(OH 2

Weigh 85.32 mg 65.88 mg

Rating 24.67 19.04

(ula)

CHARGE (C/2)

Act. Cur. 12.33 mA 9.52 mA

Time (a) 7200 7200

Cur/wt. 0.144 0.144

(A/S)

DISC iARGE (3C/2)

Act Cur. 35.99 mA 28.56 mA

Time (a) 1200 4800

Cur/wL 0.433 0.433

63.80 mg 120.5 mg 44 mg

18.44 34.84 12.7

9.22 mA 17.42 mA 6.36mA

7200 7200 7200

0.144 0.144 0.144

27.66 mA 52.26 mA 19.1 mA

1120 4800 7440

0.433 0.433 0.433

Capa-hr

acity 45.3 38.08

(m

% Util. 50% 200%

8.61 69.68 39.43

36.67° 200° 310°

57
r•, .4

ZA

',

i_ >^.itt.`

IVAM I iftV 1

Figure 1. SEM micrograph of a bare nickel

fiber composite electrode.

Figure 2. SEM micrograph of an impregnated nickel

fiber composite material.

58
1.00

0.75

C/)

0.50

0 0.25

0.00

O

-0.25

- . 0.50 4

0

Charge (C/2)

EPI-FIBREX Discharge (3C/2)

1.52 g/cc

Nickel fiber electrode

1.69 g/cc

^ Nickel fiber electrode

EPI-FIBREX

2000 4000 6000 8000 10000 12000

Time (sec)

Figurr 3. Charge and discharge curves for impregnated Eagle-Picher (FIBREX

based) and nickel fiber composite electrodes 

in 26 wt% KOH.

350

300 - - ------- --- -

250 ..... ... ....... ....

. ... ............ .

200 . . ............ . ....... ...

.. ................

..................... . ... ....... .................................

X

150

Composite NHE (0.66 g/cc)

100 ............. -

Composite NHE (1.69 g/cc)

50 . . ..... -AE--

EPI-FIBR-EX (1,52 g/cc)

0

0.0 0.5 1.0 1.5 2.0 2.5 3.0 3.5

C-Rate

Figure 4. Plot of the percent utilization versus discharge rate for an

impregnated Eagle-Picher (FIBREX-based) and an 

impregnated

nickel fiber composite electrode at two different loading levels III

26 wt% KOH. Charging perfonned at C/2.

59
300

250

200

N

150

100

50

0 200 400 600 800 1000

Number of Cycles

Figure 5. Plot of the percent utilization versus the cycle number for an

impregnated nickel fiber composite electrode in 26 wt% KOH.

Charging perfonned at C/2 and discharging at 3C/2.

A. Chemical Reaction:

Lower Oxide + Oz/OH- H Higher Oxide

B. Growth and conditioning of active material formed on

the substrate.

C. Injection of H+/H generated during HER at sites on the

composite electrode into the active material.

Increased ability to discharge residual capacity by

raising proton concentration which supports

conductivity.

Influx of protons from remaining charged active

material or from Hz generated at current collector.

D. Cycle tests performed in open atmosphere

Intercalation of anions-charge neutralization within

active material.

Figure 6. Possible reaction schemes to account for the

"anomalous" capacity observed.

60
PA S 6 N94-23352

OVONIC NICKEL METAL HYDRIDE BATTERIES

FOR SPACE APPLICATIONS

S. Venkatesan, D. A. Corrigan, M. A. Fetcenko,

P. R. Gifford, S. K. Dhar, and S. R. Ovshinsky

OVONIC BATTERY COMPANY,

Troy, Michigan 48084

ABSTRACT

Ovonic nickel-metal hydride (NiMH) rechargeable batteries are easily adaptable to a variety of

applications. Small consumer NiMH cells have been developed and are now being manufactured by

licensees throughout the world. This technology has been successfully scaled up in larger prismatic

cells aimed at electric vehicle applications. Sealed cells aimed at satellite power applications have

also been built and cycle tested by OBC and other outside agencies. Prototype batteries with high

specific energy (over 80 Wh/kg), high energy density (245 Wh/L), and excellent power capability

(400 W/kg) have been produced. Ovonic NiMH batteries have demonstrated an excellent cycle life of

over 10,000 cycles at 30% DOD, Presently, Ovonic Battery Company is working on an advanced

version of this battery for space applications as part of an SBIR contract from NASA.

INTRODUCTION

Space power requirements include high gravimetric and volumetric energy densities, long cycle life,

and high reliability. Nickel-cadmium and nickel-hydrogen batteries are currently used for most space

applications. However, these systems do have several disadvantages. The environmental impact

related to the manufacture of nickel-cadmium batteries is coming under increasing scrutiny. The high

cost of nickel-hydrogen batteries makes it difficult to justify their use of many applications. Further,

improvements in gravimetric and volumetric energy densities for each system would be desirable.

The Ovonic nickel-metal hydride battery offers a viable alternative to the existing batteries with an

opportunity for significant gains in energy density [1-3]. They have roughly twice the gravimetric

and volumetric energy density of nickel-cadmium batteries. They are made with nontoxic environmentally acceptable materials. The viability of this technology has been demonstrated in consumer

cell applications with batteries now being manufactured and sold around the world through licensees.

Licensees of Ovonic Battery Company include Varta (Germany), Hitachi Maxell (Japan), Gold Peak

(Hong Kong), Samsung (Korea), Gates Energy Products and Harding Industries.

Considerable ongoing development is aimed at larger prismatic cells for electric vehicle applications

[1-3]. Last year, Ovonic Battery Company was awarded the first development contract by the United

States Advanced Battery Consortium, a consortium of the big three auto companies, DOE, and EPRI

to develop electric vehicle batteries.

In this paper, we describe the performance characteristics of Ovonic nickel-metal hydride cells as they

apply to space applications.

61
RESULTS AND DISCUSSION

Previously [4], the high energy density of Ovonic nickel-metal hydride technology was demonstrated

in small wound consumer cells. Gravimetric and volumetric energy densities of 70 Wh/kg and 210

Wh/L, respectively, were attained. More recently, even higher energy density was demonstrated in

larger prismatic cells. Results from a 50 Ali cell showing 80 Wh/kg and 245 Wh/L are

shown in Fig. 1. Presently, several proprietary approaches are being pursued at Ovonic Battery

Company aimed at a specific energy well in excess of 100 Wh/kg.

In collaboration with Eagle-Picher Industries [5], we have produced prismatic cells specifically

designed for satellite applications. Figure 2 shows a diagram of the OBC aerospace cells. These

cells delivered 5 Ah discharged at 5 A (C-rate), which corresponds to a specific energy of 55 Wh/kg.

The energy density delivered represents a two-fold improvement over the NiCd version of the cell in

Fig. 2 which is rated at 2.5 Ah. The next generation of cells being fabricated at OBC under the

ongoing NASA contract work will have a specific energy of 70 Wh/kg due primarily to improvements

in the positive electrode energy density.

Ovonic NiMH cells also exhibit excellent high rate capability. Previously, we have demonstrated up

to 400 W/kg in large prismatic cells aimed at electric vehicle applications [1-3]. The high rate

performance of OBC aerospace cells is shown in Fig. 3. The dependence of discharge capacity on

the discharge current up to 20 A (4-C rate) is shown. These cells deliver 80% of their rated capacity

at the 4-C discharge rate.

Ovonic NiMH have a wide operating temperature range as previously reported [1-3]. Figure 4 shows

the discharge performance of OBC aerospace cells over a range of temperatures. The discharge

performance is virtually identical between 10 and 30°C.

Ovonic NiMH cells have been tested by Eagle-Picher [5] and Rockwell International [6,7] to

determine cycle life under simulated LEO conditions (35% DOD). Rockwell International reported

having achieved over 10,000 cycles when the tests were terminated for other reasons [7]. From this

test, they projected a cycle life well over 17,000 cycles for Ovonic NiMH cells. In tests of Ovonic

aerospace cells, Eagle-Picher attained over 7,000 cycles to 37% DOD. Charge-discharge curves are

compared for cycles 89, 4900, and 7900 in Fig. 5.

Resistance to abusive overcharge and overdischarge conditions is a unique feature of the Ovonic

NiMH technology which should provide for improved reliability in actual operation in spacecraft

applications. Intrinsic overcharge protection is provided via an oxygen recombination cycle where

oxygen evolved at the positive electrode is recombined at the negative electrode. Intrinsic overdischarge protection is provided via a hydrogen recombination cycle where hydrogen evolved at the

positive electrode is recombined at the negative electrode. Previous results [4] have shown Ovonic

NiMH cells can be repeated overcharged and/or overdischarged by 15% or more for hundreds of

cycles without damage.

Ovonic nickel-metal hydride system is a viable alternative to provide for improved energy densities

beyond those currently provided by nickel-cadmium and nickel-hydrogen systems now in use. There

should be no sacrifice in cycle life or reliability. However, disadvantages such as environmental

concerns related to the NiCd system or high costs related to the nickel-hydrogen system can be

overcome.

62
REFERENCES

1. S. R. Ovshinsky, P. R. Gifford, S. Venkatesan, M. A. Fetcenko, D. A. Corrigan, and S. K.

Dhar, Proceedings of the 11th International Electric Vehicle Symposium, Florence, Italy,

September, 1992.

2. S. R. Ovshinsky, S. K. Dhar, S. Venkatesan, M. A. Fetcenko, P. R. Gifford, and D. A.

Corrigan, Proceedings of the 10th International Seminar on Primary and Secondary Battery

Technology and Applications, Deerfield Beach, Florida, March 1993.

3. S. Venkatesan, D. A. Corrigan, P. R. Gifford, M. A. Fetcenko, S. K. Dhar, and S. R.

Ovshinsky, Proceedings of the 8th Annual Battery Conference on Applications and Advances,

Long Beach, California, January, 1993.

4. M. A. Fetcenko, S. Venkatesan, S. K. Dhar, S. R. Ovshinsky, and T. Tsang, Proceedings of

the 3rd International Rechargeable Battery Seminar, Deerfield Beach, Florida, March 1992.

5. J. Brill, D. Coates, P. Bemis, S. Venkatesan, M. A. Fetcenko, S. R. Ovshinsky, AIAA/Utah

State University Conference on Small Satellites, August, 1990.

6. B. M. Otzinger, NASA Battery Workshop, Huntsville, Alabama, December, 1990.

7. B. M. Otzinger, Proceedings of the 6th Annual Battery Conference on Applications and

Advances, California State University, Long Beach, California, January, 1991; B. M.

Otzinger, Rockwell International, personal communication, January 1992.

63
I 

1.4

1.3

(5 1.2 Qt

1.1

O 1

0.9

08

0 5 10 15 20 25 30 35 40 45 50 55

CAPACITY (Ah)

Figure 1. Discharge voltage versus time curve for advanced

design 50 Ah EV prototype cell demonstrating 80

Wh/kg gravimetric energy density.

1

FIGURE 2

64
_ 1.3

(D 1.2

0)

_M 1.1

O

> 1

0.9

08

OVONIC BATTERY COMPANY

CAPACITY VS. DISCHARGE CURRENT

6

5

1L

0 2 4 6 8 10 12 14 16 18 20

Discharge Current (Amps)

FIGURE 3

OVONIC BATTERY COMPANY

VOLTAGE VS. TIME

AS A FUNCTION OF TEMPERATURE

.-CQ 4

U

0

0-3

cz

U

2

1.5

1.4

0 1 2 3

Time (Hrs.)

10G 22C 30C B ----A - - - -

........p ...

FIGURE 4

4 5 6

65
i ^'! H ; _F T TT AT I [3-'N'

CHu : 5"-)MIN@ DIS-35MIN@

i FF1l - nFIo 2 - f1F)R

w

CD .

CE 1:3 JO

1.2 WU

l.i

1.5

1.4

0.9

0

1.0

10 20 30 4.0 50 50 70 80 90 100

ll'. ^E (MINUTES)

FIGURE 5

66
Pysl.^ N94 :23353

ALLOYS FOR HYDROGEN STORAGE IN

NICKEL/HYDROGEN AND NICKEL/METAL HYDRIDE BATTERIES

Anaba Anani, Arnaldo Visintin, Konstantin Petrov and Supramaniam Srinivasan

Center for Electrochemical Systems and Hydrogen Research

Texas Engineering Experiment Station

Texas A&M University System

College Station, TX 77843-3402

James J. Reilly and John R. Johnson

Department of Applied Science

Brookhaven National Laboratory

Upton, NY 11973

Ricardo B. Schwarz and Paul B. Desch

Center for Materials Science

Los Alamos, NM 87545

ABSTRACT

Since 1990, there has been an ongoing collaboration among the authors in the three laboratories to (1)

prepare alloys of the A135 and AB2 types, using arc-melting/annealing and mechanical alloying/annealing techniques;

(2) examine their physico-chemical characteristics (morphology, composition); (3) determine the hydrogen

absorption/desorption behavior (pressure-composition isotherms as a function of temperature); and (4) evaluate their

performance characteristics as hydride electrodes (charge/discharge, capacity retention, cycle life, high rate capability).

This review article presents the work carried out on representative AB5 and AB2 type modified alloys (by partial

substitution or with small additives of other elements). The purpose of the modification was to optimize the

thermodynamics and kinetics of the hydriding/dehydriding reactions and enhance the stabilities of the alloys for the

desired battery applications. The results of our collaboration, to date, demonstrate that (1) alloys prepared by arcmelting/annealing and mechanical alloying/annealing techniques exhibit similar morphology, composition and

hydriding/dehydriding characteristics; (2) alloys with the appropriate small amounts of substituent or additive

elements: (i) retain the single phase structure, (ii) improve the hydriding/dehydriding reactions for the battery

applications, and (iii) enhance the stability in the battery environment; and (3) the AB2 type alloys exhibit higher

energy densities than the AB5 type alloys but the state-of-the-art, commercialized batteries are predominantly

manufactured using AB5 type alloys.

RATIONALE FOR INVESTIGATIONS AND CRITERIA FOR ALLOY SELECTION

The interest in alloys for hydrogen storage in nickel/hydrogen and, particularly, in nickel/metal hydride

batteries is gaining momentum because of the high performance characteristics of these batteries, in respect to energy

efficiency, energy density, high rate charge/discharge capability and cycle life. Previous studies at CESHR, TEESTAMUS have shown that the high rate of self-discharge in nickel/hydrogen batteries is due to the direct reduction of

the nickel oxyhydroxide to nickel oxide (1-4). In these batteries, hydrogen is stored as a compressed gas and because

of the high pressure of hydrogen the rate of transport to the nickel oxide electrode is quite rapid. In order to elucidate

the mechanism of the self-discharge reaction and find methods for its inhibition, microcalorimetric and

electrochemical techniques were used. These studies showed that the rate of self-discharge is proportional to the

hydrogen pressure in the battery and the activity of nickel oxyhydroxide. Several types of potential inhibitors were

incorporated in the nickel oxide electrode, during its fabrication or added to the electrolyte, to elucidate their effects on

the rate of self-discharge. Though some additives such as lead, cobalt, and cadmium were found to have some effect,

the rate of self-discharge could not be reduced significantly. Thus, this work showed that the only way in which the

self-discharge could be reduced was by lowering the pressure of hydrogen in the pressure vessel. This is not feasible

because it would increase the weight and volume of the nickel/hydrogen battery; thus, it was thought at that time

that one approach would be to use hydrides for hydrogen storage in nickel/hydrogen batteries. By using this

67
approach, the hydrogen pressure could be considerably reduced and hence the rate of self-discharge considerably

decreased. The added advantage is that hydrogen stored as hydrides requires a considerably smaller volume than

hydrogen stored as compressed gas; further, there could also be a better packing and, hence, an improvernem in the

thermal management of this battery.

During the last five years, there has been a great incentive for developing nickel/metal hydride batteries.

Use of nickel/metal hydride, rather than nickel/hydrogen batteries, can reduce the weight and thereby increase the

specific energy, but the main impetus for developing nickel/metal hydride batteries is due to the impending

legislations to curtail the production of nickel/cadmium batteries, because of the highly toxic nature of cadmium.

The nickel/metal hydride batteries which are being developed, manufactured and commercialized, particularly in

Japan, have a wide variety of applications: lap-top computers, cellular phones, video-cameras, etc. There is also

great interest in developing nickel/metal hydride batteries as power sources for electric vehicles. The advantages of

the nickel/metal hydride batteries are: (1) that their energy densities, in terms of weight or volume, are better than

those for the nickel/cadmiurn battery; (2) their rate capabilities are approaching that of the nickel/cadmium batteries;

and (3) the lifetimes of these batteries, which are still in the infant stage of development, are approaching at least

1,000 cycles (5,6).

For the aforementioned type of applications, there has been, in recent times, great interest in two types of

alloys for hydrogen storage - the AB5 type and the AB2 type alloys. Since the 1960s, research on these hydrides has

been fairly extensive at Brookhaven National Laboratory in New York (7) and Phillips Research Laboratory in

Eindhoven (8). The pioneering work on the AB2 type alloys have been carried out at Energy Conversion Devices - a

subsidiary of this company, Ovonics, has been formed recently and is presently heavily engaged in developing

nickel/metal hydride batteries, particularly for the electric vehicle application (9). In these alloys, the A component

is the one which forms the stable hydride. The B component performs several functions: (1) it could play a catalytic

role in enhancing the hydriding/dehydriding characteristics; (2) it can alter the equilibrium pressures for the hydrogen

absorption/desorption and raise it or lower it to the desired levels; and (3) it could also add to the stability of the

alloys because some of the A components could be readily oxidized. The AB5 type alloys have a hexagonal structure

while the AB2 type alloys have a cubic structure. Most of the alloys of both these types, which have been

successfully used or are being developed, are multi-component. In the AB5 type alloy, the parent alloy is LaNi5,

while in the AB2 type it is TiNi2 or ZrNi2. However, either the partial substitution by or a small addition of other

elements alters the hydriding/dehydriding characteristics of these alloys, as required for gas phase hydrogen storage

and for hydride electrodes. The roles of these substituents or additives could be to: (1) increase or decrease the crystal

lattice constant and thereby change the equilibrium pressure for hydrogen absorption/desorption; (2) decrease the

hysteresis during hydrogen absorption/desorption; (3) catalyze the hydriding/dehydriding reactions; and/or (4) improve

the stability of these alloys by preventing oxide formation on one or more of the components.

Since 1990, there has been an ongoing collaboration among the authors in the three institutions to: (1)

prepare alloys of the AB5 and AB2 types using arc-melting and mechanical alloying techniques; (2) examine their

physicochemical characteristics - morphology, composition; (3) determine the hydrogen absorption/desorption

behavior (pressure-composition isotherms as a function of temperature); and (4) evaluate their performance

characteristics as hydride electrodes (charge/discharge, capacity retention, cycle life, high rate capability). This review

article illustrates the results of the several types of investigations, which have been carried out on some selected A135

and A132 alloys (10-13). Inmost cases, the alloys were prepared by the arc-melting and annealing technique, but in

some cases also by the mechanical alloying and annealing technique. A comparison of the physicochemical

characteristics, including the hydriding/dehydriding reaction, was also made on the same alloy prepared by the two

methods.

PREPARATION AND PHYSICOCHEMICAL CHARACTERIZATION OF ALLOYS

The AB2 and AB5 alloys were prepared by arc-melting of the constituent elements followed by annealing in

an argon atmosphere ; for the AB2 type alloys, annealing was carried out at 800 0 C for three days and for the AB5

type alloys it was at 10000 C for 36 hours. The alloys were pulverized mechanically into powders about 100 mesh.

The method of mechanical alloying (14) was used to prepare some selected alloys of the AB5 type. Mechanical

alloying is a high energy ball-milling process that repeatedly cold-welds and fractures powder particles. The

microstructure that forms during mechanical alloying consists of layers of the starting material; the thickness of the

material decreases with increase in mechanical alloying time leading to true alloy formation down to atomic levels.

68
Phases formed are usually meta-stable both in structure and morphology and amorphousity (nano-crystallinity) is

easily obtained by controlling the alloying time. Alloys were prepared using powder elements of puritronic quality,

obtained from Johnson Mathey Company. A laboratory size Spex 8000 mixer/mill was used to mechanically alloy

mixtures of the constituent elements contained in a tungsten carbide vial. Alloying was carried out in a high purity

helium atmosphere glove box in powder charge batches of 5 grams each. Multiple mechanical alloying runs were

combined to obtain 20 to 30 gram sample powders. In our investigations, over thirty alloys have been prepared and

characterized. A representative list of the alloys, which were prepared using the arc-melting and/or mechanical

alloying technique, is found in Table 1.

PHYSICOCHEMICAL CHARACTERIZATION OF ALLOYS

Structural Characterization Using X-ray Diffraction Techniques

Structure and phase identity of the alloy samples were characterized by x-ray diffraction utilizing a Scintag

diffractometer with Cu-K a-radiation. Typical x-ray diffractograms for the alloys, prepared by arc-melting and

annealing and by mechanical alloying and annealing, are shown in Figures 1 and 2. These figures also illustrate that

the incorporation of small quantities of substituents and/or additives do not effect the single phase structure of the

alloys. It is also worth mentioning that if the annealing treatment is not carried out on the mechanically alloyed

samples, broad diffraction peaks are observed. However, sharp peaks are obtained by annealing because a nanocrystalline (amorphous) to crystalline transformation takes place. Depending on the temperature and time of

annealing, the entire amorphous state can be transformed to an entirely crystallized state. The lattice constants and

cell volumes for the selected alloys, as extracted from the x-ray data for the alloys, are tabulated in Table 1. A

striking result is that the cell volume for the AB2 type alloy is nearly twice that for the AB5 type. This is reflected

in the considerably lower equilibrium pressures for hydrogen absorption/desorption for the AB2 type alloys than for

the AB5 type alloys. The incorporation of substituents or additives can alter the lattice constants and cell volumes,

and, hence, the equilibrium pressures for hydrogen absorption/desorption.

Morphologic and Compositional Characterization Using Scanning Electron Microscopy -

Energy Dispersive X-ray Analysis

An analysis of the morphology, composition and distribution of the elements was made using SEM-EDAX

techniques. A JSM Model 6400 Scanning Electron Microscope (equipped with a Noran 1-2 EDX unit) was used for

the study. The SEM-EDAX studies confirm the x-ray diffraction studies in respect to morphology and the

homogeneity of the alloys. A typical SEM photograph, which illustrates the homogeneity of the alloys, is shown :n

Figure 3. The compositions of these alloys were very close to that of the starting composition and the

examinations, carried out at high resolution levels, revealed the homogeneity of the alloys, in respect to structure and

composition.

Evaluation of Alloys for Hydrogen Storage from Gas-Phase Pressure-Composition-Temperature

Isotherms

The pressure-composition-temperature isotherms for hydrogen absorption/desorption are generally measured

by using a modified Sievert's type apparatus. The principles of the hydrogen absorption/desorption reaction may be

represented by Figure 4. With increase of pressure, there will be an increase in the amount of hydrogen absorption

in the single-phase (a-phase) region. Hydrogen absorption is in the solid solution region and Sievert's Law is

obeyed, i.e. the amount of hydrogen absorption is proportional to the hydrogen pressure. Once the formation of the

single-phase alloy is completed, the second phase starts forming, (i.e. the P-phase) and this forms from the outer

surface and migrates towards the interior of the alloy. This is a two-phase region and for an ideal system there would

be plateau pressure in the pressure-composition isotherm. Once the R-phase is completely formed, if any more

hydrogen is to be introduced, this will again have to be with a significant increase of pressure. During desorption

from the 0-phase solid solution region, a decrease of pressure causes a decrease of the hydrogen content. The

hydrogen desorption starts from the outer surface giving rise to the a-phase formation and this case is exactly the

reverse during the absorption phenomenon (Figure 4). For an ideal case, there will be very little hysteresis in the

hydrogen absorption and desorption isotherms, but in general there will be a slightly lower pressure in the two-phase

region during desorption. Once the (3-phase has been completely transformed into the a-phase (the solid solution

region), with decrease of pressure there will be a decrease in the hydrogen content. The AB5 type of alloy generally

69
exhibits a plateau pressure (P) that would be determined by the thermodynamics of the hydrogen

absorption/desorption reactions. With increasing temperature, the plateau pressures will increase and the capacity for

hydrogen absorption will decrease in this two phase region. Using the Vant-Hoff equation,

InP= A-H— - IS- (1)

RT R

one could calculate the enthalpy (AH) for the absorption/desorpdon reactions as well as the entropy change (AS) for

the reaction. The enthalpy change is generally of the order of about 10 kcal/mole.

With the AB2 type alloys, plateau pressures are hardly observed, which could be due to some multi-phase

alloy formation in the intermediate range. However, the pressures during which the absorption/desorption occur

could be considerably below one atmosphere. For hydrogen storage this will be a disadvantage, but for the hydride

electrodes, even though there is a considerable slope in the intermediate region of the pressure-composition isotherm,

this is not a disadvantage, because at a sacrifice of, say, less than 40 or 50 mV, the coulombic capacity for the

hydriding/dehydriding reaction corresponds close to the maximum hydrogen content (i.e., one atom of hydrogen for

each atom of the metal). In the Sievert's type experiments, one can calculate the amount of hydrogen (H) absorbed

by using the equation:

(H)= 2 uAP -Yc= 21) At cp - Ec (2)

where R and T have their usual meanings, v is the volume of reservoir, cp = 0.01098 is the slope of the pressure vs.

transducer reading curve, At is the change in transducer reading, and Ic is a correction factor accounting for hydrogen

in the gaseous phase in equilibrium with the solid hydride phase.

The useful hydrogen content for hydrogen storage, say, from the gas phase, can be considered to be in the

region between pressures of 0.5 and 5 atmospheres. In practice, for the AB5 type alloys, the range of hydrogen

pressures to obtain the maximum capacity for hydrogen absorption could be even smaller. The degree of slant may

be expressed by the equation:

S = log ( (3)

PH/M = 0.25

where PH/M = 0.75 and PH/M = 0.25 are the pressures at H/metal ratios corresponding to 0.75 and 0.25,

respectively. An optimum hydrogen storage alloy is defined in these terms to be one with a high hydrogen capacity

and a low degree of slant.

Typical hydrogen absorp6on/desorption isotherms for an alloy prepared at Brookhaven National Laboratory

(BNL 3) are shown in Figure 5. This figure illustrates a reasonably low plateau pressure and minimum hysteresis.

Further, this alloy also shows a considerable wide range over which hydrogen is absorbed and it approximately

corresponds to one atom of hydrogen for each atom of metal. Table 1 summarizes the results of the hydrogen

absorption/desorption characteristics of some alloys, prepared using both the arc-welding and the mechanical alloying

techniques. Figure 6 shows the hydrogen absorption behavior for the alloy prepared by the mechanical alloying

technique before and after annealing. Due to the amorphous nature of the alloy prior to annealing, the hydrogen

absorption/desorption characteristics are not satisfactory. However, the annealing causes a transformation from the

amorphous to the crystalline structure and produces grain boundaries, which are necessary for the migration of the

hydrogen into the alloy; thus, the absorption/desorp6on characteristics are considerably improved. Figure 7 shows a

comparison of the hydrogen absorption isotherms for the same alloy prepared by the mechanical alloying and by the

arc-melting techniques. One can note from this figure that the characteristics are very similar, in respect to the

plateau pressures; however, the alloy prepared by mechanical alloying showed a lower hydrogen content. This could,

however, be improved very probably by extending the time for mechanical alloying. The incorporation of tin

improves the hydrogen absorption/desorption behavior and cycle life (15). It also has some effect of lowering the

plateau pressure. In addition, the substitution of neodymium increases the crystal lattice parameters and also lowers

the plateau pressure (8).

70
EVALUATION OF ALLOYS AS HYDRIDE ELECTRODES

Charge/Discharge and Cycle Life Behavior

Reactions which occur during the charging and discharging of hydride electrodes in alkaline electrolytes may

be represented as follows:

M + H2O + e a M—Hads + OH int (4)

M - Hads a MH (5)

OH int + OH bulk (6)

MH MH (7)

2 Hads H2 T (8)

M + H2O + e- a MH + OH (9)

The first step (Eq. 4) is the charge transfer reaction, which gives rise to the formation of the adsorbed

hydrogen and hydroxide ions at the interface. The second step (Eq. 5) involves the dissolution of H in the alloy and

the diffusion of the absorbed hydrogen into the bulk alloy, while the hydroxide ions diffuse into the bulk of the

electrolyte (Eq. 6). The third step (Eq. 7) is the recrystallization of the condensed phase. During overcharging, one

could have hydrogen evolution (Eq. 8), and for practical purposes, this has to be avoided in the nickel/metal hydride

batteries. The complete electrode reaction is represented by Eq. 9. The phase changes, which occur during this

reaction, are identical in the gas-phase hydrogen storage. The hydride electrodes for this work were generally prepared

by mixing the powder with 100 micron particle size copper powder and pressing into pellets onto nickel screens.

Nickel wires were spot-welded onto the electrodes for current collection. In more recent work, instead of using

copper powder, high surface area carbon with Teflon was used for the fabrication of the electrodes. The main purpose

of the copper or carbon additive is to enhance the electronic conductivity of the electrode because during hydriding the

electronic conductivity decreases. The Teflon serves the purpose of a binder. The typical charge/discharge behavior

of the electrode, prepared with the BNL 3 material, is shown in Figure 8. This material exhibited very favorable

characteristics. In terms of cycle life, a limited number of studies have been carried out to date, but one example is

given in Figure 9 to show the advantages of using high surface area carbon, rather than copper, as an additive (16).

Very probably, the advantage of carbon over copper is that copper could undergo some dissolution in the strong

alkaline environment, whereas carbon is quite stable. There is also an enhanced capacity when using carbon as an

additive. This could be connected with the very high surface area of carbon (BET surface area of about 100 or 200

M2/g)_

Charge/Discharge Rate Capabilities

For several applications, it is desirable for the hydride electrodes to be able to function at relatively high

charge/discharge rates. A study was conducted on the LANL-3 material to demonstrate the performance as a function

of the discharge rate (Figure 10). This figure shows that even at the 3C rate of charge that about 90 010 of the

capacity could be obtained without any significant loss in the potential of the electrode. Such studies have been

conducted on several of the electrodes with the other alloys listed in Table 1 and one may conclude that fast discharge

rates, nearly as high as those in the nickel/cadmium battery, can be obtained with these materials. Of the materials

tested to date, the one which has shown the best behavior is the BNL 4 material which is of the AB2 type.

Capacity Retention

The capacity retention is an essential criterion for secondary batteries; this parameter gives a measure of the

self-discharge rate of the batteries. As mentioned in the first section, the self-discharge rate on the nickel/hydrogen

battery is quite high, about 10% per day. This value is about 1% per day for the nickel/cadmium battery. Capacity

retention studies were conducted by fully charging the battery and leaving it for different periods of storage time.

71
After the desired periods, the remaining capacities were determined by completely discharging the batteries. Typical

plots of the capacity retention and of the electrode potential, as a function of storage time, for the PRL I material are

shown in Figure 11. Figure 12 shows the capacity retention as a function of storage time for several of the alloys

tested in this work. An intuitive guess is that there should be a relationship between the dissolution pressure and the

self-discharge rate. Table 2 demonstrates the correlation between the desorption pressures and the self-discharge rate

after five days for some of the alloys. Here again, the AB2 type alloys may have some advantages over AB5 type

ones because of their considerably lower plateau pressures, but as seen from Table 1, it is clear that even the AB5

type alloys can be modified with substituents and/or additives to lower the plateau pressure to a level of about 0.1 to

1.0 atm. From this article, one may conclude that hydride electrodes could function optimally with alloys having

plateau pressures or even sloping pressures in the range of .1 and .9 atm.

CONCLUSIONS

1. Modified alloys, prepared by partial substitution of the A and B components and by incorporation of

minimal amounts of additives, retain the single phase structure, and provide the necessary changes in the

physicochemical characteristics to attain high hydrogen storage capacities at relatively low equilibrium pressures.

2. Alloys prepared by the two methods (arc-melting/annealing and ball milling/annealing) yield very similar

hydrogen absorption/desorption characteristics both in the gas phase and as hydride electrodes. The alloy prepared by

mechanical alloying attained the maximum electrochemical capacity without prior activation.

3. The AB5 type alloys, unlike the AB2 type alloys, exhibit minimal hysteresis and degree of slant of the

pressure-composition isotherm for hydrogen absorption/desorption. Thus, for storage of hydrogen from the gas

phase (as in Ni/H2 batteries), the AB5 type alloys have a distinct advantage.

4. The AB2 type alloys absorb/desorb hydrogen at lower pressures and have higher hydrogen storage

capacities that the AB5 type alloys and hence show a higher level of performance as hydride electrodes.

5. A modified LaNi5 alloy incorporating Ce or Nd and Sn exhibits excellent hydrogen storage behavior

from the gas phase and also as a hydride electrode.

ACKNOWLEDGEMENTS

This work was sponsored by the Office of Research and Development under contract No. 88*F674300 and

the NASA Center for Space Power at Texas A&M University (Contract P.O. 58-529211-X113). Special thanks are

also due to Dr. Hong S. Lim of Hughes Aircraft Company and Mr. Joe Stockel of the U. S. Government Office of

Research and Development for their encouragement and discussions of this work.

REFERENCES

1. Y.J. Kim, S. Srinivasan and A. J. Appleby, J. Appl. Electrochemistry , 20 (1990) 377.

2. K. A. Murugesamoorthi, S. Srinivasan and A. J. Appleby, J. Appl. Electrochemistry , 21 (1991) 95.

3. YJ. Kim, A. Visintin, S. Srinivasan and A. J. Appleby, J. Appl. Electrochemistry, 139 (1991), 351.

4. Z. Mao, A. Visintin, H. S. Lim and A. J. Appleby, J. Appl. Electrochemistry, 22, (1992) 409.

5. T. Sakai, K. Muta, H. Miyamura, N. Kuriyama, and H. Ishikawa, in D. A. Corrigan and S. Srinivasan (Ms),

Proceedings of the Symposium on "Hydrogen Storage Materials, Batteries and Electrochemistry", The

Electrochemical Society, Pennington, NJ, 92-5 (1992), 59.

6. I. Matsumoto, M. Ikoma, A. Ohata, H. Matsuda and Y. Toyoguchi, Extended Abstract of 183rd Meeting of the

Electrochemical Society, The Electrochemical Society, Princeton, NJ, 93-1, (1993) 39.

72
7. J. J. Reilly, in D. A. Corrigan and S. Srinivasan (eds) Proceedings of the Symposium on "Hydrogen Storage

Materials, Batteries and Electrochemistry", The Electrochemical Society, Pennington, NJ, 92-5 (1992), 24.

8. H. H. Van Mal, K. H. J. Buschow, and A. R. Miedema, J. Less Common Metals, 35, (1974), 65.

9. M. A. Fetcenko, S. Venkatesan and S.R. Orkinsky,in D. A. Corrigan and S. Srinivasan (eds), Proceedings of the

Symposium on "Hydrogen Storage Materials, Batteries and Electrochemistry", The Electrochemical

Society, Pennington, NJ, 92-5 (1992), 76.

10. A. Anani, A. Visintin, S. Srinivasan and A. J. Appleby, in D. A. Corrigan and S. Srinivasan (eds), Proceedings

of the Symposium on "Hydrogen Storage Materials, Batteries and Electrochemistry", The Electrochemical

Society, Pennington, NJ, 92-5 (1992), 105.

11. A. Anani, A. Visintin, S. Srinivasan and A. J. Appleby, Extended Abstract of 183rd Meeting of the

Electrochemical Society, The Electrochemical Society, Princeton, NJ, 93-1 (1993) 82.

12. A. Visintin, K. Petrov, S. Srinivasan A. J. Appleby J. J. Reilly, J. R. Johnson and

H. S. Lim, Extended Abstract of 183rd Meeting of the Electrochemical Society, The Electrochemical

Society, Princeton, NJ, 93-1, (1993) 79.

13. A. Anani, A. Visintin, S. Srinivasan, P. B. Desch and R. B. Schwarz, to be submitted to the Journal of the

Electrochemical Society.

14.R. B. Schwarz, R. R. Petrich and C. K. Saw, J. Non-Crystalline Solids, 76, (1985) 281.

15. D. Chandra and F. E. Lynch, Rare Earths, Extraction, Preparation and Applications, ed. R. G. Bautista and M.

M. Wonf. The Minerals, Metals and Materials Society, (1989), 83-89.

16. K. Petrov, A. Visintin, S. Srinivasan and A. J. Appleby, Extended Abstract of 183rd Meeting of the

Electrochemical Society, The Electrochemical Society, Princeton, NJ (1993) 41.

73
TABLE 1

Crystallographic Parameters, Hydrogen Storage Characteristics

and Coulomhic Capacities of Electrodes for Some Selected

AB5 and AB2 Alloys

COMPOSITION I D Lattice Const.

a/c (AO)

Cell

Volume

(Ao)3

Plateau

Pressure

(atm)

H? Storage

H/atom

Pressure Range

0.5 - 5.0 atm

H?

Capacity

(mAh/g)

LaNi S LANL-1 5.009/3.97 86.31 1.6 0.916 341

MmNi Mn A Co MBC-1 4.977/4.05 86.99 0.8-1.5 0.833 316

Lao gNdnNi Co? 4 SiO I PRL-1 5.020/3.98 87.08 0.96 0.721 273

Lao gNdO Ni Co? 4 Mon j Si O PRL-2 5.028/3.98 87.35

LaNi Co Si BNL-1 5.012/4.00 87.05 2.4 0.99 354

L Cc Ni Co Mo Sin 1 BNL-2 5.004/4.01 86.95 0.6 0.716 207

Lan X Ce.() N i4 gSn BNL-3 1.08 0.933 335

Lao gNdoNi Co Sn LANL-2 5.042/4.02 88.48 0.57 0.75 257

Lao gNdpNi Con S Sn() BNL-8 5.044/4.02 88.76 0.6 0.792 272

L Nd Ni4 SCon 5SnO LANL-3 5.044/4.02 88.56 0.8 0.75 259

Tin 5 j zr V Ni Cr BNL4 4.959/8.06 172.3 0.82 350

Table 2

Relationship between Desorption Pressures and Self-Discharge

for PRL-1, BNL-1, 2, and 3 Alloys

Alloy Identification Desorption pressure

(atm)

Self-Discharge

after 5 days

BNL-2 0.61 10

BNL-3 0.96 12

PRL-1 0.96 15

BNL-1 2 40

74
ry

--r

rn

n

E

CD

cn

COUNTS

.,n nn 9?_31 184.62 276.92 369.23 461.54 553.8

Figure 1.—X-ray diffraction pattern for BNL-1 alloy, prepared by arc-melting and annealing.

Numbers above peaks represent d-spacing in A°.

75
- --

.1

.6

co \ \

.1

^

^+

\ N

6 a a \ n

a ♦ N

N w .. .. \

6-

N

^

1 ^i t* ^

m

CU

(a) Prepared by mechanical alloying.

CPS 884 Q 56 1 54 1 14 %

485• lee

n

chi

436. 90

388. m 8e

^' a

339. .^ a

N

a

` ^ 78i

\ f f \

w

gid^ii 291. ^ \ OD ^• \ 1 ^ ^., \ 68

a

„

242. ^'

..

. • ! 5e

194.e145.5- 30

^ N \ \

n ..

'

_ \

CD

\ \

Q

97. \

..

..

..

4S. C 10

(b) Prepared by mechanical alloying and annealing.

Figure 2.—X-ray diffraction pattern of LaNi 5 alloy.

CPS

49

41

36

-33

V

E2

16

11,

5

ea

9e

80

70

60

5e

4e

3e

ee

1e

76
Figure 3.—Scanning electron micrographs for BNL-1 alloy, prepared by

arc-melting.

77
U)

o_

M .--s (3 - (M H) H

I I

a- M H (M H)

[H] in Solids

Solid Solution : Sievert's Law

PH2 = f([HI)

Mixed Phase : van't Hoff Relation

PH, # f([H])

Figure 4.—Principles of hydrogen gas-phase absorption/desorption.

78
10

8

P

ci 6

U)

w

4

2

0

0 1 2 3 4 5 6 7

H / Mole of Alloy

Figure 5.—Typical gas phase absorption/desorption isotherms showing

minimal hysteresis effect. BNL-3; 25 °C open circle denotes

desorption; solid circle denotes absorption.

8

6

E

^n

4

a^

n

2

0 191 i i i i i

0 1 2 3 4 5 6

H/mol

Figure 6.--Comparison of hydrogen absorption isotherms at 25 °C

for the AB5 type material. Open circle denotes arc-melted

BNL-8. Square denotes mechanically alloyed and annealed

LANL-2.

79
4

2

0

'

1 O 1 CYCLE p

E V 2 CYCLE ;/

f 0 6 CYCLE s

1

Ll

Ii

f

0 

10

w

8

Na

6

k

0.4 0.6 0.8 1.0 1.2 1.4 1.6

H/MOL

(a) Prepared over a 20 hr period using tungsten carbide vial.

E 14

- O ABSORPTION Q

12

v DESOPTION

^ 

C

En

a 10

8

6

4

2

0 1 2 3 4 5 6

H/MOL

(b) Annealed at 700 °C for 15 min.

Figure 7.--Gas-phase absorption/desorption isotherms for LaNis,

prepared by mechanical alloying.

80
0

10MA 0 5 m •

o •

o •

• o

0 0 •

o •

• 0 •

Charge( Discharge I•o

• o CfP•

oD *0 •

0 0 p • • • • • • • • •

^ 0 p 0 0' •• 0 0 0 0 at

0 0 00

Charge I Discharge

-0.3

O

z

-0.5

-0.7

c^

c

O

O

-0.9

LO

U

O

UJ -1 .1

-5 0 5 10 15 20 25 30

Time (Hrs.)

Figure 8.--Charge/discharge behavior of the electrode prepared with the BNL 3 material.

300

M 250

00

200

p4

150

4

U

100

50

0 10 20 30 40 50 60 70 80 90 100

NUMBER OF CYCLES

Figure 9.—Effect of additive to active material (LANL-3) on the cycle life

of electrode. Open circle denotes LANL-3 + Teflonized acetylene black;

solid circle denotes LANL-3 + Cu powder; 25 °C; 8 M KOH.

81
0

900

00

v

x

q

m

0

a

0

b

0

c.

U

..rm

W

800

700

600

1000

25 C; 8 M KOH

3C 1C 1/2C

80mA/cmE

20mA/cmE lOmA/cin

500

400

0 30 60 90 120 150

Time, min.

Figure 10.—Electrode potential versus time during discharge at different rates for electrodes with

LANL-3 hydride material. Additive–Vulcan + 40% PTFE.

82
10 20 30 40

Storage Time (Days)

C

80 0

c

0

tz

60

U

C^

Q

cl^

U

-+40

50

-0.925

-0.950.+

0

cE

0

O

n

U

0

00

W

300 60

Q

E 250

U ^

Q 200

cl j

U

150

F

E

^ 100

0 10 20 30 40

Storage Time (Days)

(a) Capacity retention.

rO -0.900

0

0

40

c^

U

20 L

U

U)

-+0

50

100

00

(b) Electrode potential and capacity retention.

Figure 11.—PRL-1 sample as a function of storage time.

83
0- 0

BNL-2

PRL- 1

Electrolyte :31% KOH

Temperature : 25 °C

r0

C

a^

a^

^0

Cz

Q

cts U

80

60

BNL-1

100

40

0 4 8 12

Storage Time (Days)

Figure 12.--Comparison of the capacity retention capabilities for the PRL-1,

BNL-1, and BNL-2 alloy electrode samples.

84
439019Ps/Z * ^

N94- 2 ^)4

STATUS OF THE DEVELOPMENT OF RECHARGEABLE LITHIUM CELLS

G. Halpert, S. Surampudi, D. Shen, C-K Huang, S. Narayanan, E. Vamos and D. Perrone

Jet Propulsion Laboratory

California Institute of Technology

Pasadena, California 91109

ABSTRACT

The progress in the development of the ambient temperature Lithium - Titanium Disulfide

rechargeable cell under development at the Jet Propulsion Laboratory is described in this paper.

Originally aimed at achieving a specific energy of 100 V;Ug, 'AA' cells have demonstrated 125

W Ag at the C/3 discharge rate. The results of evaluating cell design parameters are discussed

and cycling test data are also included in the paper. Safety tests results at various overcharge and

over discharge conditions and rates proved to be uneventful. The test results of cell with built-in

overcharge mechanism proved the concept was feasible. Replacing the lithium foil electrode with a

LiX C resulted in a capacity at 1mA/cm 2 of 200 mAh/gm and 235 mAh/gm at 0.167 mA.

INTRODUCTION

Ambient temperature Lithium - Titanium Disulfide (Li-TiS2) rechargeable cells have been under

development at the Jet Propulsion Laboratory. The effort sponsored by Code C at NASA

Headquarters is aimed at high specific energy rechargeable cells that can reduce mass, and/or

volume and or increase mission capability in planetary spacecraft applications. The need for a

long cycle life lithium rechargeable battery has also been a goal of this effort. The program goals

were 100 WUg, 1000 cycles at 50% depth of discharge at the C/5 rate, 5 years of active storage

life and safe under electrical abuse conditions.

The approach has been to prepare materials, components and cells to gain an understanding of how

the variations in design, composition and structure affect performance and life. During the

development a number of cathode materials were studied in a comprehensive evaluation process

before selecting titanium disulfide for its specific energy and cycle life capability. More than

seventy-five different and solvent combination electrolytes were subjected to complex impedance,

microcalorimetry and electrochemical measurements before selecting a solution of 2-methyl

tetrahydrofuran/ethylene carbonate/2-methyl furan (2-McTHF/EC/2-MeF) containing 1.5 moles of

lthium hexafluoroarsenate (LiAsF6). The initial evaluation cells were a 150 mAh spiral wound

design . These were scaled-up to 1 Ali spiral wound cells. The volume of electrolyte was

found to be essential to long life.

85
The objectives of this effort were to determine the effect of stack compression, overcharge

protection and electrical abuse on the performance and safety of the Li-TiS 2 cell. Additionally an

alternate LixC electrode prepared from commercial graphite was evaluated as a replacement for

Lithium foil as the anode.

PERFORMANCE

Examples of the performance of these cells is given in Figures 14. Figure 1 shows the voltage on

charge (C/10) and discharge (C/5). Charge is terminated at 2.5 volts and discharge at 1.7 volts.

The sloping voltage has the advantage of providing a measure of state of charge. Figure 2

provides a measure of the specific energy at various discharge rates. Note that the 1 Ali

achieved a specific energy of greater than 125 Wh/Kg at rates as high as C/2. Figure 3 illustrates

965 cycles achieved at 50% depth of discharge (to 2.2 volts) in a cell comprising the 2-

McTHF/EC/2-McF/LiAsF6 electrolyte and a cathode:anode ratio of 4:1. The cycle life at 100%

depth of discharge (to 1.7 volts) shown in Figure 4 was 335 cycles. This is not significantly less

than that achieved with commercial Ni-Cd cells.

CELL DESIGN PARAMETERS

The parameters addressed during the development process this year included: pack tightness, final

selection of electrolyte composition, separator type, electrode capacity ratio and case polarity.

More than 100 experimental 1 Ali wrapped cells were assembled in house and 50 'AA' (I Ali)

cells were procured from EIC to determine the effect of these parameters on performance.

The effect of varying the pack tightness on end of discharge voltage is given in Figure 5. The

pack was tightened using a 10 mil Teflon sheet wrapped around the spiral. The tight pack utilized

2 layers around the spiral, the medium 1 layer and the loose pack did not use a Teflon layer.

Although there is a significant difference in the early cycles, the performance appears to be similar

after 250 cycles, probably because of the expansion of the cathode and film on the lithium foil.

However, the tight pack voltage was higher throughout and therefore was judged to be a superior

configuration.

Results of the evaluation of four different separators is shown in Figure 6. The Celgard

microporous polyethylene (PE) and polypropylene (PP) appear to provide similar results. The

microporous polypropylene was selected to minimize the lithium dendrite growth. The thermal

separator, a wax -coated polypropylene material, was selected to evaluate the properties of a

separator that would serve as a fuse and shut down the cell if the temperature became unsafe.

These also provided more than 400 cycles at 50% DOD.

The summary of the design parameter study, including work previously performed is given in

Table 1. The anode:cathode ratio of 4:1 has been shown to provide consistent long cycle life cells.

Three electrolytes including that described above have proven to capable of providing long life.

Only 6 cc of the electrolyte were found to be necessary. The optimum ethylene/propylene/diene

ter-polymer (EPDM) binder in the TiS2 cathode was found to be 1%. There was little difference

in performance with a floating or case positive polarity.

86
SAFETY EVALUATION TESTS

Experimental and EIC 1 Ali were subjected to short, circuit, overcharge, over discharge, and

combinations thereof The results of the short circuit test is given in Figure 7. Peak current rose to

25 amps within 2 minutes and then dropped to near zero current immediately after. The

temperature followed the current, rising to 220 0F, then tapering to 95 0 F within 30 minutes.

There were no ventings, explosions or other safety incidents during these tests. Table 2 provides

a summary of the results of inadvertent electrical abuse tests which occurred during cell testing

when a computer failed. Even after overcharging 5 times and over discharging 10 times the

capacity after 200 - 500 cycles, the cells were in tact and did not vent.

OVERCHARGE PROTECTION STUDIES

From an operational point of view cells in a series string require overcharge protection to prevent

a cell from being overcharged when degradation results in cell imbalance in the string. Ni-Cds

cells utilize their inherent oxygen recombination capability to allow cells in a string to be balanced

during overcharge. Lithium rechargeable cells do not inherently have this capability. Three

additives were included in Li-TiS2 cells to create a shuttle reaction that could provide overcharge

protection. Tetracyanoethylene (TCNE), N-butylferrocene, suggested by EIC and JPL's

tetramethylphenylenediamine (TMPD) were evaluated. Figure 8 shows the result of adding

TMPD to a cell. The additive provided 10-15% capacity at 2.9 volts thus protecting the cell

during that period. TCNE did not provide adequate protection. N-butylferrocene from, early

measurements appears to be comparable to or better than TMPD.

LITHIUM/CARBON ANODES

Replacement of lithium foil electrodes with an alternative lithium anode to enhance safety has been

of interest. The performance targets for this material were 250 mAh/g and >2mA/cm 2 with good

mechanical stability. One of areas JPL's effort has centered on is the development of a

lithium/carbon alloy (LixQ. Several carbon materials, including pitch coke, petroleum coke,

graphite fiber and graphite were evaluated as potential candidates. The rate capability of Li x C

using the graphite material is shown in Table 3. At 1 mA/cm 2 a capacity of 200 n"g was

achieved. At the lower rates, higher capacities to 235 mAh/g were found. A discharge curve is

given in Figure 9. As expected the voltage of this experimental cell was below the lithium foil

electrode cells. The LixC electrode was only 72 my below the Lithium electrode potential at the

start of the discharge and increased to 264 my at the end of the discharge. Thus, it had a discharge

curve comparable to that of the lithium foil cells.

SUMMARY

Ambient temperature rechargeable Li-TiS 2 cells have demonstrated a capability of 125 VaAg at

rates greater than C/2 and almost 1000 cycles at 50% depth of discharge. Design studies

identified that a tight pack provided a more consistent and higher operating voltage. Microporous

PE or PP were found to be suitable as separator material. The electrolyte and anode:cathode ratios

among other parameters were also specified. Significant electrical abuse did not result in a venting

or explosion. Overcharge protection was shown to a possible in these cells. A L1 6C alloy was

found to be a suitable replacement for lithium foil with predictable reduction in specific energy for

the loss in lithium metal.

87
ACKNOWLEDGMENT

The work described herein was carried out at the Jet Propulsion Laboratory, California Institute of

Technology, under a contract with the National Aeronautics and Space Administration (Code Q.

88

(—
2 

2 

2 

2 

v

0 2.1

L

T 2

6

1.9

1.8

1.7

1. 6

0 100 200 300 400 500 600 700 800 900 1000

ELAPSED NINUTFS

Figure 1: Li-TiS2 Charge/Discharge Curve

135

132

c^

O

129

cn

0 126

s

123

120

5 HR 10 1.7 VOLTS

f / 4 HR

i

i

/ ) HR

i 2.5 HR

;2 HR

25 30 35 40 45 50 55 60 65

WATTS PER KILOGRAM

Figure 2: Performance of a Typical 1 Ah Li-TiS2 Cell

89
So ,00 ,SO zoo zS0 300' 350

CICL(S

A 0.6

N

P

N

R

S 0.6

04

0.z

0

2.5

2

J

O

i . u

^ 1.5

Q

UC/)

r 1 1O

Z

- 0.5

0^

0 100 200 300 400 500 600 700 800 900 1000

50% DOD CYCLES (C/5 DISCHARGE AND C/10 CHARGE TO 2.6 VOLTS)

Figure 3: Cycle Life of a Li-TiS2 Cell at 50% Depth of Discharge

.rr^i^ Ill rmhlhilU` I1,IIlgiull^lulnwLLlmulpn^mulI!,-1fI !I^dI'LFm,du1-.... -!

Figure 4: Cycle Life of a LI-TiS2 Cell at 100% Depth of Discharge

90
2.2

2.1

2

W

O 1.9

1.8

1.7

1.6

EFFECT OF PACK TIGHTNESS ON END-OF-DISCHARGE VOLTAGE

TIGHT

--- — --- -----

MEDIUM

LOOSE

71

0 50 100 150 200 250 300 350 40Lj 450

CYCLES

Figure 5: Effect of Pack Tightness on End-of-Discharge Voltage

EFFECT OF SEPARATOR ON END-OF-DISCHARGE VOLTAGE

.MICROPOROUS PE

----------- I ROPOROUS PP

SUB-MICRO POROUS^^^

PIP THERMAL

2.5

2

1.5

1

0.5

0

0 50 100 150 200 250 300 350 400 450 rM

CYCLES

Figure 6: Effect of Separator On End-of-Discharge Voltage

91
> 3.30

TM PD

11

Li

Q 2.80 J

0

J 2.30

LiV

240

0

0:00 0:05 0:10 0:15 0:20 0:25 0:30

Time [HRS:MIN

Figure 7: Short Circuit of a lAh 'AA' Li-TiS2 Cell

25

20

N

0 15

E

d

C

d

^j 10

5

220

200

r

c

rJ

180 s

LL

N

160 v

a^

CJ

140

m

m

120 -c,

C.

E

a^

100 ~

80

60

70-75o OF THEORETICAL CAPACITY

(STABILIZED IN 5 CYCLES)

OVERCHARGE PERFORMANCE:

3.80

0.0 10.0 20.0 30.0 10.0

CAPACITY . III

25 CYCLES WITH STABILIZED CAPACITY

* COMPATIBLE WITH THE CELL

* 10-15o OVERCHARGE PROTECTION AT

THE NORMAL CHARGE RATE

Figure 8: Results of adding TMPD to a cell for Overcharge Protection

92
2400

CAPACITY OF ANODE TO CATHODE = 2:1

2200 1 = 0.84 mA/cm2

2000

J

vi

a^

1800

0

1600

E (alloy) vs. Li - 264 (mV) \

E (TiS2) vs. Li - 1.7 (V)

1400

0.0 1.5 3.0 4.5 6.0 7.5

Time (hrs)

Figure 9: Discharge Curve of a LiXC-TiS2 Cell

93
Table 1: Cell Design Parameters

DESIGN PARAMETERS VARIATIONS

ANODE TO CATHODE CAPACITY

RATIO

4 : 1

PACK TIGHTNESS Tight

TYPE OF ELECTROLYTE

Salt: 1.5 M LiAsF 6

10% EC + 88% 2-MeTHF + 2% 2-MeF

THE + 2-MeTHF + 2% 2-MeF,

DIOX + 2-MeTHF + THE + 2% 2-MeF

QUANTITY OF ELECTROLYTE 6 C. C.

SEPARATOR Microporous Polypropylene

Microporous Polyethylene

BINDER CONCENTRATION 1 % by weight

CASE POLARITY floating or positive.

Table 2: Results of Electrical Abuse Tests on Li-TiS2 cells

CYCLE # TYPE OF ABUSE RATE ABUSE % EVENT

200 OVERCHARGE C/10 400 NONE

400 OVERCHARGE C/10 500 NONE

550 OVERCHARGE C/10 500 NONE

400 OVERDISCHARGE C/5 1000 NONE

500 OVERDISCHARGE C/5 1000 NONE

94
Table 3: Rate Capability of a LixC Anode

CURRENT (mA/cm2) CAPACITY (mAh/gm)

0.167 235

0.333 214

0.667 208

1.000 200

95
y J^oZo

IMPEDANCE MEASUREMENTS ON A SPIRAL-WOUND NICKEL/METAL HYDRIDE

CELL CYCLED IN A SIMULATED LEO ORBIT

Margaret A. Reid

National Aeronautics and Space Administration

Lewis Research Center

Cleveland, OH 44135

ABSTRACT

A spiral-wound size C cell was cycled at 25°C in a Low Earth Orbit regime (LEO) at 50% depth of

discharge (DOD) with approximately 5% overcharge. The nominal capacity was 3.5 AH. The cell

was cycled for 2000 cycles. Capacity checks and impedance measurements over the complete range

of state of charge were made upon receipt and after 500, 1000 and 2000 cycles. The capacity of the

cell was essentially unchanged until after the impedance measurements at 2000 cycles. Only small

changes in the impedance parameters were observed, but there was somewhat more scatter in the data

after 2000 cycles. When the cell was returned to LEO cycling after 2000 cycles, only 38% of the

capacity could be obtained. It is believed that the cell failed because of an equipment failure at the

end of the final impedance measurements which allowed an overdischarge.

INTRODUCTION

Nickel-metal hydride cells have been under intensive development in the last few years, and several

manufacturers are selling these for both commercial markets [1, 21 and for space applications [3].

We are initiating testing of these for low-earth orbit applications. Since spiral-wound cells are readily

available, we began with cycling and impedance measurements on an Ovonics size C cell with a

nominal capacity of 3.5 AH.

EXPERIMENTAL

Short leads were soldered onto the two ends of the cell. The capacity was measured by charging at a

C110 rate for 16 hours and discharging at a C/4 rate to a cut-off voltage of 1.0 V. The initial

capacity was 3.38 AH. A value of 3.4 AH was used in the computer program that controlled the

cycling. Open circuit voltage decay measurements when the cell was new gave a loss of about 18%

of capacity in 24 hours and 26% in 48 hours.

After the initial capacity check, the cell was charged and the impedances measured at various depths

of discharge. Measurements were made on the fully charged cell and then at increments of 10%

DOD until the cell was fully discharged. (Additional measurements were taken at 1%, 3% and 5%

DOD after 500 cycles). A program developed by Dr. Stephen Lipka of Florida Atlantic University,

CYCLER, was used to automate the charging, discharging and impedance measurements. The

program was written in Pascal and run on a DOS 286 computer with a PAR 273 potentiostat and a

Solartron 250 Frequency Response Analyzer. A Hewlett-Packard 3455A digital voltmeter was used

to monitor the cell voltage. The computer was programmed to cut off the charging and discharge

currents if a limiting voltage was reached in order to prevent overdischarge and overcharge. The

change of voltage with time on open circuit after discharge was used to determine when the cell had

reached a quasi-equilibrium state.

97

PR+Ii GK*Nlw PAGE BLANK NOT FELLED
The desired depth of discharge was obtained by starting each measurement in the fully charged state

and discharging to the calculated depth of discharge at a 20% LEO rate (0.343 Q. Impedance measurements were taken at the quasi-equilibrium voltage described below. A true equilibrium voltage is

never reached. After the discharge current is interrupted, the voltage rises abruptly, then more

gradually, eventually falling due to self-discharge. We assume that the maximum voltage reached

corresponds to a quasi-equilibrium state. The time that it takes to reach this maximum varies with the

depth of discharge. At low depths of discharge the time is brief, and at high depths of discharge it is

very long, on the order of days before self-discharge is appreciable. In order for the computer to

determine when quasi-equilibrium has been reached, a parameter, S, is defined as the rate of change

of the voltage vs time. Ideally, quasi-equilibrium would be reached when 5 = 0. However it is

impractical to wait to take the impedance measurements until the voltage stops rising completely,

especially at high depths of discharge. Because of the different times required to reach quasiequilibrium at different states of charge, the value of 6 is set for each DOD based on prior observation of the voltage-time curve. The value selected may also depend on the type of cell under study

(and how long one is willing to wait before taking the measurement). The rate of change of voltage

with time is calculated by the computer program, and when this falls below 6, the impedance

measurements are carried out at this quasi-equilibrium voltage. For the fully charged cell, b was set

at 0.005 mV/min, and this was decreased as the depth of discharge was increased. For 10% DOD, S

was set at 0.001 mV/min, and at 50% DOD and below was 0.00022 mV/min. Typical equilibration

times using these values were about 20 minutes for a fully charged cell, 1-2 hours at 30% DOD, and

4-6 hours below 50% DOD.

After each measurement the cell was recharged by adding the capacity withdrawn on the prior discharge plus an additional 10% of the capacity of the cell.

The capacity on receipt was 3.38 AH. After the initial impedance measurements, the capacity had

increased to 3.47 AH. The cell was then cycled for 500 LEO cycles at 50% DOD using a 55 minute

charge time and a 35 minute discharge time. Approximately 5% overcharge was used. This

corresponds to a 0.573 C charge rate and a 0.857 C discharge rate. The cutoff voltage on discharge

was 1.00 V. The capacity and impedance measurements were then repeated. Measurements were

repeated again after 1000 cycles and 2000 cycles.

After the cell failed, it was opened in order to determine the area and thickness of the individual

electrodes for comparison of the impedances with those of boilerplate Ni/H 2 cells on a unit area basis.

The nickel electrode was about 40 mils thick with an area of about 94 cm'. (The initial thickness was

probably smaller). The hydride electrode was about 15 mils thick with an area of about 113 cm'.

Attempts to make further measurements with the hydride electrode were unsuccessful due to its

disintegration when current was applied.

EXPERIMENTAL RESULTS

The capacity remained essentially constant, increasing slightly to 3.49 AH before cell failure at the

end of the impedance measurements at 2000 cycles.

Initial analysis of the data was done using the conventional circuit shown in Fig. 1 with the ohmic

resistance in series with a parallel combination of a capacitor representing the double-layer capacitance or adsorption capacitance and a resistor representing the kinetic resistance of the electrode. The

ohmic resistance of the cell was about 16.5 ± 0.5 mOhms and did not vary with the DOD or cycle

number. Little change in the other impedance parameters was observed until very low depths of

98
discharge, similar to observations on boilerplate Ni/HZ cells. Table I lists the parameters for the cells

at 2000 cycles.

After 2000 cycles the impedances were essentially the same as at 1000 cycles, with slightly larger

values at the lowest frequencies. At the higher states of charge there was an indication of the start of

a second RC circuit. At the lower states of charge, there was slightly more scatter. Figures 2-5

show the comparisons for 5%, 20%, 70% and 90% DOD.

DISCUSSION

Very little data is available in the literature as yet on the impedances of hydride electrodes and nickelmetal hydride cells. Kuriyama et al measured impedances of electrodes made from an alloy powder

(MmNiMCOMA1, 8) pressed with a PFE/PFP binder, both with and without a copper coating on the

alloy particles. Agerwal and Orazem measured impedances of a LaNis solid electrode. The

electrodes in the Ovonics cell studied here are made from an alloy containing Ni, Ti, V, Zr and other

trace elements.

In order to obtain an estimate of the impedance of the hydride electrode in this cell, the data were put

on a unit area basis and the complex plane plots compared to those of the boilerplate Ni/H Z cells that

we are also testing. This is shown in Fig. 6 for the two cells at a 20% DOD. The boilerplate Ni

electrodes have capacities of about 0.033 AH/cm 2 compared to about 0.037 AH/cm2 of the Ni electrodes in the metal hydride cell. If we assume that the impedance of the hydrogen electrode in the

boilerplate Ni/H2 cells is negligible [8], and that the impedances of the two nickel electrodes are about

the same at corresponding states of charge, then the increased impedance of the hydride cell over that

of the boilerplate Ni/H2 cell can be attributed to the hydride electrode. An estimate of the impedance

of the hydride electrode calculated from the differences at each frequency is shown in Figure 7.

Although these are very approximate calculations, they suggest that the impedance of the metal

hydride electrode is not negligible, in contrast to that of the gaseous hydrogen electrode in a typical

Ni/H2 cell.

A confirmation that these estimated values for the impedance of the metal hydride electrode are

reasonable can be obtained by comparison of this data with that of Agerwal and Orazem (Fig. 8).

The impedance data are of the same form at that found by these authors, (i.e., an ohmic resistance in

series with a single parallel RC circuit which includes a diffusion impedance) although they worked

with a different alloy electrode (LaNi .,) and used a solid rod rather than a powdered electrode.

Estimating the true surface area of the powdered alloy from published literature data using a surface

area of about 2000 cm'-/g and a capacity of about 400 mAH/g [8], our impedances calculated with

respect to the true surface area are of the same order of magnitude as those of Agerwal and Orazem,

where the surface area was 1.5 cm' .

The data of Kuriyama et al were measured on a bonded hydride electrode with a binder of polyperfluoroethylene/polyperfluoropropylene (10% by weight) and have an additional semicircle in the

complex plane plots, believed to be due to the contact resistance between the alloy and the current

collector. The remaining portion of the curve is similar to our curves, but insufficient data is given

to make a more detailed comparison.

Spectroscopic analysis of the Ni electrode after dismantling showed a Zr content of about 1 %. This

is consistent with the findings of earlier researchers that the Zr dissolves extensively (and preferentially) from this alloy in 31 % KOH at 70°[3]. Since nickel electrodes commonly last for thousands of

99
cycles before failure, the failure is presumably caused by changes in the hydride electrode, either as a

result of the cycling or from the failure of the equipment. It would be of interest to see if the metal

ions that have precipitated in or on the Ni electrode have an effect on the cycle life, but this cannot be

studied without measurements on the individual electrodes in situ.

More information about the changes brought about by cycling and the mechanism of failure might be

obtained if both the nickel electrode and the hydride electrode are isolated from the case so that the

case can be used as a quasi-reference electrode [8]. Studies will be initiated shortly on cells of this

type which will allow the contributions of the individual electrodes to be determined.

Since the cell impedance did not change appreciably during the 2000 cycles that the cells were cycled,

it cannot be determined from this limited study whether impedance can be used as a diagnostic criteria

for metal hydride cells, especially in this situation where the negative electrode is connected to the

case.

CONCLUSION

Impedances of the spiral-wound metal hydride cell changed only slightly over the course of 2000 LEO

cycles. Both the nickel electrode and the metal hydride electrode apparently contribute to the

impedance of the cell, with the larger portion probably due to the Ni electrode. In the absence of a

reference electrode, the contributions of each electrode to the impedance cannot be determined.

Future measurements will be made with cells constructed with the hydride electrode not connected to

the case. This arrangement allows the case to be used as a quasi-reference electrode, and thus the

effects of cycling can be determined for the individual electrodes.

REFERENCES

M. Fetcenko, S. Venkatesan, S. Dhar, S. Ovshinsky and R. Tsang, Proceedings of the 3rd

International Rechargeable Battery Seminar, Deerfield Beach, Florida, March 1992.

2. S. Ovshinsky, P. Gifford, S. Venkatesan, D. Corrigan and S. Dhar, Proceedings of the Tenth

International Seminar on Primary and Secondary Battery Technology and Application,

Deerfield Beach, FL, March 1993.

3. Sealed Metal-Hydride Batteries for Aerospace Applications, D. Coates, R.D. Wright, pp.

199-209, Proceedings of the Symposium on Hydrogen Storage Materials, Batteries, and

Electrochemistry, Eds. D.A. Corrigan and S. Srinivasan, The Electrochemical Society, Pennington, N.J., Vol. 92-5.

4. N. Kuriyama, T. Sakai, H. Miyamura, I. Uehara, H. Ishikawa and T. Iwasaki, J. Electrochem. Soc., 139 (1992) L72-L73.

5. M.G.S.R.Thomas, P.G. Bruce and J.B. Goodenough, J. Electrochem. Soc., 132 (1985) 1521-

1528.

6. B.K. Zoitos, D.L. Hudson, P.D. Bennett and V.J. Puglisi, pp. 168-178, Proceedings of the

Symposium on Hydrogen Storage Materials, Batteries, and Electrochemistry, Eds. D.A.

Corrigan and S. Srinivasan, The Electrochemical Society, Pennington, N.J., Vol. 92-5.

100
P. Agerwal, M. Orazem and A. Hiser, pp. 120-137, Proceedings of the Symposium on

Hydrogen Storage Materials, Batteries, and Electrochemistry, Eds. D.A. Corrigan and S.

Srinivasan, The Electrochemical Society, Pennington, N.J., Vol. 92-5.

M.A. Reid, J. Power Sources, 29, 467-76 (1990).

J.J. Reilly, pp 24-45, Proceedings of the Symposium on Hydrogen Storage Materials,

Batteries, and Electrochemistry, Eds. D.A. Corrigan and S. Srinivasan, The Electrochemical

Society, Pennington, N.J., Vol. 92-5.

101
Table I.—Selected impedance parameters at 2000 cycles using the standard circuit

model.

DOD Roll,„ mOhm Rk ^„ mOhm Capacitance,

Farads

Slope,

mOhm sec-112

5 15.1 14.3 —2000 2.66

20 15.3 14.3 —4000 2.95

70 5 20.9 -- 6000 2.31

90 [±1'5 3 42.1 --8000 1.44

102
double layer

—_^ i m I

1

R

kin

Z

Figure I.—Simplified circuit model for

analysis of impedance parameters.

—.0

Z"

1000 CYCLES

2000 CYCLES

0 06

Z I

Figure 2.—Complex plane plots at 5 percent DOD for

1000 and 2000 cycles. V = 1.345.

103

— i
0 .05

Z'

Z"

2000 CYCLES

1000 CYCLES

. 05

Z'

Figure 3.—Complex plane plots at 20 percent DOD for

1000 and 2000 cycles. V — 1.340.

Figure 4.—Complex plane plots at 70 percent DOD for

1000 and 2000 cycles. V = 1.262.

-.0

Z ' ,

0

104
0 .05

Z'

Z"

1.6

1.4

1.2

1

N 0.8

0.6

0.4

0.2

Figure 5.—Complex plane plots at 90 percent DOD for

1000 and 2000 cycles. V = 1.206.

-

0 0.2 0.4, 0.6 0.8 1 1.2 1.4 1.6 1.8 2

Z'

METAL HYDRIDE + BOILERPLATE CELL, 26% KOH

0 Figure 6.—Comparison of impedances of boilerplate Ni/1-1 2 cell and Ni/MH cell on a unit

projected area basis. 20 percent DOD, 1000 LEO cycles.

105
1.6

1.4

1.21

1

N 0.8

0.6 -

0.4

0.2

0

0 0.2 0.4 0.6 0.8 1 1.2 1.4 1.6 1.8 2

Z'

® HYDRIDE BY DIFFERENCE

Figure 7.—Estimated impedance of the metal hydride electrode at 20 percent DOD by difference

between Ni/MX cell and boilderplate Ni/H 2 cell. After 1000 LEO cycles.

i

106
2500

2000

1500

`N

C

L

1000

500

— 500 i

10 -1 100 101 102 103 104 105

Frequency, Hz

800

600

VI E 400 L

O

N

200

0

—200

10 -1

i

100 101 102 103 104 105

Frequency, HZ

(a) Agerwal and Orazem.

Figure 8.—Comparison of estimated impedance of metal

hydride electrode at 20 percent DOD (estimated true

surface area of about 20 000 cm 2 ) with literature

data of Agerwal and Orazem for a partially charged

electrode ( 1.5 cm rod).

107
0.6

0.5

0.4

N 0.3

0

z

N 0.2

0.1

-0.

-1 0 1 2 3

LOG FREQUENCY (HZ)

© Z' + -Z11

(b) This study.

Figure 8.—Concluded.

108
N-23351

NASA LEWIS ADVANCED IPV NICKEL-HYDROGEN TECHNOLOGY

John J. Smithrick

and

Doris L. Britton

National Aeronautics and Space Administration

Lewis Research Center

Cleveland, Ohio 44135

ABSTRACT

Individual pressure vessel (IPV) nickel-hydrogen technology was advanced at NASA Lewis and

under Lewis contracts. Some of the advancements are as follows: (1) to use 26 percent potassium

hydroxide electrolyte to improve cycle life and performance, (2) to modify the state-of-the-art cell design

to eliminate identified failure modes and further improve cycle life, and (3) to develop a lightweight

nickel electrode to reduce battery mass, hence reduce launch and/or increase satellite payload.

A breakthrough in the LEO cycle life of individual pressure vessel nickel-hydrogen battery cells

was reported. The cycle life of boiler plate cells containing 26 percent KOH electrolyte was about 40,000

accelerated LEO cycles at 80 percent DOD compared to 3,500 cycles for cells containing 31 percent

KOH. Results of the boiler plate cell tests have been validated at NWSC, Crane, Indiana. Forty-eight

ampere-hour flight cells containing 26 and 31 percent KOH have undergone real time LEO cycle life

testing at an 80 percent DOD, 10°C. The three cells containing 26 percent KOH failed on the average at

cycle 19,500. The three cells containing 31 percent KOH failed on the average at cycle 6,400.

Validation testing of NASA Lewis 125 Ah advanced design IPV nickel-hydrogen flight cells is

also being conducted at NWSC, Crane, Indiana under a NASA Lewis contract. This consists of

characterization, storage, and cycle life testing. There was no capacity degradation after 52 days of

storage with the cells in the discharged state, on open circuit, 0°C, and a hydrogen pressure of 14.5 psia.

The catalyzed wall wick cells have been cycled for over 22,694 cycles with no cell failures in the

continuing test. All three of the noncatalyzed wall wick cells failed (cycles 9,588, 13,900 and 20,575).

Cycle life test results of the Fibrex nickel electrode has demonstrated the feasibility of an

improved nickel electrode giving a higher specific energy nickel-hydrogen cell. A nickel-hydrogen boiler

plate cell using an 80 mil thick, 90 percent porous Fibrex nickel electrode has been cycled for 10,000

cycles at 40 percent DOD.

INTRODUCTION

The state of development of individual pressure vessel nickel-hydrogen battery cells is such that they

are acceptable for geosynchronous orbit (GEO) applications since not many cycles are required over the

life of the battery system (1000 cycles, 10 years). There are 20 communication satellites in GEO using

IPV nickel-hydrogen batteries [1]. For the demanding Low-Earth-Orbit (LEO) applications, however, the

current cycle life at moderate-to-deep depths-of-discharge (40 to 80 percent) should be improved.

Battery cycle life has a major impact on life cycle cost for LEO applications such as Space Station

Freedom (30 year life). The primary drivers are transportation to orbit and battery cost.

IPV nickel-hydrogen technology was advanced at NASA Lewis and under Lewis contracts. Some

of the advancements are as follows: (1) to use 26 percent potassium hydroxide electrolyte to improve

cycle life and performance, (2) to modify the state-of-the-art cell designs to eliminate identified failure

modes and further improve cycle life and (3) to develop a lightweight nickel electrode to reduce battery

mass, hence reduce launch cost and/or increase satellite payload.

The influence of potassium hydroxide electrolyte concentration on cycle life was investigated at

Hughes Aircraft Company under a NASA Lewis Contract. Hughes reported a breakthrough in LEO cycle

109
life [2, 31. Boiler plate cells containing 26 percent KOH were cycled for about 40,000 accelerated cycles

at 80 percent depth-of-discharge (DOD) at 23'C, compared to 3,500 cycles for cells containing 31 percent

KOH. These results are in the process of being validated using flight hardware and real time LEO test

under a NASA Lewis Contract with the Naval Weapons Support Center, Crane, Indiana.

The advanced design for an IPV nickel-hydrogen cell was conceived with the intention of

improving cycle life at moderate-to-deep depths-of-discharge (40 to 80 percent). The approach was to

review IPV nickel-hydrogen cell designs and results of cycle life tests conducted in-house and by others

to identify areas where improvements could result in a longer life [4 - 10]. The feasibility of the design

was demonstrated using 6 Ah boiler plate cells [11]. The advanced design is in the process of being

validated using 125 Ah flight cells and real time test under a NASA Lewis Contract with the NWSC,

Crane, Indiana.

A lightweight nickel electrode using a thick, porous, fiber substrate has been evaluated in-house

using boiler plate cells. The thickness is 80 mils, porosity 90 percent, active material loading level

1.6 g /cm3 void volume.

In this report results of validation tests will be updated [12,13].

MEASUREMENTS AND PROCEDURES

For both the 48 and 125 Ah cells, the quantities measured every 2.4 min for each cell during charge

and discharge and their accuracies are: current (± 2.0 percent), voltage (± 0.001 percent), pressure (± 1

percent), and temperature (± 1 percent). Charge and discharge ampere-hour capacities are calculated

from current and time. Charge-to-discharge ratio (ampere-hours into cell on charge to ampere-hours out

on discharge) is calculated from the capacities. Cell charge and discharge currents are calculated from

the voltage measured across a shunt, using an integrated digital voltmeter. Cell pressure is measured

using a strain gauge located on the cell dome. The temperature is measured using a thermistor located on

the center of the pressure vessel dome. The therrnistor is mounted using a heat sink compound to insure

good thermal contact.

To verify the performance of lightweight nickel electrodes in boiler plate hardware, a nickelhydrogen cell was assembled at NASA Lewis. The quantities measured during charge and discharge are:

current (± 2.0 percent), voltage (± 0.001 percent), and time (± 1 percent). The charge and discharge

ampere-hours and watt-hours are calculated from current and time on every scan. Every 18 sec, the data

is saved and stored to a data collector. The data can be retrieved for data and graphical output. The

percent utilization is calculated by using the ratio of the measured discharge capacity to the theoretical

calculated capacity based on the weight of the active material deposited. End-of-life failure is defined as

the point where the discharge voltage degrades to 0.5 V.

For the 48 Ah cells, prior to cell final hydrogen gas adjustment, the nickel electrodes were positively

charged, which results in a 0 psia hydrogen gas pressure. After completion of acceptance testing the cells

were discharged at the C/10 rate (4.8 A) to 0.1 V or less. The cells were shipped to NWSC, Crane,

Indiana, where they were stored at 10°C under trickle charge at C/200 for 31 days. After storage the

discharge ampere-hour capacity acceptance test was repeated. The capacity was measured after charging

the cells at the C/2 rate (24 A) for 2.0 hr, then C110 for 6 hr followed by a 0.5 hr open circuit stand. The

discharge capactiy was measured to 1.0 V at each of the following rates: C2, C, 1.4, and 2 C.

Prior to undergoing cycle life testing, the capacity retention after a 72 hr open circuit stand (10°C)

was measured for each cell. For the cycle life test the cells were connected electrically in series to form a

six cell pack. The cycle regime was a 90 min LEO orbit consisting of a 54 min charge at a constant

0.93 C rate (44.7 A) followed by a 36 min discharge at a 1.33 C rate (64 A). The charge to discharge

ratio was 1.048. The depth-of-discharge was 80 percent of name plate capacity (48 Ah). During the

cycle life test the cooling plate temperature was maintained at 10±2°C. Cell failure for this test was

defined to occur when the discharge voltage degrades to 1.0 V during the course of the 36 min discharge.

110
For the 125 Ali after completion of activation by the manufacturer, the precharge hydrogen

pressure was set to 0 prig (14.5 psia) with the nickel electrodes in the fully discharged state. After

completion of the acceptance testing the cells were discharged at the C110 rate (12.5 A) to 0.1 V or less

and the terminals were shorted. The cells were shipped to NWSC, Crane, where they were stored at open

circuit, and discharged at 0'C for 52 days. After storage the discharge ampere-hour capacity acceptance

test was repeated. The capacity was measured after charging the cells at the C/2 rate (62.5 A) for 2 hr

then C110 for 6 hr followed by a 0.5 hr open circuit stand. The discharged capacity was measured to

1.0 V for each of the following rates: C/2, C, 1.4 C, and 2 C.

Prior to undergoing cycle life testing the capacity retention after a 72 hr open circuit stand (10'C)

was measured for each cell. For the cycle life test the cells were connected electrically in series to form a

six cell pack. The cycle regime is a 90 min LEO orbit consisting of a 54 min charge at a constant 0.69 C

rate (87 A) followed by a 36 min discharge at a C rate (125 A). The charge-to-discharge ratio was 1.04.

The depth-of-discharge was 60 percent of name plate capacity (125 Ali). During the cycle life test the

cooling plate temperature was maintained at l Ot2°C. Cell failure for this test was defined to occur when

the discharge voltage degrades to 1.0 V during the course of the 36 min discharge.

The boiler plate cell was cycle life tested using a 90 min cycle regime at 40 percent DOD. The

continuous test regime consists of charging at a constant 0.48 C rate for 55 min immediately followed by

discharge at a constant 0.69 C rate for 35 min. Cell capacities are measured every 1000 cycles, by

charging at the C rate for 80 min followed by discharging at the 0.69 C rate to 0.5 V. Failure of the cell

occurs when the discharge voltage degrades to 0.5 V during the course of the constant current 35 min

discharge at the 0.69 C rate.

EXPERIMENTAL

TEST FACILITY - The facility is capable of testing 45 battery packs with maximum of 10 cells

electrically connected in series per pack. Each pack has its own charge and discharge power supply

controlled by a computer which is programmed to satisfy the particular test requirements. During testing,

each pack is scanned every 2.4 min to compare data such as voltage, temperature, and pressure with

programmed limits. If a parameter is out of limit, an alarm will be initiated and a message will be typed

out identifying the cell and parameter. The data is recorded on a 132 mB disc drive and if requested can

be obtained in report form. The cell temperature during a test is controlled by a recirculating cooler that

circulates a solution of water and ethylene glycol through a cooling plate.

CELL DESCRIPTION - 48 Ali Flight Cells - Six Air Force/Hughes recirculating design IPV nickelhydrogen flight cells manufactured by Hughes are undergoing testing. Three of the cells contain 26

percent KOH electrolyte (test cells). The other three (control cells) are identical to the test cells but they

contain 31 percent KOH. Both the test and control cells contain an equal number of components. The

name plate capacity is 48 A-hr. The cell is illustrated in Fig. 1. It consists of a stack of nickel electrodes,

separators, hydrogen electrodes, and gas screen assembled in a non-back-to-back electrode configuration.

In this configuration electrodes of different types directly face each other. The stack is packaged in a

cylindrical pressure vessel, with hemispherical end caps. This is made of Inconel 718 and lined with

zirconium oxide which serves as a wall wick. The components are shaped in a pineapple slice pattern.

Like electrodes are connected electrically in parallel. The separators consist of two layers of zircar,

which extend beyond the electrodes to contact the wall wick. Hence, the electrolyte which leaves the

stack during cycling will be wicked back. The gas screens are polypropylene. The nickel electrode

consists of a dry sinter plaque containing a nickel screen substrate which was electrochemically

impregnated by the alcoholic Pickett process [14].

CELL DESCRIPTION - 125 Ali Flight Cells - Six 125 Ali design IPV nickelhydrogen flight cells fabricated by Eagle-Picher, Joplin according to NASA Lewis specification are

undergoing cycle life testing. The nickel electrodes were fabricated at Eagle-Picher, Colorado Springs

and were impregnated with active material by the alcoholic Pickett process [ 14]. Three of the cells (test

cells) contain all of the advanced design features as described in [15]. The other three cells (control cells)

are the same as the test cells except they do not have catalyst on the wall wick. All six cells contain 26

rather than 31 percent KOH electrolyte.

111
The test cell design is illustrated in Fig. 2. The new features of this design which are not

incorporated in the state-of-the-art Air Force/Hughes or COMSAT/Intelsat Cells are: (1) use of 26 rather

than 31 percent KOH electrolyte which improves cycle life [2, 3, and 16], (2) use of a catalyzed wall

wick located on the inside surface of the pressure vessel wall which chemically recombines oxygen

generated at the end of charge and on overcharge with hydrogen to form water. State-of-the-art nickelhydrogen cells recombine the oxygen on the catalyzed hydrogen electrode surface in the stack. The

catalyzed wall wick should improve oxygen and thermal management [17], (3) use of serrated edge

separators to facilitate gaseous oxygen and hydrogen flow within the cell, while still maintaining physical

contact with the wall wick for electrolyte management, (4) use of a floating rather than fixed stack (SOA)

to accommodate nickel electrode expansion due to charge/discharge cycling. This is accomplished by use

of Belleville disc springs located at each end of the stack. The significant improvements resulting from

these innovations are extended cycle life, enhanced oxygen, thermal and electrolyte management, and

accommodation of some of the nickel electrode expansion.

LIGHTWEIGHT NICKEL ELECTRODE -The lightweight nickel electrodes evaluated were

fabricated from Fibrex nickel fiber substrates. The substrates contain 50 percent nickel fibers, 35 percent

nickel powder and 15 percent cobalt powder. The porosity was 90 percent and the thickness was 80 mils.

They were electrochemically impregnated by Eagle-Picher using the aqueous process to a loading level of

1.6 g/cm3.

CELL DESCRIPTION - For Testing Lightweight Nickel Electrodes - The cycle test was conducted

using a boiler plate cell with a 5- by 5-cm active area. The components of the cell consist of a stack of

nickel electrode, separator, hydrogen electrode, and gas screen. The nickel electrode is made from an

80 mil thick, 90 percent porous Fibrex nickel electrode loaded to 1.6 g/cm 3 void volume. The separator

consists of three layers of 6-mil thick beater treated asbestos separator. The hydrogen electrode was

made by Life Systems, Incorporated and consists of a mixture of platinum and Teflon applied to a gold

plated nickel screen. The gas screen, located behind the hydrogen electrode, is 60 mil thick nickel Exmet

which was compressed to 40 mil to fit into the cavity designed for the hydrogen electrode and gas screen.

An electrolyte reservoir plate (ERP), which is incorporated into the cell, is made of 125 mil-thick foam

metal. The current collectors are gold plated nickel. A graphic representation of the unit cell cross

section is shown in Fig. 3. After assembly, the entire stack is vacuum-filled with 26 percent KOH

electrolyte. The electrolyte is allowed to stand in the stack overnight, then the excess electrolyte is

drained. The stack is placed in a pressure vessel which is evacuated and filled with hydrogen to 3.4 atm

(50 psi).

RESULTS AND DISCUSSION

48 Ah FLIGHT CELLS - Storage Test - The nickel-hydrogen battery could undergo a planned or

unplanned storage due to delays prior to launch. This may have an effect on performance. The influence

of storage (31 days, trickle charged at C/200, 10°C) on the capacity of the 48 Ah IPV nickel-hydrogen

flight cells containing 26 and 31 percent KOH electrolyte is shown in Fig. 4. The spread in the data

indicate there is no significant capacity loss after 31 days for either the 26 or 31 percent KOH cells.

48 Ah FLIGHT CELLS - Performance Test - A comparison of the average discharge voltage (three

cells) as a function of time for the cells containing 26 and 31 percent KOH was made and is shown in Fig.

5. The voltage for the 26 percent KOH cells is higher than for the the 31 percent KOH cells up to about

an 82 percent DOD. The discharge rate was 1.4 C (67.2 A) and the cell temperature was maintained at

10'C. The ampere-hour capacity for these cells is shown in Table 1 (1.4 C, 10'C). The capacity on the

average for the 26 percent KOH cells was about 10 percent lower than the 31 percent KOH cells. The

tradeoff for this relatively small decrease in initial capacity is a significant increase in cycle life. It should

be noted that the data in Table I is for a 100 percent DOD. In an actual application the DOD will be

much less. For instance the DOD for Space Station Freedom will be about 35 percent. In this operating

region the cells containing 26 percent KOH have a higher discharge voltage, and still have adequate

capacity reserve.

48 Ah FLIGHT CELLS - Cycle Test - The influence of LEO cycling at 80 percent DOD on the end

112
of discharge voltage for the 48 Ah IPV nickel-hydrogen flight cells containing 26 percent KOH is

summarized in Fig. 6. The three cells containing 26 percent KOH failed on the average at cycle 19,500

(cycle 15,314, 19,518, and 23,659). The influence of cycling on the end of charge pressure for the 26

percent KOH cells is shown in Fig. 7. The pressure increase per 1000 cycles if 23.3 PSI. The pressure

increase could be indicative of nickel plaque corrosion which converts nickel to active material. The

increase in pressure will result in a shift in the beginning of life state-of-charge versus pressure curve.

The influence of LEO cycling at 80 percent DOD on the end of discharge voltage for the cells

containing 31 percent KOH is shown in Fig. 8. The three cells containing 31 percent KOH failed on the

average at cycle 6,400 (cycles 3,729, 4,165 and 11,355). The failure mode for each cell was

characterized by degradation of discharge voltage to 1.0 V. No cell failed due to an electrical short. A

comparison of the discharge curve at the beginning and end of life for Cell 1, which failed at cycle 3,729,

is shown in Fig. 9. This information also shows a voltage degradation. The ampere-hour capacity

decrease for cell 1 was about 33 percent (1.4 C rate, 10°C), for cell 2, 33 percent, and for cell 3, 36 percent. The influence of cycling on the end of charge pressure for the 31 percent KOH cells is shown in

Fig. 10. The pressure change can be correlated with the discharge voltage change due to cycling. The

pressure increase per 1000 cycles is 23.3 PSI. The pressure increase is the same as for the 26 percent

KOH.

The cycle life of the cells containing 26 percent KOH was a factor of 3 to 4 better than those with

31 percent KOH. The superior performance of the 26 percent KOH cells compared to the 31 percent cells

is in agreement with boiler plate cell results reported previously (2 and 3). It is attributed to

crystallographic change of active material [ 16]. Gamma NiOOH is converted to beta MOOH in 26

percent KOH. Beta NiOOH has a lower capacity but longer life.

48 Ah FLIGHT CELLS - Destructive Physical Analysis - Destructive physical analysis (DPA) of

all three of the 31 percent KOH cells was completed and documented at Hughes under a NASA Lewis

contract [18]. DPA of the 26 percent cells is in process. A summary of the DPA results of the

31 percent KOH cells is as follows: All three cells failed during cycling due to a decrease in voltage and

nickel-electrode capacity. The capacity decrease was confirmed by measuring nickel-electrode capacity

in flooded electrolyte cells. Some observations which could cause the capacity decrease are nickelelectrode expansion, rupture and corrosion of the nickel-electrode substrate, active material redistribution,

and accumulation of electrochemically undischarged active material with cycling. Cell 3 appears to have

failed by gradual wear-out due to these changes. Some of the electrodes from cells 1 and 2 showed a

premature capacity fading which was responsible for early failure. However, chemical analysis of these

electrodes did not show anomalous results. The mechanism of the premature capacity fading is not fully

understood by the present DPA. No cells failed due to an electrical short. All cells showed some

increase in internal resistance after the cycle test; however, this increase itself does not appear to be the

direct cause of failure. All cells showed a decrease in discharge voltage and an increase in charge voltage

after the cycle test.

125 Ah ADVANCED FLIGHT CELLS - Cell Performance - For a representative 125 Ah advanced

catalyzed wall wick nickel-hydrogen flight battery cell the voltage and pressure during charge and

discharge are shown in Fig. I 1 (beginning of life). The discharge rate was 0.69 C (87 A) and the

temperature was a nominal 10°C. The mid-discharge voltage was 1.248 V. The pressure, as expected,

varies linearly with the state-of-charge. It should be noted, however, that the pressure could increase with

charge/discharge cycling causing a shift in the state-of-discharge curve.

The effect of discharge rate on ampere-hour capacity for a representative cell of each type is shown

in Fig. 12. The capacity decreased slightly (1 percent) over the range of C/2 to 1.4 C, after which point it

decreased rapidly. In a nickel-hydrogen cell the gaseous hydrogen comes into contact with the nickel

electrodes resulting in a capacity loss due to self discharge. The capacity retention of the cells after a

72 hr open circuit stand at 10°C is shown in Fig. 13. The data shows no significant difference in capacity

retention between the catalyzed and noncatalyzed wall wick cells. The capacity retention for the

catalyzed wall wick cells on the average is 84 percent and for the noricatalyzed wall wick cells is

85 percent.

113
125 Ah ADVANCED FLIGHT CELLS -Storage Test - The effect of storage (52 days, discharged,

open circuit, 0T) on the capacity of the six 125 Ah flight IPV nickel-hydrogen cells is summarized in

Fig. 14. The spread in the data shows no significant capacity loss for either the catalyzed or noncatalyzed

wall wick cells due to the 52 day storage. Actually, there was a slight average increase in capacity for

both the catalyzed and noncatalyzed wall wick sells.

125 Ah ADVANCED FLIGHT CELLS - Cyclee Test - The influence of LEO cycling at 60 percent

DOD on the end of discharge voltage for the 125 Ah catalyzed wall wick IPV nickel-hydrogen flight cells

is summarized in Fig. 15. After 22,694 cycles there has been no cell failure in the continuing test. The

influence of cycling on the end of charge pressure for the catalyzed wall wick cells is shown in Fig. 16.

No pressure for cell 2 is available because the cell had a bad strain gauge. For cells 1 and 3 the pressure

increased relatively rapidly up to about cycle 1400 then decreased. The average pressure increase at

cycle 1400 is about I 1 percent higher than at the beginning of life.

The influence of LEO cycling at 60 percent DOD on the end of discharge voltage for the 125 Ah

noncatalyzed wall wick IPV nickel-hydrogen flight cells is shown in Fig. 17. All three of the

noncatalyzed wall wick cells failed (cycles 9,588, 13,900, and 20,575). The failure was characterized by

degradation of end of discharge voltage to 1.0 V. The cells did not fail due to an electrical short. The

influence of cycling on the end of charge pressure for the noncatalyzed wall wick cells is shown in

Fig. 18. The pressure for the three cells increased up to about cycle 2000 then decreased. The average

pressure increase at cycle 2000 is about 9 percent higher than at the beginning of life.

The cycle life testing will continue until cell failure. A post-cycle teardown and failure analysis will

be conducted to evaluate the cause of failure. This information will be used to affect further

improvements.

BOILER PLATE CELL - For Testing Lightweight tweight Nickel Electrodes - The cell cycled for 10,000

cycles, as shown in Fig. 19, before the test was terminated. The effect of cycling on the end of discharge

voltage is shown in Fig. 20. An end-of-discharge voltage of about 1.175 V is observed for the first 1000

cycles. The average end-of-discharge voltage gradually dropped to about 1.060 V at 9000 cycles and

remained constant until the end of the cycle test.

Final characterization testing will be conducted and will be reported later. Failure analysis will also

be conducted to evaluate the cause for failures. CONCLUDING REMARKS

A breakthrough in the LEO cycle life of individual pressure vessel nickel-hydrogen battery cells

was reported. The cycle life of boiler plate cells containing 26 percent KOH electrolyte was about 40,000

accelerated LEO cycles at 80 percent DOD compared to 3,500 cycles for cells containing 31 percent

KOH. Results of the boiler plate cell tests have been validated at NWSC, Crane, Indiana. Forty-eight

ampere-hour flight cells containing 26 and 31 percent KOH have undergone real time LEO cycle life

testing at an 80 percent DOD, 10'C. The three cells containing 26 percent KOH failed on the average at

cycle 19,500. The three cells containing 31 percent KOH failed on the average at cycle 6,400.

Validation testing of NASA Lewis 125 Ali design IPV nickel-hydrogen flight cells is also

being conducted at NWSC, Crane, Indiana under a NASA Lewis contract. This consists of

characterization, storage, and cycle life testing. There was no capacity degradation after 52 days of

storage with the cells in the discharged state, on open circuit, 0'C, and a hydrogen pressure of 14.5 psia.

The catalyzed wall wick cells have been cycled for over 22,694 cycles with no cell failures in the

continuing test. All three of the noncatalyzed wall wick cells failed (cycles 9,588, 13,900 and 20,575).

Cycle life test results of the Fibrex nickel electrode have demonstrated the feasibility of an

improved nickel electrode giving a higher specific energy nickel-hydrogen cell. A nickel-hydrogen boiler

plate cell using an 80 mil thick, 90 percent porous Fibrcx nickel electrode has been cycled for 10,000

cycles at 40 percent DOD.

114
REFERENCES

1. L. Miller, "The Ni-H2 Battery System: A Space Flight Application Summary," 1988 IECEC;

Proceedings of the Twenty-Third Intersocieryy Energy Conversion Engineering Conference, Vol. 2,

ASME, New York, 1988, pp. 489-492.

2. H. S. Lim, and S. A. Verzwyvelt, "KOH Concentration Effect on Cycle Life of Nickel-Hydrogen

Cells, I1I. Cycle Life Test," Journal of Power Sources, Vol. 22, Mar.-Apr. 1988, pp. 213-220.

3. H. S. Lim, and S. A Verzwyvelt, "KOH Concentration Effect on the Cycle Life of Nickel-Hydrogen

Cells, IV. Results of Failure Analysis," Journal of Power Sources, Vol. 29, Feb. 1990, pp. 503-519.

4. E. Adler, T. Duranti, P. Frisch, T. Jacewicz, H. Rogers, L. Samoss, S. Stadnick, and L. Tinker,

"Nickel-Hydrogen Battery Advanced Development Program," AFWAL-TR-80-2044, Hughes

Aircraft Company, 1980.

5. D. Warnock, "Life Test of 50 Ah NiH2 Battery," The 1981 Goddard Space Flight Center Workshop,

G. Halpert, ed., NASA CP-2217, 1982, pp. 487-500.

6. J. J. Smithrick, "Cycle Life Test and Failure Model of Nickel-Hydrogen Cells," IECEC '83:

Proceedings of the Eighteenth Intersociery Energy Conversion Engineering Conference, Vol. 4,

A1ChE, New York, 1983, pp. 1535-1542.

7. D. H. Fritts, "Testing the Mechanical Characteristics of Sintered Nickel-Battery Plaque and Their

Relationship to Nickel-Electrode Performance," Journal of Power Sources, Vol. 6, Apr. 1981,

pp. 171-184.

8. D. F. Pickett, H. H. Rogers, L. A. Tinker, C. A. Bleser, J. M. Hill, and J. S. Meador, "Establishment

of Parameters for Production of Long Life Nickel Oxide Electrodes for Nickel-Hydrogen Cells,"

Energy to the 21st Century; Proceedings of the Fifteenth Intersociery Energy Conversion

Engineering Conference, Vol. 3, AIAA, New York, 1980, pp. 1918-1924.

9. V. C. Mueller, "Failure Analysis of Nickel Hydrogen Cell Subjected to Simulated Low Earth Orbit

Cycling," The 1983 Goddard Space Flight Center Battery Workshop, D. Baer and G. W. Morrow,

eds., NASA CP-2331, 1983, pp. 523-538.

10. K. M. Abbey, and D. L. Britton, "Electrolyte Management in Porous Battery Components - Static

Measurements," NASA TM-83073, 1982.

11. J. J. Smithrick, "Initial Performance of Advanced Designs for IPV Nickel-Hydrogen Cells,"Power

Sources 11: Research and Development in Non-Mechanical Electrical Power Source. L. J. Pearce,

ed., International Power Sources Symposium Committee, Leatherhead, England, 1986, pp. 215-226

(Also, NASA TM-87282).

12. J. J. Smithrick, and S. W. Hall, "Effect of KOH Concentration on LEO Cycle Life of IPV NickelHydrogen Flight Cells - An Update," Proceedings of the 26th Intersociery Energy Conversion

Engineering Conference, IECEC'91, D. C. Black, ed.,Vol. 3, American Nuclear Society, 1991,

PP• 276-281.

13. J. J. Smitli ick, and S. W. Hall, "Effect of LEO Cycling on 125 Ah Advanced Design IPV NickelHydrogen Flight Cells - An Update," Proceedings of the 25th Intersociery Energy Conversion

Engineering Conference, IECEC'91, D. L. Black, ed., Vol. 3, American Nuclear Society, 1991,

pp. 311-317.

14. D. F. Pickett, "Preparation of Nickel Electrodes," U. S. Patent 3,827,911, Aug. 1974.

115
15. J. J. Smithrick, M. A. Manzo, and O. Gonzalez- S anabri a, "Advance Designs for IPV NickelHydrogen Cells," IECEC '84: Advanced Energy Systems - Their role in Our Future; Proceedings of

the Nineteenth Intersociety Energy Conversion Energy Conference Vol. 1 American Nuclear

Society, New York, 1984, pp. 631-635 (Also, NASA TM-83643).

16. H. S. Lim, and S. A. Verzwyvelt, "Electrochemical Behavior of Heavily Cycled Nickel Electrodes

in Ni H2 Cells Containing Electrolytes of Various KOH Concentration," Proceedings of the

Symposium on Nickel Hydroxide Electrodes, D. A. Corrigan, A. A. Zimmerman, eds.,

Electrochemical Society Proc., Vol. 90-4, Electrochemical Society, 1990, pp. 341-355.

17. O. D. Gonzalez-Sanabria, "Effect of NASA Advanced Designs on Thermal Behavior of Ni-H2

Cells," 1988 IECEC; Proceedings of the Twenty-Third Intersociety Energy Conversion Engineering

Conference, D. Y. Goswami, ed., Vol. 2, ASME, New York, 1988, pp. 453-456.

18. H. S. Lim, G. R. Zelter, T. T. Smithrick, and S. W. Hall,"Destructive Physical Analysis Results of

Ni/H2 Cells Cycled in LEO Regime," Proceedings of the 26th Intersociety Energy Conversion

Engineering Conference, ECEC'91, D. L. Black, ed., Vol. 3, American Nuclear Society, 1991,

pp. 304-310.

116
TABLE 1. - CAPACITY OF HUGHES FLIGHT

CELLS CONTAINING 26 AND 31 PERCENT

KOH ELECTROLYTE

Cell Capacity,a

A hr

KOH 

concentration,

percent

1 59.0 31

2 59.9 31

3 59.0 31

4 53.8 26

5 53.2 26

6 52.3 26

-PRESSURE VESSEL --STACK WALL. WICK GAS SCREEN (NOT SHOWN) HYDROGEN OHYGASO SEPARATOR ELECTRODE RECIRCULATING STACK --' --- NICKEL ELECTRODE DROGEN

FIGURE 1 - ILLUSTRATION OF HUGHES RECIRCULATION

STACK INDIVIDUAL PRESSURE VESSEL NICKEL- a HYDROGEN CELL

Discharge at 1.4 C rate, 10 °C.

3. 2 4 5. 6 1. GAS SCREEN SEPARATOR HYDROGEN ELECTRODE BELLEVILLE SPRING NICKEL ELECTRODE 10. 9. 8 7. OXYGEN SEAL CATALYZED STRIP ZIRCONIUM OXIDE STRIP WALL WICK END PLATE

0 - FLOATING STACK - ACCOMMODATES NICKEL USE OF 26% KOH - IMPROVES CYCLE LIFE 10 X SOA SERRATED EDGE SEPARATOR - FACILITATES GAS MOVEMENT CATALYZED WALL WICK IMPROVES THERMAL AND OXYGEN MANAGEMENT ELECTRODE EXPANSION 109 - BACK-TO-BACK ELECTRODES-DIRECT PROPER STACK ELECTROLYTE ELECTROLYTE VOLUME TOLERANCE - MAINTAINS • COMPATIBLE WITH SOA AIR FORCE/HUGHES OXYGEN TO CATALYZED WALL WICK DESIGN - MINIMIZES DEVELOPMENT COST AND TIME

FIGURE 2 - NASA ADVANCED DESIGN IPV NICKEL-HYDROGEN CELL-CATALYZED WALL WICK

CURRENT COLLECTOR

GAS SCREEN

HOUSING HYDROGEN ELECTRODE

SEPARATOR

NICKEL ELECTRODE

ELECTROLYTE RESERVOIR PLATE (ERP)

CURRENT COLLECTOR

FIGURE 3 - CROSS-SECTIONAL UNIT OF THE NICKEL-HYDROGEN CELL

117
60

50

40

CAPACITY

A-HR 30

20

10

0

INITIAL CAPACITY

® CAPACITY AFTER 31 DAYS

CHARGE)

1 2 3 4 5 6

31% KOH CELL 26% KOH

FIGURE 4 - EFFECT OF STORAGE ON CAPACITY OF 48 A-hr HUGHES

IP Ni/11 2 FLIGHT CELLS.

1.4

w

a 1.3

J

J WJ 'Ji 1.2

U

w

t9

4

2

> 1.1

a

1.0 0 10 20 30 40 50

DISCHARGE TIME, MIN.

FIGURE 5 - COMPARISON OF HUGHES 49 A-hr IPV NUHy FLIGHT CELLS CONTAINING

26 PERCENT AND 31 PERCENT KOH ELECTROLYTE

60

a

118
1.5

1.4

269E KOH, 10-C

O CELL 4

o CELL5

n CELL6

1.3

CELL VOLTAGE

END OF 1.2 DISCHARGE

1.1

1.0

0.9

0 2000 4000 6000 8000 10000 12000 14000 15000 18000 20000 22000 24000 26000 28000 30000

CYCLENUMBER

FIGURE 6 - EFFECT OF LEO CYCLING AT 80 PERCENT DOD ON 48 A-hr IPV HUGHES

FLIGHT CELLS CONTAINING 26 PERCENT KOH ELECTROLYTE, 10-C.

1200

O CELL4 1100 O CELLS

0 CELL 6

1000

900

END OF 800

CHARGE

PRESSURE 700 PSI

600

Leo

Soo

400

300

0 2000 4000 6000 8000 10000 12000 14000 16000 18000 20000 22000 24000 26000 28000 30000

CYCLE NUMBER

FIGURE 7 - EFFECT OF LEO CYCLING AT 80 PERCENT DOD ON 48 A•hr IPV

HUGHES FLIGHT CELLS CONTAINING 26 PERCENT KOH.

119
1.4

1.3

CELL VOLTAGE 12

END OF

DISCHARGE

1.1

1.0

0.9

31 % KOH. 1 O'C

o CELL 4

o CELL 5

o CELL 6

1.5

0 2000 4000 6000 B000 10000 12000 14000 16000 18000 20000 22000 24000 26000 28000 30000

CYCLE

FIGURE 8 - EFFECT OF LEO CYCLING AT 80 PERCENT DOD ON HUGHES FLIGHT

CELLS CONTAINING 31 PERCENT KOH ELECTROLYTE, 10'C.

CELL 1.1,4C RATE

1.4

w

c^

a

b' 1.2

J

J

w

U

1.0

0 10 20 30 40 50 60

DISCHARGE TIME, MIN.

FIGURE 9 • COMPARISON OF HUGHES 48 A•hr IPV NUH FLIGHT CELLS

CONTAINING 31 PERCENT KOH ELECTROLITE

120
1200

1100

R' I0oo

Ld

Y^

soo

W

LL

d 600

W

a

S

a 700 SULL

O 600

0

Z

W 500

400

300

0 2000 4000 6000 8000 10000 12000 14000 16000 16000 20000 22000 24000 26000 28000 30000

CYCLE NUMBER

FIGURE 10 • EFFECT OF LEO CYCLING AT 80 PERCENT DOD ON HUGHES

FLIGHT CELLS CONTAINING 31 PERCENT KOH.

2.0

1400

1.6

ya 1200

ui (D 1.2

¢ a

^ J

N

al000

.e ujU

W

U

B00 .4

600 0'

TIME, MIN

54 90

FIGURE 11 • CELL VOLTAGE AND PRESSURE DURING CHARGE AND DISCHARGE FOR A REPRESENTATIVE

125 A-hr ADVANCED CATALYZED WALLWICK IPV NUH 2 FUGHT BATTERY CELL

170

t

a 160

a

a

a

U

J 150 W

U

140

C2 C 1.4C 2C

DISCHARGE RATE

FIGURE 12 - COMPARISON OF EAGLE-PICHER 125 A-hr NUH 2 CELLS

CATALYZED AND NONCATALYZED WALLWICK

121
INITIAL

® CAPACITY AFTER 52 DAYS STORAGE

(DISCHARGE, Voc. O'C) 180

160

140

L

Q 120

U 100

Q

Q 80

U

60

40

20

0

CATALYZED WALL WICK

100

rr

Z 80 WUQW

a

Z 60

O

F

Z

w

ij 40

cc

r

U

20

Q

U

0

1 2 3 4 5 6

CELL

FIGURE 13 - CAPACITY RETENTION OF 125 A-hr EAGLE-PICHER ADVANCED

IPV NVH 2 FLIGHT CELLS AFTER 72 hr OPEN CIRCUIT STAND

1 2 3 4 5 6

CATALYST NON-CATALYST

CELL

FIGURE 14 - EFFECT OF STORAGE ON CAPACITY OF 125 A-hr EAGLE-PICHER

ADVANCED IPV Ni/H 2 FLIGHT CELLS, CATALYZED AND

NONCATALI2ED WALLWICK, 26 PERCENT KOH.

122
1.5

O CELL 1

1.4 0 CELL 2

0 CELL 3

1.3

12

CELL VOLTAGE

END OF 1.,

DISCHARGE

1.0

0.9

0.8

0.7

0 2000 4000 6000 8000 10000 12000 14000 16000 18000 20000 22000

CYCLE NUMBER

FIGURE-15- EFFECT OF LEO CYCLING ON 125 A-hr NASA LEWIS ADVANCED CATALYZED WALLWICK IPV

N"H2 CELLS MANUFACTURED BY EAGLE-PICHER, 26 PERCENT KOH, 60 PERCENT DOD, 10-C.

1400

O CELL 1

q CELL 3

1300

1200

1100

CELL PRESSURE

END OF CHARGE

PSI 1000

900

800

700

0 2500 5000 7500 10000 12500 15000 17500 20000 22500 25000

CYCLE NUMBER

FIGURE 16 - EFFECT OF LEO CYCLING ON 125 A-hr NASA LEWIS ADVANCED CATALYZED WALLWICK IPV

Ni/H2 CELLS MANUFACTURED BY EAGLE-PICHER, 26 PERCENT KOH, 60 PERCENT DOD, 10'C.

123
1.5

O CELL 4

1.4 O CELL 5

Cl CELL 6

1.3

1.2

CELL VOLTAGE

END OF 1.1

DISCHARGE

1.0

0.9

0.8

0.7

0 2000 4000 6000 8000 10000 12000 14000 16000 18000 20000 22000 24000

CYCLE NUMBER

FIGURE 17 - EFFECT OF LEO CYCLING ON 125 A-hr NASA LEWIS ADVANCED NONCATALYZED WALLWICK IPV

Ni/H2 CELLS MANUFACTURED BY EAGLE-PICHER, 26 PERCENT KOH, 60 PERCENT DOD, 10'C.

O CELL 4

O CELL 5

0 CELL 6

CELL PRESSURE

END OF CHARGE

PSI

0 2500 5000 /b00 1UUW 1i^W 1tlJW 1/DW CWUU CLDw ZDuuu

CYCLE NUMBER

FIGURE 18 - EFFECT OF LEO CYCLING ON 125 A-hr NASA LEWIS ADVANCED NONCATALYZED WALLWICK IPV

Ni/H 2 CELLS MANUFACTURED BY EAGLE-PICHER, 26 PERCENT KOH, 60 PERCENT DOD, 10-C.

124
110

100

90

x 80 f

A

70

60

4096 DOD

L1pMweipnt nickel electrode:

Plaque: Fbrex Iber, 80 mil thkk, 909E porosity

Loadk>0 level 1.6 Vcc void

5- by 5-cm active area

50 l r 1 1 r i r

0 2000 4000 6000 8000 10000

CYCLE NUMBER

FIGURE 19 - UTILIZATION VERSUS CYCLE NUMBER OF A NICKEL-HYDROGEN CELL USING

A FIBEREX NICKEL ELECTRODE.

1.4

W

O

12 0 1.175 V 1.135 V

W

1.080 V

1.060 V

a

1

u.o 0 2000 4000 6000 8000 1000

CYCLE NUMBER

FIGURE 20 - END OF DISCHARGE VOLTAGE VERSUS NUMBER OF CYCLES FOR A

NICKEL-HYDROGEN CELL USING A FIBEREX NICKEL ELECTRODE.

125
Section

Two

Fuel Cells

P C t^!{^!£ PAGE BLANK NOT WMED
OZZ

r 4 }o N94ftw2 Ayr

PRIMARY AND SECONDARY ELECTRICAL SPACE POWER BASED ON

ADVANCED PEM SYSTEMS

N. E. Vanderborgh, J. C. Hedstrom and K. R. Stroh

Advanced Engineering Technology Group

Los Alamos National Laboratory

Los Alamos, New Mexico 87545

and

J. R. Huff

Ballard Power Corporation

Albuquerque, New Mexico 87109

INTRODUCTION

For new space ventures, power continues to be a pacing function for mission

planning and experiment endurance.

Although electrochemical power is a well

demonstrated space power technology,

current hardware limitations impact future mission viability. In order to document and augment electrochemical technology, a series of experiments for the

National Aeronautics and Space

Administration Lewis Research Center

(NASA LeRC) are underway at the Los

Alamos National Laboratory that define

operational parameters on contemporary

proton exchange membrane (PEM) hardware operating with hydrogen and oxygen

reactants. Because of the high efficiency

possible for water electrolysis, this

hardware is also thought part of a secondary battery design built around

stored reactants -- the so-called regenerative fuel cell. An overview of stack

testing at Los Alamos, and of analyses

related to regenerative fuel cell systems

are provided in this paper.

Finally, this paper describes work

looking at innovative concepts that remove complexity from stack hardware

with the specific intent of higher system

reliability. This new concept offers the

potential for unprecedented electrochemical power system energy densities.

STACK TEST PROGRAM

Although contemporary designs are

not space qualified, available PEM stacks

are being tested to ascertain the utility

of this fuel cell design for long term

space power applications. Tests are conducted that operate multi-kilowatt stacks

on hydrogen and oxygen. Experiments

are run that explore individual cell voltages, the effects of reactant stoichiometry, thermal management, and water balances during stack operation. To date,

two fuel cell stacks manufactured by

Ballard Power Systems, Vancouver, B.C.,

have been evaluated: 1) a stack using

DuPont NafionTM 117 electrolyte membranes; and 2) a similar stack formed

using experimental polyperfluorosulfonic acid polymer membranes manufactured by Dow Chemical Company.

Fuel cell stack testing requires the

establishment of a safe and regulated test

environment. Under NASA sponsorship,

Los Alamos designed and fabricated two

test stations to obtain reliable test data.

These test stations monitor inlet and

outlet parameters, and accomplish a full

mass and energy balance around the test.

One test station utilizes a passive load

bank to facilitate long-term, unattended

testing under rather continuous conditions. The second utilizes a computercontrolled electronic load that permits

programmed load changes. Because the

objective of these tests is to provide data

on stack life and performance decay

FI ft^aAC*46 PAGE BLANK NUT FUN4CI)

129
rates, long term testing is necessary.

The tests are fully automated with a

control computer that is able to diagnose

faults and, if necessary, terminate a

particular test.

The test configuration is shown

schematically in Fig. 1. The experiment

is designed to measure reactant flows,

product flows, current and voltage, and

heat flows. The determination of water

production rates and locations helps in

the evaluation of candidate membrane

materials used to form PEM stacks.

Electrochemical water production is frequently used to supply makeup water,

making this diagnostic measurement difficult. Water flows were partitioned in

the experimental design specifically to

determine stack water balances.

Electrochemical water production

occurs in the cathode compartment as the

result of two processes, hydrogen oxidation and water generation as the result of

electro-osmosis (transport of water from

the anode to cathode membrane faces as

the result of proton motion). Water

transport occurs through motion of a hydrated proton, H(H2O)n + . This type of

transport is frequently termed "water

drag." Water also moves from the wet

cathode to the anode through the membrane controlled by membrane mass

transfer rates.

The model predicts that appreciable fractions of water can be produced in

the anode compartment, depending on the

value selected for ks and n, where ks is a

Recirculation Humidifiers

Stack Heater Water Pump

Recirculation

Fan Oxygen

Inlet

f

Cooled

Condersor

Stack

Cooling Cathode Product

Water

Ion Exchange

Filter

Hydrogen Fan

Inlet

—►

Cooled

Condenser Stack

Anode

€1

Water

Makeup

Water

Fig. 1 Stack test schematic -- recirculation of both H2 and 02 and condensation

of product water allows full measurement of mass and energy flows.

Stack Performance Models- An

energy- and mass-balance computer

model has been developed, assuming

simple stack reactant and cooling flow

geometry. The partitioning of the

chemical hydrogen energy into heat and

electrical energy is modeled. This fuel

cell stack code also includes implicit

terms for membrane water transport.

water mass transfer coefficient that contains both diffusive and convective

terms, and n is the stoichiometry factor

that describes the average proton hydration number. Model development is discussed in Refs. 1-5. Additional discussion, related to the data presented below,

can be found in Ref. 1.

130
Stack Test Results. Early experiments showed relatively short stack

lifetimes. The cause of failure was typically excess cross over, that is, excessive

leakage anode to cathode resulting from

one or more failed membranes. Ballard

worked to correct these failures and

supplied a series of modified stacks with

improved membranes, electrodes, gaskets, and seals. A variety of stack

clamping designs have been tested. Each

time a stack was refitted to solve a known

technical problem, the clock was set to

zero and testing started on that new configuration. Initially failures occurred

near 200 hours; the last stack tested,

fitted with Nafion membranes, obtained a

test duration of 1250 hours.

Likewise, ancillaries have also been

improved during this test program.

Initially recirculation pumps used a

metal bellows design that typically

lasted less than 24 hours. A whole series of recirculation pumps for both hydrogen and oxygen have been evaluated,

and currently lifetimes in excess of 500

hours are predicted. The water circulation pump, used for thermal control, also

now shows useful, long lifetime. These

stack tests will continue until the 5,000

hour continuous operating goal is

achieved.

To date, two improved 4-kW stacks

have been tested -- the Ballard Mark V

"Nafion" and "Dow" stacks. The Nafion

stack tested has 46 single cells, using

graphite bipolar plates. Each cell has an

active area of 232.3 cm2 (0.25 ft 2 ). The

higher performance of the Dow membrane

allows that stack to contain only 40 cells.

The stacks are water cooled through the

use of integrated cooling hardware.

Cooling water also flows through the

humidification section that forms one

end of the stack.

Some representative data from these

tests can be found in Table I. These data

were obtained operating the stacks at a

constant current of 125 amps resulting

in a current density of 538 mA/cm 2 (500

amp/ft 2 ). Under these conditions, the

average cell voltage for the Nafion stack

was 0.752 V and was 0.799 V for the Dow

stack. The apparent electrical efficiency

(electrical energy/HHV fuel energy) was

50.7% for the Nafion stack and 54.1% for

the Dow stack.

Measurement of voltage performance under various load conditions with

time is straightforward and gave, in general, expected results. The successful

partition of chemical energy into electrical energy and exhaust heat routinely

showed internal and external leak rates

were very low on the improved stacks.

Additional stack testing and performance

model details can be found in Refs. 1 and

6.

POWER SYSTEM DESIGN

The preceding sections described

Los Alamos activities for NASA LeRC related to contemporary PEM fuel cell components. The following sections describe

system analyses and conceptual design

work related to complete fuel cell power

systems. The first section describes a

completed study of regenerative fuel cell

systems configured with evolutionary

hardware. The last sections present an

innovative concept for PEM fuel cell configuration and power system integration

that offers the potential for both higher

reliability (through use of fewer active

components and reduced complexity) and

unprecedented energy density in an

electrochemical power system.

Regenerative Systems, Los

Alamos has completed a technology assessment and trade-off study of fuel cell

and electrolyzer technologies suitable

for use in a regenerative fuel cell (RFC)

system for Project Pathfinder [Ref. 7, see

also Refs. 8-10]. The major components

of an RFC system can be seen in Fig. 2.

The technology assessment provides the

basis for selecting the most promising

candidate fuel cell and electrolyzer

technologies for the Pathfinder RFC energy storage system. The trade-off study

identifies the general characteristics of a

25-kW, 20,000-h energy storage system

utilizing the fuel cell and electrolyzer

131
technologies identified in the technology

assessment phase. The resulting system

incorporates PEM, metal-hardware fuel

cell stacks, PEM electrolysis stacks, and

compressed-gas reactant storage in filament-wound, lined-composite, high-performance tanks.

ment and system integration, there would

be no difference between the space and

terrestrial components. Total system

considerations would also be very similar. Only later would development efforts become application specific.

TABLE I Energy Balances and Water Balances from Operating PEM Fuel Cell Stacks

Energy Balance

Membrane

Type

Fuel Flow Reject Heat Stack Heat

Loss

Anode Heat

Loss

Cathode

Heat Loss

Electrical

Energy

Nafion 8532W

(HHV)

3581W 187W 165W 284W 4327W

Dow 7395W

(HHV)

3006W 160W 187W 287W 3997W

Water Balance

Membrane

Type

Water In

("makeup")

Product Water

(from H2 flow)

Anode Water

Flow

Cathode Water

Flow

Nafion 252 /h 1,935 /h 181 /h 2,027 /h

Dow 345 /h 1,674 /h 453 /h 1,633 /h

Energy

\ Source

Reactant

Storage

Rejection

Load

Chemical " Electro- Elec

Bus chemical Bus

Components

Fig. 2 Major components of a

regenerative fuel cell (RFC)

system.

During program deliberations, it

became apparent that some terrestrial

applications could profitably use an RFC

system. In the initial stages of developSimple, Modular PEM

Systems. A space power system requires a weight- and volume-constrained,

reliable and rugged power supply. Some

applications have the benefit of low

gravity (for example, on the lunar surface), but other applications must function in the microgravity of space.

Demonstrated contemporary electrochemical systems are quite complex

(especially in ancillary components

contrasted to electrochemical components), making the high reliability requirements of future space missions

quite challenging. However, new designs

offer the potential to eliminate heavy

bipolar plates and greatly simplify stack

thermal management, thereby reducing

system complexity. Moreover, such designs can, in concept, handle freeze-thaw

cycles with little problem.

Given the basic concept of a proton

exchange membrane (PEM) fuel cell op132
Hydrogen

Flow

Field

erating on hydrogen, we have developed a

concept l that has the potential for exhibiting the necessary and desirable

characteristics [Refs. 11 and 12].

Modularity permits both flexibility in

mission planning and increasing overall

reliability. Appropriate power system

designs tend to integrate into human

habitat modules, simplifying heat management in space environments.

The selection of compressed hydrogen storage in a composite pressure vessel presents the system designer with a

sizable cylindrical section. In considering ways to optimize the packaging of an

entire power system, a concept was developed that wraps the PEM fuel cell

components around the exterior of the

hydrogen tank. In this concept, the necessary flow fields, anode, membrane, and

cathode are assembled in an annular

configuration. This configuration is

modular, and useful voltages can be obtained by electrically connecting the anode and cathode of adjacent modules in

series. An efficient solid-state DC-DC

converter raises the "stack" voltage to

the desired system output voltage.

Higher system powers can be obtained by

incorporating additional annular PEM

modules.

An individual module consists of

an annular arrangement of PEM fuel cell

components. From the inside out there

are successive layers performing the

function of the porous electrocatalytic

anode, the proton exchange membrane,

and the porous electrocatalytic cathode.

Pressurized oxygen is recirculated over

the exterior of the cathode providing the

oxidant. Hydrogen is fed from the storage tank through a pressure regulator to

a plenum feeding the anode flow field.

The fuel cell stack is assembled on

the exterior of the hydrogen storage tank,

which completely supports the annular

Work performed for the US Army's

Belvoir Research, Development, and

Engineering Center, Fort Bel voir, VA.

fuel cell modules. As a result, the fuel

cell components can be very thin, providing high performance, low weight, and

ease of heat removal. System start-up

can be rapid because the low mass of the

cell allows the membrane and catalysts to

rapidly attain the optimum operating

temperature. The gas-diffusion portion

of the electrodes is constructed from

carbon felt embedded, if necessary, with

a wire-screen conductor. An optimized

membrane fabricated with integral thinfilm catalyst layers is hot-pressed between the carbon felt electrodes to complete the cell. No clamping forces or associated hardware are required. We envision fabricat?on of two carbon-felt

electrode sections on a continuous electrical conductor, where one electrode

section becomes the anode of the

"upstream" cell and the next electrode

section becomes the cathode of the

"downstream" cell. The continuous conductor electrically connects the anode

and cathode of adjacent modules. (The

transition must also maintain anode-toanode and cathode-to-cathode electrical

separation, continuity of the anode hydrogen flow field, and separation of the

anode and cathode gas flows.) The interconnection concept can be see in Fig. 3.

Oxygen

Flow

Field

Gasket

Layer

Cell Stacking Scheme

Fig. 3 Concept for interconnection

of modular annular fuel

cells.

The composite tank is wrapped with

a non-conducting layer to prevent elec133
trical connection between the individual

module anodes. This layer is fabricated

from a low-density, high-porosity, opencell polymer that acts as the anode flow

field. An additional layer of the same

material covers the exterior of the stack

assembly, mechanically protecting the

annular modules and preventing any exterior contact from shorting cathode-tocathode, while allowing the free flow of

oxygen, product water, and excess heat

through the interconnected porosity.

The recirculating cathode oxygen

flow serves to remove heat caused by fuel

cell inefficiencies from the stack and to

sweep product water from the system.

Subsequent heat exchange from the recirculating cathode flow will condense the

product water vapor to liquid where it

can be collected and drained to the electrolyzer water source. (In the absence of

gravity, centrifugal separation of the

liquid water will be required.)

A close-fitting shroud encloses the

system. This shroud channels the cathode flow, supports those components not

directly mounted on the tank, and provides the pressure boundary for the fuel

cell system. Some fraction of the waste

heat will dissipate through the shroud.

Ribs on the inside of the shroud contact

and retain the porous layer covering the

cathode. The use of composite material

for the shroud could provide exceptional

stiffness and strength with little weight

impact.

A connecting boss on the upper end

of the composite storage tank feeds the

hydrogen pressure regulator. The porous

annular anode flow field is connected to

the pressure-regulated hydrogen supply

by a plenum. Experimental results suggest that hydrogen humidification is not

required if cell current densities are

kept below about 500 mA/cm 2 . A connecting boss on the bottom of the tank locates a safety valve and a valved fill port

for refueling from the electrolyzer.

Because water will slowly accumulate in

the dead-end anode flow field, a periodic

purge of a "sump" is required, with a

corresponding valve and control circuit.

Safety considerations are expected

to require the incorporation of hydrogen

leak sensing and automatic shut-off of

the high-pressure gas supply. Hydrogen

sensors can be relatively simple and robust, consisting of a catalyst to enhance

hydrogen oxidation and a thermocouple

to sense that the exothermic reaction is

occurring. This function does not significantly complicate the microprocessor

control system.

Figure 4 displays the annular

module PEM H2-02 fuel cell concept.

When integrated with an electrolyzer, the

resulting configuration forms the reversible energy storage/generation portion of a regenerative power system.

Examplee System, To illustrate the

concept, we have performed some scoping

calculations on a 5-kW, 20,000 W-h

system. To increase the power density,

we have extended the design concept

(described above and in Refs. 11 and 12)

by "folding" the PEM membrane and electrode assemblies into two concentric

cylinders. Here, one mono-layer of cells

covers the hydrogen tank surface (as

above), but then "crosses over" and folds

back on itself with a continued monolayer of cells lining the inside of the

shroud/pressure boundary. The inner

and outer PEM cells are arranged so that

the cathodes face a common oxygen flow

field.

For reactant storage at 3,000 psi,

the resulting configuration has a shroud

outer diameter of approximately 36.3 cm

(14.3 in) and an overall length of about

133 cm (52.3 in). The corresponding

oxygen tank has an outer diameter of

25.4 cm (10 in) and an overall length of

88.9 cm (35 in). The calculated weight of

the tanks, reactants, fuel cell stack,

shroud/pressure boundary and ducting

is 28.7 kg (63.2 lb). Higher hydrogen

storage pressure will decrease volume,

but have little impact on overall weight.

Including a rough estimate for the weight

of those necessary components not included above, it appears that a primary

(but refuelable) "battery" of this design

could have an energy density (excluding

134
Hydrogen

Storage

adiator or

Meat Exchanger

Oxygen

Supply

^,..iuid Water

Storage

the heat rejection components) exceeding

500 W-h/kg (227 W-h/lb). Higher stored

energy will increase the energy density,

with the systemtem energy density approaching that of the reactant storage

subsystem. Figure 5 displays a crosssection through the example fuel cell

module.

Direct radiators are large, therefore, and

consequently heavy. In an application

where living or equipment space must be

conditioned, however, use of this "waste"

energy for heating, with subsequent rejection of excess heat from the larger

surface of the conditioned volume, may

prove attractive.

Fig. 4 Simple modular system concept using annular PEM fuel cells.

Subsystem OD

c 136.26 cm

Tank ID

Cathode

Flow Passage

Fig. 5 Horizontal cross-section

through example fuel cell

system showing relative

dimensions.

Heat rejection is a challenge for

this technology because the optimum operating temperature of the fuel cell components is no more than about 90°C.

Regenerative System Integration.

The concept and example above require

additional functions and components to

complete a regenerative system (see

Fig. 2 and Ref. 7). Electrolysis could

proceed at the reactant storage pressure

in a separate electrochemical component,

with a water pump raising the water from

the storage reservoir to the electrolyzer

pressure. Alternately, electrolysis

could occur in a reversible fuelcell/electrolyzer at the low fuel-cell operating pressure, with the product reactants subsequently pumped up to the required storage pressure. Each approach

involves challenges with high pressure

pumps and interfaces. Additional design

work and trade-off studies are required

to optimize the overall system in terms of

reliability, weight, and volume.

Concept Summary. In summary, the

series-integrated annular module PEM

H2-02 regenerative fuel cell system

030.97

135
consists of the following major

components:

broad uses in both the civilian and military sectors.

• A fuel cell subsystem consisting

of several annular PEM modules,

the anode and cathode flow-field

layers, the anode feed plenum,

and gasketing;

A hydrogen delivery system consisting of a high-performance

graphite-composite reinforced

compressed-gas storage tank with

a shut-off valve, pressure regulation, a safety valve, a valved fill

port, and an anode-flow-field

purge valve;

An oxygen delivery system consisting of a high-performance

graphite-composite reinforced

compressed-gas storage tank with

a shut-off valve, pressure regulation, a safety valve, a valved fill

port, a humidifier, a recirculation fan, a heat exchanger with

provision for product water collection, and a shroud/pressure

boundary;

• An electrolyzer system consisting of a high-pressure electrolyzes with necessary pumps,

valving, water storage, and photovoltaic array;

• An electronic package containing

DC-DC power conditioning, sensors, and controls;

• Mechanical supports; and

• A heat rejection radiator or connection with an integrated thermal management system (for example, providing cabin heat).

We believe that the development of

the annular modular configuration described above could lead to a system that

exhibits the necessary and desirable

characteristics for application to a variety of space or planetary missions. This

same technology would certainly have

FUTURE WORK

As part of the ongoing effort, the

Los Alamos National Laboratory will assist the NASA Lewis Research Center by

conducting an electrochemical cell test

program to provide critical life, reliability, and performance data currently

lacking in the data base of PEM fuel cells

and water electrolyzers. Los Alamos will

provide technical support for fuel cell

programs managed by NASA LeRC at the

discretion of the NASA program manager.

Los Alamos will also continue to develop,

in collaboration with NASA, the single

cell stand that was delivered to NASA

LeRC during Fiscal Year 1992.

REFERENCES

Vanderborgh, N. E., Kimble, M. C.,

Huff, J. R., and Hedstrom, J. C., "PEM

Fuel Cell Stack Heat and Mass

Management," Proc. 27th

Intersociety Energy Conversion

Engineering Conference (IECEC),

Vol. 3, pp. 407-411 (1992).

Nguyen, T., Hedstrom, J., and

Vanderborgh, N., "Heat and Mass

Transfer Design Issues in PEM Fuel

Cell Hardware," Proc. American

Institute of Chemical Engineers

(AIChE) Meeting, San Francisco CA

(November 1989).

3. Vanderborgh, N., Huff, J., and

Hedstrom, J., "Heat and Mass

Transfer Effects in PEM Fuel Cells,"

Proc. 24th Intersociety Energy

Conversion Engineering Conference

(IECEC), p. 1637 (1989).

4. Rieke, P., and Vanderborgh, N.,

"Temperature Dependence of Water

Content and Proton Conductivity in

PolyperfIuorosulfonic Acid

Membranes," J. Membrane Science,

Vol. 32, pp. 313-328 (1987).

Nguyen, T., Guante, J., and

Vanderborgh, N., "Ion Transport

136
Reactant Delivery Systems, and

System Integration for Individual

Power Sources," Los Alamos National

Laboratory LA-UR-93-345 Revised

(January 1993).

Through Polyperfluorosulfonic Acid

Membranes," Materials Research

Society Fall Meeting, Boston, MA,

(November 28-December 2, 1988).

Vanderborgh, N., Hedstrom, J., and

Huff, J., "Thermal Management of

Advanced Fuel Cell Power Systems,"

Proc. 25th Intersociety Energy

Conversion Engineering Conference

(IECEC), p. 149 (1990).

7. Huff J., Vanderborgh, N., and

Hedstrom, J., "Technology

Assessment and Trade-off Study of

Fuel Cell and Electrolyzer

Technologies for the Project

Pathfinder Energy Storage System,"

Los Alamos National Laboratory LAUR-90-3244 prepared for the

National Aeronautics and Space

Administration Lewis Research

Center (May 1991).

8. Huff, J., Hedstrom, J., Vanderborgh,

N., and Prokopius, P., "Regenerative

Fuel Cell Systems for Project

Pathfinder," Proc. European Space

Power Conference, Madrid, Spain

(October 2-6, 1989). Published by

ESA Pub. Div., Noordwijk,

Netherlands, pp. 217-219.

9. Vanderborgh, N., Hedstrom, J., and

Huff, J., "Electrochemical Energy

Storage Using PEM Systems," Proc.

European Space Power Conference,

Florence, Italy (September 1991).

10. Vanderborgh, N., Hedstrom, J., and

Huff, J., "Advanced Space Power PEM

Fuel Cell Systems," Proc. European

Space Power Conference, Madrid,

Spain (October 2-6, 1989).

Published by ESA Pub. Div.,

Noordwijk, Netherlands.

11. Vanderborgh, N. E., and Stroh, K. R.,

"High Energy Density Stack Design,"

Los Alamos National Laboratory

PATENT DISCLOSURE (December 11,

1992).

12. Vanderborgh, N. E., Stroh, K. R., and

Meier, R., "Analysis of Fuel Cells,

137
y3^v^ 3 P 16

J4®2333

PEM REGENERATIVE FUEL CELLS

Larry L. Swette, Anthony B. LaConti, Stephen A. McCatty

Giner, Inc.

14 Spring Street

Waltham, MA 02154-4497

ABSTRACT

This paper will update the progress in developing electrocatalyst systems and electrode structures primarily

for the positive electrode of single-unit solid polymer proton exchange membrane (PEM) regenerative fuel cells.

The work was done with DuPont Nafion 117 in complete fuel cells (40 cm electrodes). The cells were

operated alternately in fuel cell mode and electrolysis mode at 80°C. In fuel cell mode, humidified hydrogen

and oxygen were supplied at 207 kPa (30 psi); in electrolysis mode, water was pumped over the positive

electrode and the gases were evolved at ambient pressure. Cycling data will be presented for Pt-Ir catalysts and

limited bifunctional data will be presented for Pt, Ir, Ru, Rh and Na,^O, catalysts as well as for electrode

structure variations.

INTRODUCTION

In prior work 11, 2, 31, a large number of candidate bifunctional positive electrode catalysts were

evaluated for chemical and electrochemical stability and for catalytic activity in 30% KOH at 80°C. As a result

of this work, two potentially bifunctional catalyst systems were identified: 1) Na,^O, [4] and 2) metal/metal

oxide combinations of Rh, Pt and Ir [3]. More recently the scope of the program was expanded to include

development of bifunctional positive electrode catalysts for the regenerative PEM fuel cell, focusing on many of

the same catalysts.

In order to alternate between oxygen reduction and oxygen evolution on the same electrode, in addition to

bifunctional catalysis, it is necessary to have an electrode structure that can perform in both of these modes.

Some of the effort on this program was directed to the development of such structures. The approach taken was

to develop catalyst/binder compositions optimized separately at the particulate level for either oxygen reduction

(more hydrophobic) or oxygen evolution (more hydrophilic) and then combine these at an optimal ratio in a

single electrode, referred to as an "integrated dual-character" (IDC) electrode. In this approach the catalyst for

each function can be the same material if it shows bifunctional activity (e.g., Na "Pt3 04 ), or two different

monofunctional catalysts (e.g., Pt for 02 reduction and IrO2 for 02 evolution). This approach, developed for

the alkaline system, was extended to the PEMRFC by introducing a two-layer structure to achieve an

appropriate interface to the more confined two-dimensional character of the proton exchange membrane.

EXPERIMENTAL

Catalyst Materials

Pt black, used in both the positive and negative electrode structures, was fuel cell grade material obtained

from Engelhard Industries. Metal oxide catalysts were prepared by proprietary modifications of Adams-type

fusions of the metal salts in a nitrate flux [5]; metal/metal oxide catalysts (Pt-MO x) were prepared by co-fusion

of the metal salts in a similar manner followed by electrochemical reduction of the Pt0 2 to Pt, or were blends of

separately prepared materials (Pt-Na.Pt304).

Na^O4 was typically prepared by firing mixtures of Na2CO3 and Pt02 under oxygen at 650°C or higher

for several hours. In most cases it was necessary to use a proprietary activation procedure to develop significant

catalytic activity and acceptable conductivity.

139
The surface areas of all materials were determined by the BET nitrogen adsorption method using a

Micromeritics Flowsorb 11 2300 instrument. The electrical conductivities of materials were estimated by a

2-point method measuring the resistance of the powders under moderate compression [2]. The compositions of

the metal oxides prepared by nitrate fusion were determined in earlier work by X-ray diffraction analysis.

Iridium chloride fusions yield IrO 2 , which becomes slightly oxygen deficient Ir0 2.x (x = - 0.05) if electrochemically reduced. Rhodium chloride fusions yield Rh O 2 . The ruthenium salt fusions yield RuO2 , which also

becomes oxygen deficient RuO2_ x (x = - 0.1) if electrochemically reduced. The NaxPt3O4 preparations were

also analyzed by XRD for chemical characterization and determination of phase purity. The sodium content in

Na,^O4 can be determined from the unit cell parameter, a.:

x = (a. - 5.59)/0.11

Recent preparations of Na.Pt3 O4 have exhibited moderate surface areas and marginally acceptable conductivities

without any post-synthesis treatment and have been tested in that form. Materials characterization data is

summarized in Table 1.

PEM Regenerative Fuel Cell Testing

A rechargeable fuel cell requires 1) a negative electrode that is active for both H 2 oxidation and H2

evolution and 2) a positive electrode that is active for both O2 reduction and 02 evolution. The general approach

to achieve bifunctional electrode structures was to partially distribute the charge and discharge functions to

separate catalyst layers of a composite electrode each of which is optimized primarily for one function. A

hydrophilic PEM-bonded layer supports the gas evolution functions; for oxygen evolution the selected catalyst,

at a loading of 1-3 mg/cm2 , was blended with 10-20 wt% of the ionomer and pressed onto the surface of the

membrane. The gas-consumption discharge functions are supported by a separate more hydrophobic catalyst

layer which was physically pressed against the PEM-bonded layer by the mechanical compression of the cell

assembly.

Platinum is a good catalyst for the negative electrode reactions in both alkaline and acid electrolytes. We

used the composite electrode structure for most of the PEM cell configurations. For the negative electrode the

combination consisted of a platinum layer bonded directly to the PEM (optimized for hydrogen evolution) and a

separate IDC electrode (optimized for hydrogen oxidation). The "free-standing" component of the composite

electrode consisted of an IDC platinum composition bonded to wetproofed carbon paper.

The requirements for the positive electrode are more difficult to meet because of the scarcity of efficient

bifunctional oxygen catalysts. In addition, a titanium current collector is required because carbon support

materials are not stable under oxygen evolution conditions. Platinum is an excellent 0 2 -reduction catalyst for

the fuel-cell mode, but is relatively poor for oxygen evolution. IrO 2 is an excellent 0 2 -evolution catalyst, but a

poor 02 -reduction catalyst. For electrolysis cells it is mixed or alloyed with Pt to improve the conductivity of

the IrO2 . The general approach was to use Pt black for the oxygen reduction function in the free-standing

component and to incorporate the more favorable oxygen evolution catalyst in the PEM-bonded layer, with Pt

added to improve conductivity. For baseline performance data we used Pt or Pt-IrO 2 as the oxygen electrode

catalyst. Three baseline cell types were fabricated and tested: a Pt/Pt fuel cell (also tested as an electrolyzer), a

Pt/Pt-IrO2 electrolyzer, and several Pt/Pt-IrO2 bifunctional cell configurations; subsequently, Pt/Pt-Na,,Pt3O4

bifunctional cells of similar construction were also tested [6]. More recently, as the oxygen evolution catalyst,

we have evaluated Pt-RuO X , RuO X -IrO, and RhO2 , and we examined the performance of Pt-IrO. and Na,^O4

with repeated cycling.

The baseline Pt/Pt fuel cell and the experimental bifunctional cells were fabricated with a Nafion 117 PEM

and utilized composite electrodes (catalyst layer bonded to the PEM plus a free-standing IDC electrode) on both

sides. The baseline fuel cell and all of the experimental bifunctional cells had the same negative (H 2) electrode

configuration. On the oxygen side, the PEM-bonded electrode was Pt, Pt-IrO 21 Pt-RuOz , RuO,,-IrO,,, RhO2 or

Pt-Na,.Pt3 O4 , and the free-standing electrode was Pt or Pt-Na xPt3 O4 on titanium mesh (Pt-plated). The baseline

electrolysis cell had only the PEM-bonded electrode layers, optimized for gas evolution only (without ionomer

140
added). In electrolysis mode, the cells were run at ambient pressure with heated water pumped through the

positive electrode chamber; in fuel-cell mode the gases were supplied at 207 kPa and the H 2 stream was

humidified at 90-120°C. The majority of testing was done at a cell temperature of 80°C.

RESULTS

As described above, three different types of PEM cells were constructed and tested to establish a

performance baseline for single-function cells (fuel cell or electrolyzes) and a bifunctional cell. Subsequently,

bifunctional PEM cells with Pt-Na xPt3 04 positive electrodes were fabricated and tested to determine the relative

performance of this potential catalyst in acid electrolyte. More recently we have evaluated Pt-RuO,,, Rh0 2 and

RuO,,-IrO,, as the oxygen evolution components in bifunctional cells and investigated the performance of Pt-IrO.

and Pt-Na,^04 under limited cycling conditions.

Baseline Cells

Baseline Pt/Pt Fuel Cell. - In fuel-cell mode, this cell ran at 0.874 V at 100 mA/cm 2 and at 0.714 V at

500 mA/cm2 at 80°C. As an electrolyzer, performance was poor, as expected, primarily because Pt is not a

very good catalyst for O2 evolution; 1.77 V at 100 mA/cm2 , 2.00 V at 500 mA/em2 at 80°C. This cell was

provided with the same positive electrode structure as used in the bifunctional cells (an IDC platinum

composition at a loading of 8-10 mg/cm 2 on a Ti current collector mesh) so that the electrolyzer performance of

Pt could also be measured. More recently a non-reversible Pt/Pt cell, optimized for fuel cell operation, was

built and tested. In this cell the anode and cathode were of the same composition and configuration; the PEMbonded layers consisted of Pt at a loading of 1 mg/cm 2 with 10 wt %o ionomer, and the free-standing electrodes

consisted of an IDC composition of Pt at a loading of 4 mg/cm 2 on wetproofed carbon paper (Toray HO-30).

The fuel cell performance of these two cells is compared in Figure 1. It can be seen that the cell optimized for

fuel cell operation gave slightly better performance.

Baseline Pt/40Pt-60lrO- Electrolyzer. - This cell ran at 1.50 V at 100 mA/cm 2 and at 1.671 V at

500 mA/cm2 . This is fairly typical performance for a dedicated electrolyzer cell with Nafion and Pt/Pt-IrO2

catalysts. This cell was not tested in fuel-cell mode. The performance is presented in Figure 2 in comparison

to the baseline Pt/Pt fuel cell and the baseline Pt/Pt-IrO 2 bifunctional cell.

Baseline Bifunctional Pt/Pt-IrO-, Cell. - Several Pt/Pt-lr02 cells with varying IrO2 contents and

electrode structures were evaluated initially to optimize bifunctional performance [6]. A Pt/40Pt-60IrO 2 cell (#5

in Tables 1 & 2) cell exhibited the best performance in both modes. At 500 mA/cm2 , this cell ran at 1.587 V in

electrolysis mode and at 0.723 V in fuel-cell (#5) mode. The comparable voltages for the baseline cells are

1.671 V for the electrolyzer and 0.714 V for the fuel cell. The performance is presented in Figure 2 in

comparison to the baseline Pt/Pt fuel cell and the baseline Pt/Pt-Ir0 2 electrolyzer.

Catalyst Evaluation

Bifunctional Pt/Pt-Na,,Pt3 04 Cells. - Two different bifunctional cells were constructed for the initial

evaluation of Na,,Pt3 04 as a component of the oxygen electrode catalyst. The first cell (#6, Na,,Pt3 04 prep.#21,

see Tables 1 & 2) used a two-layer oxygen catalyst structure bonded to the PEM consisting of very hydrophilic

Na,^O, (10% ionomer) and a less hydrophilic Pt transition layer. The balance of the composite electrode was

a free-standing IDC electrode consisting of a 33/67 mix of Na^O, and Pt. This cell gave excellent fuel cell

performance (0.74 V at 500 mA/cm 2 ), as shown in Figure 2 and Table 2, better than the Pt/Pt cell (0.714 V)

and the best Pt/40Pt-60IrO2 cell (0.723 V). The cell performance in electrolysis mode was 1.697 V at

500 mA/cm2 , better than Pt/Pt (2.00 V), but not as good as the best Pt/40Pt-60IrO2 cell (1.587 V).

The second Na.Pt3 04 -containing bifunctional cell (#7, Na.Pt3 04 prep.#34A3) was constructed using a

single layer of Na.,Pt3 04 with 10% ionomer bonded to the PEM and the standard IDC Pt free-standing IDC

electrode as the composite oxygen electrode. This cell showed reduced performance in both modes compared to

141
the first Na^0, cell (Figure 2 and Table 2). In the fuel-cell mode, the voltage at 500 mA/cm 2 was 0.657 V.

In electrolysis mode, the voltage was 1.766 V at 500 mA/cm2 , which is still a significant improvement over a

pure Pt/Pt cell. The differences in performance between these two cells may be partly due to differences in

catalytic activity since the materials, although very similar in physical characteristics, are from two separate and

somewhat different preparations (see Table 1); a larger part of the differences, however, may be attributed to

the changes in electrode structure in Cell #7 resulting in a higher resistance interface.

Pt/Pt-ILO-,. - Cell #8 was an attempt to repeat the build of the baseline bifunctional Pt/Pt-IrO, cell (#5) but

with two slight differences: slightly more ionomer in the PEM-bonded layer and the introduction of a thin Pt

transition layer between the bonded layer and the free-standing electrode. The cell had somewhat higher

resistance and did not perform quite as well as Cell #5; the performance curves are shown in Figure 3.

RuO,,-IrO X was used as the PEM-bonded layer for oxygen evolution in Cell #9 and the Pt transition layer

was included for comparison to Cell #8. The electrolysis performance was better than for Cell #8 but the fuel

cell performance was not as good (see Figure 3). It is anticipated that a structural reconfiguration of the cell

duplicating Cell #5 would significantly upgrade fuel cell performance with this catalyst.

Pt-Ru% was used as the PEM-bonded layer for oxygen evolution in Cell #10 and the Pt transition layer

was again included for comparison. The performance was quite similar to the performance observed for RuO,,-

IrO, (see Figure 3).

RhO,, was used as the PEM-bonded layer for oxygen evolution in Cell #11 but the Pt transition layer was

eliminated. This material was difficult to work with because of the very high surface area and more development would be required for a fair evaluation. Two performance points are shown in Table 2, but the cell was

not very stable in either mode.

Repetitive Cycling Tests

Cells #12 and #13 were attempts to reproduce the electrode structures and performance of the baseline

bifunctional Pt/4013 t-60IrO, cell (#5) for the purpose of investigating tolerance to repetitive cycling. The first

build (#12) had somewhat higher than typical internal resistance and showed lower performance than Cell #5

(Table 2). The cell was subjected to six charge/discharge cycles at 500 mA/cm 2 , shown in Figure 4, after

which it developed a leak and testing was terminated.

Cell #13 came somewhat closer in performance to Cell #5; in fuel cell mode it was not quite as good as

Cell #5 but in electrolysis mode it was somewhat better (see Table 2). This build was subjected to fifteen

charge/discharge cycles, at which point it also developed a leak. The performance at 500 mA/cm 2 , shown in

Figure 5, remained fairly constant for both modes of operation over the fifteen cycles.

Cells #14, #15 and #16 were attempts to fabricate bifunctional positive electrodes similar to that used in

Cell #5 but based on Na X Pt3 0,, also for the purpose of investigating tolerance to repetitive cycling. The first

two builds (#14 & #15) did not give very stable performance and were not cycled (see Table 2). Cell #16 did

not show any significant improvement in performance but was somewhat more stable (see Figure 6) and was

subjected to five charge/discharge cycles. Neither the performance of Cell #5 nor the performance of the earlier

Na,^O, build (#6) was reproduced, but the cyclic performance was stable and better than Pt alone.

CONCLUSIONS

Regenerative PEM fuel cells showing efficient bifunctional performance can be fabricated with traditional

Pt and Pt-Ir catalysts if the electrode structures are properly designed.This investigation has also identified RuO,,

as a catalyst with good potential for improving the oxygen evolution component of a bifunctional positive

electrode. Na,,Pt3 0, was identified in our earlier work as a catalyst with potential for bifunctional operation in

142
both alkaline and acid electrolytes; the current work has confirmed this with limited charge/discharge cycling

but additional development will be required to achieve more reproducible catalyst properties and performance.

ACKNOWLEDGEMENTS

This phase of the work was supported by NASA Lewis Research Center under the direction of Dr. Patricia

Loyselle on Contract No. NAS3-24635. We would also like to acknowledge the assistance of Prof. B.L.

Chamberland of the University of Connecticut in the selection and identification of numerous materials tested in

this program.

REFERENCES

1. L. Swette and J. Giner, J. of Power Sources, 22 (1988) 399-408.

2. L. Swette and N. Kackley, J. of Power Sources, 29 (1990) 423-436.

3. L. Swette, N. Kackley and S.A. McCatty, J. of Power Sources, 36 (1991) 323-339.

4. L. Swette and N. Kackley, "Bifunctional Alkaline Oxygen Electrodes," IECEC Conf. Proceedinas, Boston,

MA, Aug. 4-9, 1991.

5. R. Adams and R.L. Shriner, J. Am. Chem. Soc., 45 (1923) 2171-2179.

6. L. Swette, N.D. Kackley and A.B. LaConti, "Regenerative Fuel Cells," Presented at the 27th Intersociety

Energy Conversion Engineering Conference, San Diego, CA, Aug. 3-7, 1992

143
TABLE 1

PEMRFC positive electrode: composition and characteristics of PEM-bonded layer

Cell # Catalyst Ratio Loading Ionomer Surface Area

(wt%) (mg/cm2) (wt%) (M2 /g)

5 Pt-IrOx 40-60 1 10 104

6 Pt overlay 33 1 0 28

Na.Pt3 O4 #21* 67 2 10 12

7 NaxPt3O4 #34A3* 100 2 10 7

8 Pt-IrOx 40-60 1 15 117

9 RuO-lrox 50-50 2 25 114

10 Pt-RuOx 40-60 1 15 69

11 Rh OZ 100 1 18 224

12 Pt-IrOx 40-60 1 10 117

13 Pt-IrOx 40-60 1 10 117

14 Pt black + 33 1 10 29

Nax Pt3 O4 #41 * 67 2 10 18

15 Pt black + 33 1 10 29

Na,,Pt3 O4 #43* 67 2 10 35

16 Pt black + 33 1 10 29

Na x Pt3O4 #43* 67 2 10 35

* Na,,Pt3O4 Characteristics

Prep Surface Area Conductivity Unit Cell Na,-value

No. (mz/g) (ohm-cm-1) Parameter

#21 12 33 5.671(3) 0.74

#34A3 7 30 5.674(1) 0.76 ("activated ")

#41 18 0.3 5.675(3) 0.77

#43 35 3.4 5.688(1) 0.89

144
TABLE 2

Cell voltage in fuel-cell and electrolysis modes for baseline cells and bifunctional PEM cells.

80°C, electrolysis at ambient pressure, fuel cell at 207 kPa.

Cell Performance in ---> Electrolysis Mode Fuel-Cell Mode

Positive Electrode 100 mA/cm2 500 mA/cm2 100 mA/cm2 500 mA/cm2

Red./Evol. Cat. (V) (V) (V) (V)

Baseline Fuel Cells

Pt/Pt(reversible) 1.770 2.000 0.874 0.714

Pt/Pt (ftiel cell) Not Tested Not Tested 0.901 0.733

Baseline Electrolvzer

Pt/40Pt-60Ir02 1.500 1.671 Not Tested Not Tested

Bifunctional Cells

#5 Pt/40Pt-6OIrO2 1.456 1.587 0.901 0.723

#6 Pt/Na.Pt 3

04 -Pt 1.545 1.697 0.903 0.740

#7 Pt/Na X N04 -Pt 1.621 1.766 0.839 0.657

#8 Pt/40Pt-60lr0„ 1.477 1.624 0.888 0.704

#9 Pt/50RuO,-50IrO,, 1.406 1.627 0.778 0.565

#10 Pt/40Pt-6ORuO,, 1.417 1.589 0.784 0.566

#11 Pt/Rh02 ---- 1.675 ---- 0.460

#12 Pt/40Pt-60IrOx 1.489 1.620 0.858 0.640

#13 Pt/4013t-60I0, 1.455 1.550 0.870 0.714

#14 Pt/33Pt-67Na.Pt 3 O4 ---- 1.873 0.873 0.698

#15 Pt/33Pt-67Na,,Pt 3 04 1.599 1.797 0.866 0.704

#16 Pt/33Pt-67Na .,Pt3 04 1.618 1.799 0.838 0.597

145
2.20

2.00

1.80

1.60

0

> 1.40

1.00

U

0.80

0.60

0.40

50 100 Current Density, mA/cm2 1000

Fig. 2.—Performance of bifunctional Pt/Na.Pt3 04 PEM cells.

Electrolysis: mode,

ambient Pressure, 8a*C :+:

+

+

+

0:

:O:

a 1 + Pt/Pt

Fuel Cell

• Pt/40Pt-60Ir

Electrolyzer

0 Pt/40Pt-6OIr

2 Cell #5 0

• Pt/NaxPt304

w

Cell #6

W I ID Pt/NaxPt304

Cell #7

1807 IEC92-4.TC

1

FC#2: Fuel Cell Only, C-ppr @ cath, Pt ionomer layers

0.9

W

anM

0 0.8

a^ #2 U #1

0.7

^ FC#1: Baseline reversible cell , Ti @ cath., no ionomer

0.100 500 1000

Current Density (mA/cm2)

Fig. I.—PEM Pt/Pt fuel cell performance (fuel cell mode only; 80°C, 30 psi).

146
1.8

1.6

1.4

v

M 1.2

j 1

v 0.8 U

0.6

0.4

gative: Pt/C-paper; Positive: Pt(Ti + Cat

100 500 1000

Current Density (mA/cm2)

F^EvolCat-

--w- Pt-lr#B -+- Pt-Ru#10 =+- Ru-try9 -E9-- Bslny5

Fig. 3.—Bifunctional oxygen catalyst evaluation (Nation 117, 80°C, 30 psi).

1.80

1.60

y

W 1.40

V

FnO 1.20

1.00

W

U

0.80

0.60

1 2 3 4 6 6

NUMBER OF CYCLES AT 500 mA/cm2

^— FUEL CELL —A— ELECTROLYSIS

Fig. 4.—Cyclic performance of Cell #/12 (Pt/Pt-IrO, 80°C).

147
1.8 0

—V—f—V—V—^—V,V—f—V —V—t—V-^,V ....

_ 1.60

>

W 1.4 0

V

H

a0

1.2 0

1.00

W

U

0.80

0.60

1 4 7 10 13 16

NUMBER OF CYCLES AT 500 mA/cm2

—^— FUEL CELL -A- ELECTROLYSIS

Fig. 5.—Cyclic performance of Cell #13 (Pt/Pt-IrOx , 800C).

1.8

6

1.6

1.4

v

1.2 egative: Pt/C-paper; Positive: Pt/Ti + Cat

1

v 0.8

0.6

04

100 500 1000

Current Density (mA/cm2)

- ^— NaPt0x#6 —+— #16-[nit E b16-Final

Fig. 6.—Cell #16 before and after five cycles (Nation 117, 80°C, 30 psi).

148
316zV

99511 N94-23359

HIGH POWER DENSITY PROTON EXCHANGE MEMBRANE FUEL CELLS

Oliver J. Murphy, G. Duncan Hitchens and David J. Manko

Lynntech, Inc.

7610 Eastmark Drive, Suite 105

College Station, Texas 77840

ABSTRACT

Proton exchange membrane (PEM) fuel cells use a perfluorosulfonic acid solid polymer

film as an electrolyte which simplifies water and electrolyte management. Their thin electrolyte

layers give efficient systems of low weight, and their materials of construction show extremely

long laboratory lifetimes. Their high reliability and their suitability for use in a microgravity

environment makes them particularly attractive as a substitute for batteries in satellites utilizing

high-power, high energy-density electrochemical energy storage systems. In this investigation, the

Dow experimental PEM (XUS— 13204. 10) and unsupported high platinum loading electrodes

yielded very high power densities, of the order of 2.5 W cm -2 . A platinum black loading of 5 mg

per cm2 was found to be optimum. On extending the three-dimensional reaction zone of fuel cell

electrodes by impregnating solid polymer electrolyte into the electrode structures, Nafion was

found to give better performance than the Dow experimental PEM. The depth of penetration of the

solid polymer electrolyte into electrode structures was 50-70% of the thickness of the platinumcatalyzed active layer. However, the degree of platinum utilization was only 16.6% and the

roughness factor of a typical electrode was 274.

INTRODUCTION

There is increasing interest in the development of fuel cell systems as power sources in

space, on land and for marine applications [1-3]. Since fuel cells bring about the direct conversion

of stored chemical energy to electrical energy without the intermediate generation of thermal

energy, they are not limited by the Carrot cycle, unlike heat engines [4]. Thus, energy conversion

efficiencies for fuel cell systems (45%-65%) are generally a factor of 2 higher than those for heat

engines, e.g., diesel generators. Although considerable advances have been made in fuel cell

systems over the past three decades [5,6], they are still considered as an emerging technology.

However, for manned spacecraft applications (Gemini, Apollo and Space Shuttle Orbiter),

relatively small (1 kW-12 kW) and highly reliable fuel cell power plants have been fabricated and

used successfully. Similarly, for terrestrial electrical power generation, relatively large systems

(40 kW-11 MW) have been built and demonstrated to be viable candidates for commercial power

generation [7].

Traditionally, satellite energy storage has been provided by nickel-cadmium batteries and,

more recently, by rechargeable nickel-hydrogen batteries [8], where the primary power source is

solar photovoltaic panels. However, with increasing satellite applications, and the associated

growing demand for electrical energy storage, in particular, for systems having high peak power

outputs, has led to a search for alternatives to rechargeable aqueous-based battery systems [8].

Operating conditions in respect to the charge regime and cycle life of satellite power sources have

become much more demanding, in particular, for long-lived, low earth orbit (LEO) satellites in

equatorial orbit [9]. For such an application, the number of cycles required of an electrochemical

energy storage system is 5,840 per year, or over 50,000 for a ten-year satellite lifetime. This

149
extremely high cycle life has led to the consideration of lightweight, efficient, high power density,

regenerative fuel cells for such satellite applications [8,10].

High power densities, and high energy conversion efficiencies, are a prerequisite for the

development of lightweight, small-volume and low-cost fuel cell systems having high specific

energies (Wh/kg) and high energy densities (AWI). At present, the only fuel cell systems that have

the capability of achieving high power densities (> 1 W cm-2) are the proton exchange membrane

(PEM) and the alkaline fuel cell (AFC) systems [11, 12]. Of these, the PEM system is the leading

candidate fuel cell-based power source for demanding satellite applications. The attractiveness of

this fuel cell system stems from a number of advantageous characteristics that it possesses,

compared to AFC systems. These include: (1) low-temperature operation; (ii) cold start-up

capability; (iii) negligible electrolyte management problems, since it uses a proton exchange

polymer membrane as an electrolyte; (iv) infinite life on open-circuit stand; (v) higher reliability;

(vi) can tolerate high reactant gas pressure differentials across the PEM electrolyte layer, and (vii)

ease of thermal and fluid management. Because of the attainment of high power density and high

energy conversion efficiency, together with the CO2-rejecting nature of the strongly acidic PEM

electrolyte layer, this fuel cell system readily satisfies the requirement of a dual-use fuel cell

technology, e.g., as the power source for electric vehicles where air, rather than oxygen, is used as

the oxidant feed gas.

Research and development areas presently being addressed by many research groups, in

order to make PEM fuel cells cost-competitive and more effective in meeting various mission

requirements, involve optimizing the structure, chemical composition and configuration of fuel cell

components [13-17]. These include: (i) electrode structures (platinum-based catalyst loadings and

thicknesses of gas diffusion and active layers) [13,18]; (ii) PEM electrolyte material (thickness,

conductivity, stability and mechanical characteristics) [11, 13,17]; (iii) preparation of membrane

and electrode (M&E) assemblies (hot-pressed membranes and gas diffusion electrodes, cast active

layers on membranes in physical contact with gas diffusion electrodes and electrodeposited noble

metal catalysts in the outer layers of PEM-impregnated gas diffusion electrodes) [ 13,15,19]; (iv)

lightweight, readily mass produced, bipolar plates, cell frames and gasket materials [20], and (v)

thermal and water management [16].

In the present investigation, various proton exchange membrane materials, together with

unsupported, high platinum loading electrodes (1-10 mg Pt cm 2), were used as the basic cell

components for carrying out performance evaluations in 5 cm2 active area single-cell test fixtures.

To increase the utilization of expensive platinum electrocatalysts, electrode structures were

impregnated with a solubilized form of various proton exchange membrane materials. Membrane

and electrode assemblies were prepared by a proprietary hot-pressing procedure. Humidified

oxygen and hydrogen gases at various pressures were used as reactants. Performance

characteristics (V-i curves) were recorded at various cell operating conditions (temperature and

reactant gas pressures). Using representative M&E assemblies, a cyclic voltammetric technique

was used to measure the degree of platinum catalyst utilization in porous gas diffusion electrodes.

Also, energy dispersive spectroscopy (EDS) was applied to the cross-sections of tested M&E

assemblies, in order to determine the thickness of active electrode layers and the depth of

penetration of solubilized PEM material into electrode structures.

EXPERIMENTAL

Electrodes And Proton Exchange Membranes

150
High surface area, pure platinum black electrocatalyst materials were used for both the

hydrogen and oxygen electrodes. The geometric area of each electrode was 5 cm 2. Commercially

available (Johnson-Matthey, Inc.) fuel cell gas diffusion electrodes, fabricated using unsupported

platinum black (surface area of 35-50 m2 g- 1 ), with platinum loadings of either 1, 2.5, 5 or 10 mg

cm-2 were used.

Three proton exchange membrane materials, two from Du Pont and one from Dow

Chemical Company, were investigated. Of these, only Nafion® 117 from Du Pont is

commercially available. The others, Nafion ® 115 and the Dow experimental membrane (XUS13204.10) were made available in small quantities for experimental purposes. The acid

membranes consist of a thin film of perfluorinated sulfonic acid polymer material that possesses

good chemical and thermal stabilities [21,22]. The structural formula for these cation exchange

membranes is given in Figure 1. Both Nafion® 117 and Nafion® 115 have equivalent weights of

1100 and thicknesses in the dry state of 175 gm and 125 μm, respectively. The Dow experimental

membrane has an equivalent weight of approximately 800, a thickness in the dry state of 125 μm

and a conductivity of approximately 0. 12 Q- 1 cm 1 , as measured with an AC bridge [23].

Before being used, PEM materials were washed in various solutions to remove trace

organic and inorganic contaminants. The purification procedure involved boiling the proton

exchange membranes in a 10 vol % aqueous H2O2 (30 vol %) solution for one hour, followed by

boiling for one hour in pure water and subsequently boiling for a second hour in a fresh sample of

pure water. Membranes were then removed and boiled for one hour in 0.5 M H2SO4, followed

by boiling for one hour each in two different samples of pure water. On completing the

purification procedure, membrane specimens were stored in the last sample of washwater at room

temperature prior to use.

A 5 wt % Nafion solution in a mixture of lower aliphatic alcohols and 10% water was

obtained from Aldrich Chemical Company. Portions of this solution (or a similar in-house

prepared solution of the Dow PEM material) were brush-coated onto the catalyzed surfaces of fuel

cell electrodes. On allowing the first coat to dry in air for a few minutes, a second coat was then

applied in a similar fashion. Afterward, the PEM-impregnated electrodes were placed in an oven at

80°C and allowed to dry for 45 minutes in an air atmosphere. From the difference in weight before

and after impregnation, PEM loadings (approximately 0.6 mg cm- 2) within the catalyzed electrode

structures were determined.

PEM-impregnated electrodes were bonded onto both sides of purified proton exchange

membranes, using a proprietary hot-pressing procedure. Pressed M&E's were immediately

removed from the hot press and mounted in a test fuel cell fixture.

Single Cell Test Fuel Cells

Test fuel cell fixtures were fabricated using machined graphite blocks, having rib-channel

patterns on one side, which facilitated the distribution of humidified gases to the porous gas

diffusion electrodes, as well as providing current collection from the electrodes. The active cell

area, represented by the rib-channel patterns, was 5 cm 2. A membrane and electrode (M&E)

assembly, which is the principal component of a cell, was obtained on pressing porous electrode

materials onto both sides of a proton exchange membrane, using a proprietary method. Tefloncoated fiberglass gasket materials, placed on both sides of the membrane, provided sealing on

151
bolting the cell components together. Copper endplates enabled test cell fixtures to be bolted

together, and electrical cartridge heaters inserted into the walls of the copper plates allowed various

operating cell temperatures to be selected.

After allowing 24 hours to condition a new M&E assembly in a test fuel cell fixture under

predetermined gas pressures (usually 0 psi H2 and 0 psi 02) and a temperature of 50°C, the

performance characteristics of a cell were determined. This involved placing an electronic load, set

to operate in the constant load mode, between the terminals of a cell. Performance characteristics

(V-i curves) were then recorded at various operating conditions of a cell (temperature and

pressure). On completing performance characterizations for some of the test cells, the oxygen gas

inlet was removed from the cathode compartment and replaced by an inert gas, e.g., argon. The

cathode compartment was then purged overnight with argon, and cyclic voltammetric profiles for

the oxygen electrode were subsequently recorded at a scan rate of 50 mV s- I between the potential

limits of hydrogen and oxygen gas evolution at a cell temperature of 50°C with hydrogen gas at 1

atm flowing over the anode and reference electrodes, and argon gas at a pressure of 1 atm flowing

over the cathode. The latter profiles were recorded using an EG&G PAR Instruments model 362

scanning potentiostat, in conjunction with a Yokogawa model 3025 X-Y recorder.

Test Station Hardware

Fuel cell test stations used in this investigation had: (i) microprocessor control of hydrogen

and oxygen gas temperatures entering the fuel cell; (ii) control of gas pressures from atmospheric

to 200 psi; (iii) flow meters for the control and measurement of gas flow rates; (iv) electronic loads

with a maximum capability of 300 W, and (v) a data acquisition system, with appropriate

interfacing for the automated collection and processing of data. An important component of each

test station was an attached temperature-controlled environment chamber that housed the test fuel

cell fixture and the hydrogen and oxygen gas humidification vessels. The temperature of the

controlled environment chamber could be raised or lowered on activating an electrical heating pad

placed at the bottom of the chamber. Appropriate electrical connections to components within the

chamber were made by isolated electrical feedthroughs, mounted on the walls of the chamber.

RESULTS AND DISCUSSION

Nature Of PEM Electrolyte Material On The Performance Of M&E Assemblies

Since ohmic overpotentials are largely determined by the thickness and by the conductivity

of proton exchange membranes, the performance of M&E assemblies using membrane materials

thinner than Nafion® 117 (e.g., Nafion® 115) and alternative membranes having the same

thickness as Nafion® 115 but at least a factor of 2 higher in conductivity (e.g., Dow experimental

membrane, XUS-13204.10) were evaluated. The more highly conducting Dow PEM also

possesses better water retention characteristics than Nafion®.

Usingg unsupported high platinum loading electrodes (5 mg Pt per cm2) and segments of

Nafion 117, Nafion® 115 and the Dow experimental membrane, M&E assemblies were

prepared. V-i profiles for the resulting assemblies, tested at a fuel cell temperature of 95°C with

H2/O2 reactant gas pressures of 30 psi/40 psi, are presented in Figure 2. A considerable

improvement in PEM fuel cell performance was obtained on using the Dow experimental proton

exchange membrane, as compared with Nafion ® proton exchange membranes.

152
In the absence of mass transport overpotentials, and taking the hydrogen oxidation reaction

to be kinetically fast, the three V-i profiles presented in Figure 2 may be expressed by [24,25]:

V = VO – b log i – Rii (1)

where

VO = Vr + b log i0 (2)

Vr is the thermodynamic reversible potential for the fuel cell reaction, 10 and b are the exchange

current density and Tafel slope for the oxygen reduction reaction and Ri is the differential

resistance of the cell. When activation overpotential (primarily due to the oxygen reduction

reaction) and ohmic overpotential are the major sources of power loss, the data points

corresponding to these regions in the V-i profiles can be fitted to Equation [1], and the parameters

V0, b and Ri extracted using the method of least squares fit. Excellent agreement between the

experimental data points and simulated V-i profiles from the curve fitting was obtained up to

current densities where mass transport effects became noticeable. The values of V0, b and Ri for

the three M&E assemblies, represented by the V-i profiles presented in Figure 2, are summarized

in Table 1.

The values for V0, which correspond to the cell potential at a current density of 1 mA cm-2,

are similar for the three M&E assemblies. The values for the Tafel slope, b, are close to those

reported by other investigators [26]. The interesting result from the table is given by the values for

Ri, which decrease in an expected manner on taking into account the physico-chemical

characteristics of the Nafion® and Dow membrane materials [23,27]. Based on these values of Ri,

the exceptionally good performance obtained from the M&E assembly containing the Dow

membrane can be readily understood in terms of low ohmic overpotentials.

The value of 0.075 Q cm2 for the M&E assembly containing the Dow membrane consists

of the resistance across the membrane itself and contact resistances between other components in

the test fuel cell fixture. In the absence of a Dow PEM, the contact resistances for the identical fuel

cell fixture, and similar electrodes were measured at 95°C and yielded a value of 0.032 Q cm2.

Thus, the resistance across the PEM itself was 0.043 Q cm 2. On taking the reported value for the

specific resistance of the Dow PEM, 8 Q cm [23] and the actual measured thickness (45 μm) of the

Dow PEM in the hot-pressed M&E assembly yields a membrane resistance of 0.036 Q cm2.

Thus, using the developed proprietary hot-pressing procedure for the preparation of M&E

assemblies, ohmic losses across M&E assemblies approach their theoretically limiting values.

V-i profiles for an M&E assembly (Dow PEM and unsupported, high platinum loading

electrodes-5 mg Pt per cm2), fabricated using the proprietary hot pressing procedure and tested

using a fuel cell temperature of 95°C and H2/02 gas pressures of 30 psi/40 psi, 60 psi/70 psi and

160 psi/160 psi, are presented in Figure 3. As expected, increasing gas pressures gave rise to

better fuel cell performance due to the effect of pressure on electrochemical reaction rates.

Variations of power density with cell potential and with current density for this M&E

assembly tested at a fuel cell temperature of 95°C with H2/02 gas pressures of 160 psi/160 psi,

using both humidified and dry oxygen gas are presented in Figure 4(a) and 4(b), respectively. It

can be seen from the figure that the fuel cell performed surprisingly well on dry oxygen gas and

may be able to function satisfactorily in this manner, at least for short periods of time, e.g., under

153
pulse power conditions. More importantly, it is apparent from the data presented in Figure 4 that

extremely high power densities of the order of 2.5 W cm-2 were obtained at a cell potential of 0.48

V and a current density of 5.1 A cm-2.

Impact Of Electrode Structure On The Performance Of M&E Assemblies

To maximize electrocatalyst utilization in PEM fuel cells, previous researchers [26,28]

impregnated catalyzed electrode structures with solubilized Nafion® PEM material. This treatment

was found to enhance the three-dimensional nature of the active catalyzed layers of fuel cell

electrodes and was observed to give rise to a reduction in activation and mass transport

overpotentials [ 13,25,26]. Because of the improved fuel cell performance obtained on using the

Dow experimental PEM as the solid electrolyte film, it was decided to investigate if further

improvements in fuel cell performance could be obtained if solubilized Dow membrane solution

was impregnated into the catalyzed electrode layers instead of the Nafion® solution. Thus, M&E

assemblies were fabricated using the Dow experimental PEM as electrolyte and unsupported, high

platinum loading electrodes (5 mg Pt per cm- 2). However, for one M&E assembly, the electrodes

were impregnated with the commercially available Nafion ® solution, while the second M&E

assembly was constructed using identical electrodes but impregnated with in-house-prepared Dow

membrane solution [29,30].

V-i profiles for the two similar M&E assemblies tested using a fuel cell temperature of 95°C

and H2/O2 gas pressures of 100 psi/100 psi are presented in Figure 5. It can be seen from the

figure that the Nafion ®-impregnated electrodes gave better performance than the corresponding

Dow-impregnated electrodes. This was also observed from data derived from other similar M&E

assemblies. This unexpected result may be due to the fact that the lower equivalent weight of the

Dow experimental membrane with its associated higher sulfonic acid group content, giving rise to

increased water retention in the membrane, brings about flooding within the electrode structures,

which impedes the access of reactant gases to platinum catalyst sites. However, the higher — CF2

— content of Nafion® compared to the Dow PEM (cf., Fig. 1) facilitates a greater solubility of

dissolved oxygen in the Nafion® impregnated PEM electrolyte material [31 ]. The higher localized

concentration of dissolved oxygen at platinum particle/PEM electrolyte interfaces within electrode

structures can support higher oxygen reduction reaction rates accounting for the trend observed.

Hence, while the characteristics of the Dow PEM are most suitable when it is used as the proton

exchange membrane itself, they may not be as appropriate as Nafion® for impregnation within

electrode structures. Thus, it seems that the best combination for an optimized M&E assembly is

the use of the Dow experimental PEM as the electrolyte layer and dissolved Nafion ® as the

electrolyte material to be used for impregnating catalyzed electrodes, so as to increase the threedimensional character of fuel cell electrodes.

On operating PEM fuel cells at high power densities, most of the electrochemical reactions,

giving rise to high cell currents, will occur near the front surface of the catalyzed electrodes. A

simple analysis of electrode pore models shows that high IR drops within the length of pores

dictates that most of the current is localized in the outer catalyzed layers of electrodes. Hence, to

achieve high power densities, it is advantageous to localize a large fraction of the electrocatalyst

material in the outer layers of fuel cell electrodes. Therefore, platinum electrocatalyst loading

would be expected to have a significant effect on the performance of PEM fuel cells.

To investigate the effect of platinum electrocatalyst loading on PEM fuel cell performance, a

number of M&E assemblies were fabricated using the Dow experimental PEM and Nafion®-

154
impregnated, unsupported platinum electrodes with loadings of 1, 2.5, 5 or 10 mg Pt per cm 2 . V-i

profiles for the resulting assemblies, tested using a fuel cell temperature of 95°C and H2/O2 gas

pressures of 100 psi/100 psi, are presented in Figure 6. Fuel cell performance was observed to

increase with increasing platinum catalyst loading, reaching a maximum at a loading of 5 mg Pt per

cm2. This can be more clearly seen in the variation of power density with cell potential and with

current density, shown in Figure 7, for these assemblies. The dependence of maximum power

density on platinum catalyst loading was in the order:

5 mg cm-2 > 10 mg cm-2 > 2.5 mg cm- 2 > 1 mg cm-2

Relatively thick platinum-catalyzed active layers, as found in electrodes having a Pt loading of 10

mg per cm2, may impede the access of dissolved PEM material into the electrode structures or

introduce mass transport overpotentials, which hinder high fuel cell performance. The behavior of

the M&E assembly with the highest Pt loading electrodes at high current densities [cf., Fig. 61

supports these conclusions.

In the case of unsupported, high surface area platinum black catalyzed electrodes, active

layer thicknesses corresponding to a platinum loading of 5 mg per cm 2 are optimum for PEM fuel

cells yielding the highest performance. For aerospace applications where high performance and

reliability are of the highest priorit^, the cost associated with the platinum, approximately $200 per

kW at a current density of 1 cm [19], can be justified. However, for terrestrial applications,

e.g., fuel cell power sources for electric vehicles, this cost figure is too high and carbonsupported, low platinum loading electrodes will have to be used.

Scanning Electron Microscopy And Energy Dispersive Spectroscopy Of M&E Assemblies

Scanning electron micrographs (SEM's) at various magnifications of the cross-section of

an M&E assembly, fabricated using Nafion®-impregnated, unsupported, high platinum loading

electrodes (5 mg Pt per cm2) and the Dow experimental proton exchange membrane (XUS13204.10) are presented in Figure 8. The fibers in the weaves of the carbon cloth backing can be

readily seen in Figure 8(a). The bonding between the membrane and the electrocatalyst layers

appears to be excellent, and no degradation was observed after completing a short-duration lifetime

test (240 hours). The cross-section of the M&E assembly, shown in the SEM given in Figure

8(b), is centered on the interface between the proton exchange membrane and the catalyzed layers

of one electrode. Some of the locations along the cross-section, where energy dispersive

spectroscopy (EDS) elemental analyses were performed, are numbered on the micrograph.

Exposure of the electron beam spot on the proton exchange membrane during EDS data collection

brought about decomposition of the PEM polymer material [cf., Fig. 8(b)].

The three key layers of the M&E assembly, namely, the gas diffusion layer, the active

platinum-catalyzed layer and the PEM electrolyte film, were analyzed using EDS. Representative

energy dispersive spectrographs for each of these layers are presented in Figure 9. The gas

diffusion layer [cf., Fig. 9(a)] shows high-intensity elemental carbon and fluorine signals, arising

from the Teflon-bonded, high surface area carbon powder material used in fabricating this layer.

The weak elemental platinum signal observed is probably due to carryover of platinum particles

from the active layer on slicing the M&E assembly. The spectrograph for the active platinum

electrocatalyst layer [cf., Fig. 9(b)] is dominated by a high-intensity elemental platinum signal, as

expected. Also, in this figure, weak signals from carbon and fluorine (associated with Teflon

bonding of the platinum electrocatalyst particles) and sulphur are to be noted; the latter is probably

155
associated with Nafion® impregnated into the active platinum-catalyzed Iayer. In the spectrograph

derived from the PEM electrolyte film [cf., Fig. 9(c)], high-intensity signals were obtained,

corresponding to carbon, fluorine and sulphur. This is in agreement with a proton exchange

membrane consisting of a perfluorosulfonic acid polymer material. No elemental platinum

dissolved into the membrane during fuel cell operation, as evidenced by the lack of a platinum

signal.

Elemental peak height ratios as a function of distance across the cross-section of the M&E

assembly, shown in the electron micrograph presented in Figure 8(b), are given in Figure 10. The

profile derived from the Pt/C elemental peak height ratios shows that the thickness of the platinumcatalyzed active layer is approximately 15-20 Rm. Examination of the profiles associated with S/C

and F/C elemental peak height ratios, given in Figure 10, demonstrate that the impregnated

Nafion® does not penetrate the entire thickness of the unsupported platinum-catalyzed active layer.

A rough estimate of the Nafion® penetration depth is of the order of 50-70% of the thickness of

the platinum-catalyzed active layer. This depth of penetration of solubilized Nafion® into the

platinum-catalyzed active layer is likely to be sufficient for PEM fuel cell operation at high power

densities (cf., Figs. 3 and 4). However, for high energy conversion efficiencies at lower current

densities, new methods which facilitate greater penetration of solubilized Nafion ® into active

catalyzed layers need to be developed.

Electrocatalyst Utilizations And Roughness Factors For Unsupported, High Platinum Loading

Fuel Cell Electrodes

The charge associated with the formation of a hydrogen monolayer or the removal of a

monolayer of oxide film, measured using the cyclic voltammetric technique, was used to estimate

the degree of platinum utilization and the roughness factor of unsupported, high platinum loading

(5 mg Pt per cm2) fuel cell electrodes. Cyclic voltammetric profiles recorded for Nafion ®

-impregnated, unsupported platinum black electrodes in test fuel cell fixtures are given in Figure 11.

The voltammetric profiles are similar to those obtained for smooth platinum in aqueous acid

solutions [32]. The poorly resolved hydrogen adsorption and desorption regions, the double layer

region and the oxide film formation and reduction regions are easily discernible. From the

voltammetric charge associated with either the monolayer oxide film reduction peak or the adsorbed

hydrogen monolayer region, information can be derived concerning the electrochemically active

surface area and the degree of platinum electrocatalyst utilization under typical fuel cell operating

conditions. The relevant data is presented in Table 2.

It can be seen from the table that satisfactory agreement between the adsorbed hydrogen

monolayer charge and the oxide film reduction charge was obtained. From the known platinum

catalyst loading (5 mg Pt per cm2) and assuming spherical platinum particles, 85 A in diameter, a

maximum electrochemically active surface area was estimated. For the unsupported, high platinum

loading, oxygen electrode, comparison of the calculated and experimentally determined

electrochemically active areas or adsorbed hydrogen charges indicate that the degree of platinum

utilization was only 16.6% and that the roughness factor for this electrode was 274. Clearly, either

some of the surface of the platinum particles are not in contact with the Nafion®-impregnated

electrolyte material or not all of the surface is electrochemically active. The non-electrochemical

EDS technique tends to overestimate the thickness of the platinum-catalyzed layer taking part in

electrochemical reactions. Further investigations are necessary to resolve this issue.

156
CONCLUSIONS

To obtain high power densities in PEM fuel cells, activation, ohmic and mass transport

overpotentials were minimized by the use of optimized M&E assemblies.

2. The Dow experimental PEM (XUS-13204.10) gave much superior performance than

Nafion PEM materials, which are manufactured by Du Pont.

3. Extremely high power densities of the order of 2.5 W cm- 2 were obtained at a cell potential

of 0.48 V and a current density of 5.1 A cm- 2 on using the Dow PEM and unsupported,

high platinum loading electrodes (5 mg Pt per cm2).

4. Electrode structure factors, e.g., nature of the three-dimensional active catalyzed layer,

electrocatalyst loading, etc., had a significant effect on PEM fuel cell performance, that is,

in obtaining high power densities.

Elevated reactant gas pressures and fuel cell temperatures considerably enhanced fuel cell

performance.

6. There was no evidence of deterioration or degradation in any of the tested M&E assemblies

or other fuel cell components after completing fuel cell performance evaluations.

ACKNOWLEDGMENT

Proton exchange membrane fuel cell research and development work performed at

Lynntech, Inc. is supported by the U.S. Air Force (Contract No.: F04611-90-C-0095) and by the

U.S. Navy (Contract No.: N00024-91-C-4067). This support is gratefully acknowledged.

157
REFERENCES

1. M. Warshay and P. Prokopius, J. Power Sources, 29 (1990) 193.

2. R. A. Lemons, J. Power Sources, 29 (1990) 251.

3. W. H. Kumm, J. Power Sources, 29 (1990) 169.

4. J. O'M. Bockris and S. Srinivasan, "Fuel Cells: Their Electrochemistry," McGraw-Hill,

New York (1969).

5. S. S. Penner, Energy: The International Journal, 11 (1986) 1-229.

6. A. J. Appleby and F. R. Foulkes, "Fuel Cell Handbook," Van Nostrand Reinhold, New

York (1989).

7. J. J. Hirschenhofer, Mech. Eng., Sept. 1992, p. 82.

8. G. Halpert and A. Attia, in: "Proc. 24th IECEC," ed., W. D. Jackson, IEEE, 345 East

47th Street, New York, NY 10017, Vol. 3, p. 1429 (1989).

9. H. S. Lim and S. A. Verzwyvelt, J. Power Sources, 22 (1988) 213.

10. A. J. Appleby, J. Power sources, 22 (1988) 377.

11. K. Prater, J. Power Sources, 29 (1990) 239.

12. S. Srinivasan, J. Electrochem. Soc., 136 (1989) 41C.

13. S. Srinivasan, O. A. Velev, A. Parthasarathy, D. J. Manko and A. J. Appleby, J. Power

Sources, 36 (1991) 299.

14. D. S. Watkins, K. W. Dircks and D. G. Epp, U.S. Patent 4,988,583, Jan. 29, 1991.

15. M. S. Wilson and S. Gottesfeld, J. Electrochem. Soc., 139 (1992) L28.

16. A. P. Meyer, in: "Proc. 24th IECEC," ed., W. D. Jackson, IEEE, 345 East 47th Street,

New York, NY 10017, Vol. 3, p. 1619 (1989).

17. K. Straber, Ber. Bunsenges. Phys. Chem., 94 (1990) 1000.

18. M. S. Wilson and S. Gottesfeld, J. Appl. Electrochem., 22 (1992) 1.

19. E. J. Taylor, E. B. Anderson and N. R. K. Vilambi, J. Electrochem. Soc., 139 (1992)

L45.

20. S. M. Misiaszek, "Low Pressure Air Breathing Fuel Cell Stack Development," Paper on

Fuel Cell Developments at Energy Partners, West Palm Beach, FL 33407, Dec. 1992.

158
21. A. Eisenberg and H. L. Yeager, eds., "Perfluorinated Ionomer Membranes," ACS Symp.

Series, No. 180 (1982).

22. U.S. Patent No. 4,417,969 (1983), assigned to Dow Chemical Company.

23. G. A. Eisman, Extended Abstract, Int. Seminar on Fuel Cell Technology and

Applications," Netherlands, Oct. 26-29, 1987, p. 287.

24. M. Ratcliff, F. L. Posey, D. K. Johnson and H. L. Chum, J. Electrochem. Soc., 132

(1985) 577.

25. S. Srinivasan, E. A. Ticianelli, C. R. Derouin and A. Redondo, J. Power Sources, 22

(1988) 359.

26. E. A. Ticianelli, C. R. Derouin and S. Srinivasan, J. Electroanal. Chem., 251 (1988) 275.

27. R. S. Yeo and H. L. Yeager, in: "Modern Aspects of Electrochemistry," eds., B. E.

Conway, R. E. White, and J. O'M. Bockris, Plenum Press, New York (1985), Vol. 16,

pp. 437-504.

28. I. D. Raistrick, U.S. Patent 4,876,115 (Oct. 1989).

29. M. N. Szentirmay, L. F. Campbell and C. R. Martin, Anal. Chem., 58 (1986) 661.

30. R. B. Moore and C. R. Martin, Anal. Chem., 58 (1986) 2569.

31. S. Srinivasan, Personal communication, March 1993.

32. L. D. Burke and M. B. C. Roche, J. Electroanal. Chem., 164 (1984) 315.

159
TABLE 1. VALUES OF V0, b AND R- FOR M&E ASSEMBLIES CONTAINING EITHER

NAFION ® 117, NAFIONd 115 OR THE DOW MEMBRANE (XUS-13204.10)

AS POLYMER ELECTROLYTE

Reactant Fuel Cell

Polymer Pressures Temp. VO b Ri

Electrolyte Reactants (psi) CC) (V) (V decade- 1 ) (S2 cm2)

Nafion 117 H2/O2 30/40 95 1.012 0.052 0.208

Nafion 115 H2/O2 30/40 95 0.995 0.044 0.176

Dow Membrane H2/O2 30/40 95 1.013 0.053 0.075

TABLE 2. ROUGHNESS FACTOR AND DEGREE OF PLATINUM UTILIZATION FOR

NAFION-IMPREGNATED UNSUPPORTED HIGH SURFACE AREA

PLATINUM BLACK-BASED FUEL CELL ELECTRODE

NAFION-IMPREGNATED

ELECTRODE*

Experimentally determined adsorbed hydrogen charge, QH,exptl.

Experimentally determined oxide film reduction charge, QO,exptl.

Experimentally determined electrochemically active surface area,

ESA, from QH,exptl. (assuming 210 μC cm-2)

Calculated real surface area (assuming 85 Angstroms diameter

Pt particles)

Geometric surface area, GSA

Surface roughness factor [ratio of experimentally determined

electrochemically active surface area (ESA) to geometric

surface area (GSA))

Calculated adsorbed hydrogen charge, QH,calc. (assuming

85 Angstroms diameter Pt particles)

288 mC

566 mC

1,371 cm2

8,250 cm2

5 cm2

274

1,733 mC

Degree of electrocatalyst utilization, QH,exptl./QH,calc. 16.6%

*Unsupported high surface area platinum black-based electrode with a platinum loading of

5 mg per cm2.

160
CF2

CF2 nr

CF-C 

2 

F

OCF2 CF C 2 SO H

L12 J C F m

X

r^-1

C F2

CF2 n r

CF-OCF2 CF2 SO3H

I

C2

x^-^---^

Structure of Du Pont's Nafion Structure of Dow Experimental

n=6.6 and m =1 Membrane n=3.6-10

FIGURE 1. Structural formula for proton exchange membranes.

H2/02 : 30psi/40psi

T:95°C

C^

Q

Q^ "CQCc

C 

Dow (xuS-13204.10)

` qO1CC13Et7

cC0C

0\ Nafion®115 CC^CC^,i

Nafion® 11

1.0

0.8

J

H

W

0.6

Oa

-^ 0.4 J

W

U

0.2

0.0

0 1000 2000 3000 4000 5000

CURRENT DENSITY, mA cm-2

FIGURE 2. Fuel cell potential-current density profiles for M&E assemblies

containing Nafion O 117, Nafion @ 115 and the Dow PEM.

161
30 psi/40psi

> 0.8

J

a

w 0.6

0

a

0.4

W

U

0.2

160 psi/160 psi

.— 1p

N 2.5 ,E

2.0

r

H

1.5

z

w

1.0 cc

w

O 0.5

a

0.0

Humidified OZ

map R"WalnDry OZ

T=95°C

H 2/02:160psi/160psi

1.0 T=95°C

0.0

0 2000 4000 6000

CURRENT DENSITY,mA cm -2

FIGURE 3. Fuel cell potential current density profiles for a Dow PEM-based M&E

assembly tested at various H2/02 gas pressures.

2.5 NE

2.0

t 1.5

V)

z

w

0 1.0

w

0.5

CL

0.0

(a)Humidified O2

Dry O2

T_95°C

H2/02:160psi/160psi

0.2 0.4 0.6 0.8 1.0

CELL POTENTIAL, V

0 1000 2000 3000 4000 5000 6000

CURRENT DENSITY, mA cm -2

FIGURE 4. Variation of power density: (a) with cell potential and (b) with current density for

a PEM fuel cell supplied with humidified H2 gas and humidified or dry 02 gas.

162
t.t

t

0.9

> 0.8

Q

0.7

Z 0.6

Gz:

Q 0.5

G

.7 0.4

Lz7

U 0.3

0.2

0.1

0

H2 /Q :100Psi/IOOPsi

T=95°C

N AFION-IMPREGNATED

DOW-IMPREGNATED

0 2000 4000 6000 8000

CURRENT DENSITY, mA cm-2

FIGURE 5. Fuel cell potential-current density profiles for ME assemblies containing

Nafion-impregnated or Dow-impregnated high platinum loading

electrodes.

1.t

t

0.9

> 0.8

Q

0.7

Z 0.6

O 0.5

0.4

U 0.3

0.2

0.1

0

( ^) 1 mg Pt cm-2

(Q) 2.5 mg Pt cm 2

(+) 5 mg Pt ero2

(0) 10 mg Ptcm2

H2/02 :100PSUl00Psi

T=95°C

0 2000 4000 6000

CURRENT DENSITY, mA cm-2

FIGURE 6. Fuel cell potential-current density profiles for M&E assemblies

containing high platinum loading electrodes.

163
2.5 (a) F{2 /02

:100"'/100psl

T=95°C

N

E 2.0

U

1.5

cn

z

W 1.0 0

cc

W

p 0.5

Q.

0.0

0.0 0.2 0.4 0.6 0.8 1.0

CELL POTENTIAL, V

2.5 I (b)

N

2.0 E

>U

7

1.5

N

z

W 1.0 0

cc

W

0 0.5

a.

T=95°C

0.0

HY02:100ps1/100psi

0 1000 2000 3000 4000 5000 6000

CURRENT DENSITY, mA crri2

FIGURE 7. Variation of power density: (a) with cell potential and (b) with current density for a

PEM fuel cell supplied with humidified H 2 and O gases. M&E assemblies

consisted of the Dow experimental membrane anilunsupported high platinum

loading electrodes:(0)l mg Pt cnr 2 ; ( p) 2.5 mg Pt cn,2 ; ( ♦) 5 mg Pt cnt 2 ; and,

( n)10 mg Pt cm-2.

164
(a)

FIGURE 8. Scanning electron micrographs

of the cross-section of an M&E

assembly fabricated using

Nafion-impregnated

unsupported high platinum

loading electrodes and Dow

experimental PEM.

0 000 V5 - 312 10 240

FIGURE 9. Energy dispersive

spectrographs taken at three

different locations along the

cross-section of an M&E

assembly fabricated using

unsupported high platinum

loading electrodes and Dow

experimental membranes.

(a)----Gas diffusion — , _ --- __--

layer

ppF

?

0.000 v s - 20+8 10.2,

_(b)Pt electrocatalyst

layer

R

C F _ fj

fm VFS - 4096 10. 2

(c) _PEM electrolyte—.--

layer

--- -_ --^ _

it

r i

(b)

165
20

0.8

U

N

O

0.6

a

F-z

c^

`w

0.4 z

Y

Q

W

a

J

Q

0.2

z

w

w

J

W

U

O

Q 15

F

= U

wo

Y z 10

W U

aEL

Q

~z 5

W

2

W

J

W

0 0

50 40 30 20 10 0 10 20 30

DISTANCE, p m

FIGURE 10. Elemental peak height ratios as a function of distance along the crosssection of an M&E assembly fabricated using unsupported high platinum

loading electrodes and Dow experimental PEM.

FDIC

60

100 mvs-1

^ J t

30 i l \ 50 mvs-1

0

H

Q

-30

-60

I^

V \ ^

CATHODIC \ x`=500 C

0.100 0.425 0.750 1.075 1.400

ELECFRODE POTENTIAL, V (RHE)

FIGURE 11. Cyclic voltammograms for an M&E assembly fabricated using Nafionimpregnated unsupported high platinum loading electrodes (5 mg Pt per

cm2) and Nafion 117 PEM.

166
q3q025fys/V N94'23350

THE ELECTROLYTE CHALLENGE FOR A DIRECT METHANOL-AIR

POLYMER ELECTROLYTE FUEL CELL OPERATING AT

TEMPERATURES UP TO 200 °C

Robert Savinell, Ernest Yeager, Donald Tryk, Uziel Landau, Jesse Wainright

Dominic Gervasio, Boris Cahan, Morton Litt, Charles Rogers and Daniel Scherson

Case Center for Electrochemical Sciences

Case Western Reserve University

Cleveland, OH 44106

ABSTRACT

Novel polymer electrolytes are being evaluated for use in a direct methanol-air fuel cell operating

at temperatures in excess of 100 oC. The evaluation includes tests of thermal stability, ionic conductivity

and vapor transport characteristics. The preliminary results obtained to date indicate that a high

temperature polymer electrolyte fuel cell is feasible. For example, Nafion 117 when equilibrated with

phosphoric acid has a conductivity of at least 0.4 Q -1 cm-1 at temperatures up to 200 oC in the presence of

400 torr of water vapor and methanol vapor cross over equivalent to 1 mA/cm2 under a one atmosphere

methanol pressure differential at 135 oC. Novel polymers are also showing similar encouraging results.

The flexibility to modify and optimize the properties by custom synthesis of these novel polymers presents

an exciting opportunity to develop an efficient and compact methanol fuel cell.

INTRODUCTION

The direct methanol-air fuel cell is a very promising system for a wide spectrum of applications

where high efficiency and light weight are required. Methanol is attractive for fuel cell use for a number

of reasons; it is more convenient to store and handle than hydrogen or ammonia, and it offers a relatively

high energy density. The direct methanol fuel cell is favored over the indirect (reformed methanol)

system since the direct oxidation of methanol offers the promise of greater efficiency and simplicity with

lower weight and lower cost resulting from the elimination of the fuel reformer. Present direct methanolair fuel cell configurations are severely limited by the lack of sufficiently active catalysts for the methanol

anode, and to a lesser extent, the oxygen cathode. In addition, current methanol-air fuel cells are hindered

by the cross-over of methanol from the anode to the cathode, which results in a loss of efficiency via

chemical reaction of the fuel with oxygen and depolarization of the cathode.

The use of solid polymer electrolytes offers new opportunities to overcome the catalyst stability

and activity problems, provided the polymers are stable and retain reasonable ionic conductivity at

temperatures approaching 200 oC. The use of a polymer electrolyte membrane (PEM) at 200 oC should

allow for the identification of effective catalysts which lack sufficient activity at the lower temperatures

typical of conventional PEM fuel cells or that are significantly deactivated by the adsorption of phosphate

in high temperature phosphoric acid fuel cells. The higher temperatures will increase the catalytic

activity substantially, while depressing the self-poisoning effects associated with strongly adsorbed

intermediates. The polymer should have other properties such as high oxygen solubility and low

electrolyte adsorption which will promote better electrode performance. In addition, the use of a polymer

electrolyte as opposed to a liquid electrolyte should result in extended fuel cell life due to the elimination

of electrolyte movement within the electrodes, i.e. electrode flooding. A DARPA sponsored project is in

progress at CWRU with the objective of developing highly efficient methanol-air polymer electrolyte fuel

cells. This research is focused on 1) developing new ion-conducting polymers capable of operation at

200 oC, 2) more active and stable electrocatalysts for methanol oxidation, 3) high performance

electrocatalysts for oxygen reduction that are tolerant of methanol and its oxidation products, 4) optimized

electrode structures and 5) new high area electrocatalyst supports with greater oxidative stability than the

traditional carbon.

167
This paper will focus on recent results obtained with several promising solid polymer electrolytes.

In particular, the electrolytes are being evaluated with respect to three criteria; 1) thermal stability, 2)

proton conductivity, and 3) transport of methanol, oxygen, water and carbon dioxide. These evaluations

are being carried out under temperatures and envirorunents (humidity level, oxidizing or reducing

conditions) that are indicative of the anode and cathode conditions in an operating fuel cell.

RATIONALE OF APPROACH

The present commercially available sulfonic acid ionomers (e.g. Nafion 117, Dow 560, see

Figure 1) depend on water to solvate the protons generated by the ionization of the sulfonic acid groups.

When the polymers are subjected to temperatures above 100 oC at one atmosphere pressure, water is lost,

including the water solvating the hydronium ions. The membrane shrinks and conductivity is lost. The

use of a pressurized system can extend the useful temperature range, but at the expense of overall system

efficiency. For an operating temperature of 200 oC, the pressure required is too high to be practical. It

has been shown that the ionization of trifluoro methane sulfonic acid (CF 3 SO3H) can be produced by the

addition of phosphoric acid to yield P(OH) 4+ and CF3 SO3 - [1]. A similar effect occurs with the

equilibration of sulfonic acid ionomers in concentrated phosphoric acid. The phosphoric acid acts as a

Bronsted base, ionizing the strong sulfonic acid groups and solvating the proton in the same manner as

water. However, due to the relatively low vapor pressure/high boiling point of phosphoric acid, it would

not be expected that ionization and solvation would be lost at temperature above 100 oC. The use of

phosphoric acid has the disadvantage that H3PO4 or its dissociation products are strongly adsorbed on Pt

based catalysts, resulting in a loss of catalytic efficiency.

The equilibration of Nafion with phosphoric acid is one example of the focus of the polymer

synthesis effort within the DARPA drect methanol fuel cell project The goal of the synthesis effort is to i

produce a polymer that is stable at high temperatures, with a superacid functionality for conductivity and

enhanced catalytic activity, also containing an immobilized Bronsted base. One possibility is that the

immobilization of the base could be accomplished by bonding the base onto the polymer backbone. It may

also be possible to synthesize a polymer in which water is strongly hydrogen bonded, which would reduce

the activity of the water and allow water to be retained at temperatures above 100 oC.

Researchers at Clemson University are currently developing perfluorosulfonimide polymers

analogous to Nafion. Examples of some of the polymers currently under study are given in Figure lab.

One of the advantages of the sulfonamide functionality is that much greater flexibility in the design of the

polymer is possible. For example, the side chain length can be varied and one or more functional groups

can be included in the side chain. In this manner a wade range of ion exchange capacities can be

obtained, and the domain structure of the polymer can be varied to provide optimized conductivity and

vapor barrier properties. In addition, preliminary samples of these materials have been shown to be

soluble in dimethylformamide (DMF) and are readily cast into films. This behavior is in distinct contrast

to Nafion, which is only solublized with considerable effort.

Researchers at University of Iowa are currently developing perfluorophosphonic acid polymers

as shown in Figure lc. It is expected that these polymers will have excellent thermal stability, will be

melt formable and will have good conductivity resulting from the two ionizable protons on each side

chain. In addition, it is expected that this material will strongly hydrogen bond to water to provide high

temperature conductivity without the need for a low vapor pressure solvent such as phosphoric acid. Since

the functional groups are part of the polymer chain, they should be sufficiently sterically hindered to

prevent adsorption onto active catalyst sites. In addition, it may be possible to make co- or ter- polymers

of these materials with the perfluorosulfonimide polymers described above.

A third synthesis effort is underway at CWRU to produce films of polybenzimidazole (PBI)

which have been doped with a strong acid such as sulfuric acid and then reacted at high temperature to

produce a polymer with the acid anion bound to the phenyl ring on the polybenzimidazole. The protons

are ionized by the basic nitrogen atoms in the benzimidazole rings as shown in Figure Id.

Polybenzimidazoles are known to have excellent oxidative and thermal stability, and should be further

stabilized by the reaction with acid. Proton hopping between basic sites should provide ionic conductivity

168
and the barrier properties of these films should be better than that of Nafion since the benzimidazoles will

be a single phase material, unlike the nonpolar fluorocarbon/polar ionic biphasic nature of Nafion.

EXPERIMENTAL

Preparation of Polymer Electrolytes

Nafion 117 (acid form) membranes were obtained from duPont and hydrated by immersion in

boiling water for four hours. Following hydration, the samples were either left in distilled water until use

or were equilibrated with phosphoric acid. The phosphoric acid equilibration consisted of immersion in

85% H3PO4 at 150 oC for 12 hours. Following equilibration the samples were left in phosphoric acid

until use.

Polybenzimidazole was purchased from Celanese. Films were cast from a solution of the polymer

in dimethylacetamide using a Gardner knife. The films were then doped by immersion in 8M sulfuric

acid for 16 hours.

The perfluorosulfonimide membranes were prepared by Dr. D. DesMarteau and co-workers at

Clemson University. The details of the polymer synthesis have been given elsewhere [2]. These

membranes were hydrated and equilibrated with phosphoric acid using the same procedure as was used

for the Naflon 117 samples. The results discussed below are for the difunctional polymer (acid form)

shown in Figure lb. The ion exchange capacity of this material was 0.66 mg I . &

Four Probe Conductivity Measurements

The conductivity measurements were made with a four probe apparatus constructed at CWRU.

In this apparatus two Pt foil electrodes are clamped to the ends of a 3 cm (L) x 0.5 cm(W) sample. These

electrodes are used to pass current through the sample. The magnitude and frequency of the applied

current were controlled using a PAR 173 potentiostat/galvanostat and a Wavetek 186 signal generator.

Two platinum wires (0.5 mm dia.) spaced 1 cm apart, were used to measure the voltage drop in the center

of the sample. Voltage measurements were made at three or more current levels to ensure that the (i,v)

behavior was ohmic. The current was applied at two different frequencies, 100 and 1000 Hz. No

frequency dependence was observed. The entire apparatus was contained within a sealed stainless steel

vessel, which was placed inside an oven and connected to a gas manifold system so that the temperature,

pressure and composition of the gas phase in contact with the sample could be controlled. Measurements

as a function of water partial pressure above 100 oC were performed by partially evacuating the cell and

then injecting sufficient liquid water via a septum to yield the desired water partial pressure.

Vapor Permeability Measurements

The vapor permeability measurements were made using a technique based on ASTM D-1434. In

this experiment the membrane separates two thermostated chambers. In one chamber (the retentate) a

constant partial pressure (typically 15 psi) of the gas of interest is maintained. As the gas permeates

across the membrane the pressure in the second chamber (the permeate) is monitored. Both chambers

contain an equal background pressure of an inert gas. Assuming that the pressure in the second chamber

is small compared to that in the first chamber, the pressure rise is linear with respect to time and can be

directly related to a pseudo-steady-state flux of gas across the membrane. In our apparatus the membrane

area was n cm2 , and the volume of the second chamber, including the internal volume of the pressure

gauge, was 10 cm3 . A stainless steel screen was used to support the membrane on the low pressure side.

An MKS Baratron pressure gauge was used to measure the pressure independent of gas composition to

within one torr. A Hewlett Packard 7090A recording plotter was used to record the pressure as a function

of time. The same apparatus was also used to measure methanol fluxes in pervaporation. In the

pervaporabon experiment, one side of the membrane was contacted with pure, liquid methanol, while the

169
permeate cell on the other side of the membrane was initially evacuated. Following evacuation, the

permeate cell volume was sealed. As methanol permeates across the membrane, it evaporates upon

reaching the permeate volume. By measuring the pressure rise in the permeate volume, the methanol flux

can be determined as in the vapor phase permeability experiments. In order to prevent methanol from

condensing, the pressure in the permeate volume was never allowed to exceed the vapor pressure of

methanol at the temperature used.

RESULTS AND DISCUSSION

Conductivity

The conductivity as a function of temperature for phosphoric acid equilibrated Nafion 117

(Na.f1on/H3PO4) in three different environments is shown in Figure 2. The curves labeled 'Dry (pure N2)

and '370 torr H2O/400 torn NZ' are self explanatory. The third curve, labeled 'H20/MeOH/N2' is for a

vapor phase with a water partial pressure of 350 torr, a methanol partial pressure of 350 torn and a

nitrogen partial pressure of 100 ton, i.e., conditions similar to what might be found at the anode of the

fuel cell. From these curves it can be seen that the presence of water in the atmosphere significantly

increases the conductivity, and that the presence of water and methanol together increases the conductivity

still further. These results suggest that even above 100 oC, the conductivity is dependent on the solubility

of water in the membrane. Since the water solubility in Nafion is enhanced by the presence of methanol

[31, if the conductivity depends on the water content of the membrane, a higher conductivity would be

expected with methanol present, which is what was observed. The decrease in conductivity at higher

temperatures suggests that the water content decreases with increasing temperature.

The conductivity results for the Clemson sulfonimide polymer T' (shown in Figure lb)

equilibrated with phosphoric acid are presented in Figure 3 (Vs. temperature at 400 ton water partial

pressure) and Figure 4 (Vs. water partial pressure at 125 0Q. For comparison purposes the conductivity

of Nafion /H3PO4 under similar conditions is also shown. Although the reproducibility of the data for the

sulfonimide samples is not good, the results are promising with conductivities significantly higher than

Nafion being observed in some samples. The poor reproducibility may be the result of the sample

equilibration procedure, which was the same as that used for the Nafion samples and resulted in highly

swollen films. It may be possible to optimize the equilibration procedure for the sulfoninde material, to

yield more reproducible results.

The conductivity results for the highly doped polybenzimidazole material are shown in Figure 5,

as a function of temperature with a water partial pressure of 400 torr. The Nafion 117/11 31304 results

under the same conditions are again given for comparison. This sample contained approximately

250 mol% H2 SO4, i.e., about 2.5 H2 SO4 molecules to each repeat unit. With this doping level, the

conductivity approaches that of Nafion, being about 50% lower at temperatures above 160 oC.

The best conductivity results to date are summarized in Table I for films at 170 oC and ca.

400 torr water partial pressure in terms of conductivity and voltage drop for a fuel cell operating at

100 mA/cm2 . For comparison literature values are given for 96% phosphoric acid at 170 oC and water

equilibrated Nafion at 80 oC. While the conductivities obtained in this work are far below that of liquid

phosphoric acid, they are very comparable to that of Nafion under traditional low temperature conditions.

Permeability

The permeability of oxygen through a Nafion 117 film equilibrated with phosphoric acid, Nafion

equilibrated with water, a sulfonimide 'B' film equilibrated with water, and a PBI film doped with H2SO4

was measured using the closed volume technique. The results are summarized in Table 2a. The current

density column represents the current density equivalent to the flux across the membrane with a one

atmosphere pressure differential as a driving force. The equivalent current density was calculated as

i = nFN where n = 4 for the oxygen reduction reaction, F is Faraday's constant and N is the flux. For

comparison, literature values for oxygen permeation in phosphoric acid and in Nafion (water equilibrated

170
and dry) are also given. There is reasonably good agreement between our result for water equilibrated

Nafion and the literature value shown, indicating that our apparatus appears to be functioning properly.

The permeability of oxygen in the phosphoric acid equilibrated Nafion at temperatures above 100 oC is

similar to that in water equilibrated Nafion at 80 oC. This result indicates that the performance of the

cathode should not be hindered by poor oxygen transport to the electrode surface, which, as shown in the

table, is a limitation of liquid phosphoric acid. The sulfonimide B' material is orders of magnitude more

permeable to oxygen, most likely due to the fact that it was swelled greatly by the equilibration process.

As indicated above, a different swelling procedure may be necessary to obtain optimal results with the

sulfonimide polymers. The lowest oxygen permeability was obtained with the polybenzimidazole

material, which was expected to show good barrier properties.

The permeability of methanol vapor through Nafion 117 films equilibrated with phosphoric acid

and equilibrated with water was also measured. These results are summarized in Table 2b. The

equivalent current density was calculated as shown above, but taking n = 6 for methanol oxidation. The

results show that the permeability of methanol through the phosphoric acid equilibrated Nafion films at

temperatures above 100 oC is significantly lower than through water equilibrated Nafion at 85 oC. From

this preliminary result it appears that methanol crossover to the cathode will be less of a problem for a fuel

cell operating at the higher temperatures under consideration. Since the diffusion coefficient might be

expected to rise with temperature, the lower permeability is probably the result of a lower solubility for

methanol in Nafion at the higher temperatures. A lower rate of methanol crossover should result in

higher efficiency and a lower probability of cathode depolarization by methanol.

The results from the pervaporation experiments are given in Table 2c in terms of equivalent

current density. A literature value for methanol pervaporation through Nafion 125 is also given for

comparison. The pervaporation results are in relatively good agreement with the literature value cited.

The agreement gives added confidence to our beliefs that our permeation apparatus and experimental

procedures are producing valid results. These results show much higher permeation fluxes for the case

where the membrane is contacted by liquid methanol as opposed to the permeation when methanol vapor

is in contact with the membrane_ Thus the methanol crossover through phosphoric acid equilibrated

Nafion at temperatures above 100 oC may be considerably smaller than would be expected based on

conditions where liquid methanol is in contact with a Nafion membrane.

Thermal Stability

The thermal stability of the sulfonimide B' polymer was investigated using Differential Scanning

Calorimetry (DSC) as shown in Figure 6. An endothermic process starting at about 150 oC can be seen,

corresponding to the loss of water from the membrane. The thermal decomposition of the polymer can be

seen at temperatures above 320 oC. This result shows that under these conditions (N2 atmosphere,

heating at 10 oC/min) the polymer is stable in the temperature range of interest.

Thermogravimetric Analysis (TGA) has been used to study the thermal stability of

polybenzimidazole and to follow the reaction between PBI film and sulfuric acid. In Figure 7a the TGA

scan for an undoped PBI film in nitrogen is shown. The heating rate was 10 oC/min. The PBI film lost

water at about 80 oC and underwent thermal degradation at 550 oC. The TGA of PBI doped with

approximately 20 mol% sulfuric acid is shown in Figure 7b. This result shows a loss of water at 100 oC

and also showed two further weight losses starting at 330 oC and 415 oC. The first loss at 330 oC is

attributed to the reaction of the acid with PBI to generate sulfonic acid groups while the second loss is

probably due to loss of S02 or SO3 from the polymer. The weight loss due to thermal degradation has

been shifted to 600 oC, which indicates that the reaction with the sorbed acid improved the thermal

stability of the film.

171
CONCLUDING REMARKS

A research program is in progress at CWRU/Clemson University/University of Iowa aimed at

developing a highly efficient methanol-air polymer electrolyte fuel cell capable of operation at

temperatures up to 200 oC. Preliminary results obtained to date indicate that the electrolyte requirements

needed to realize this goal can be met In particular, three novel polymer electrolytes have shown stability

and sufficient ionic conductivity at the elevated temperatures envisioned.

ACKNOWLEDGMENTS

This work was supported under a University Research Initiative Grant from the Defense

Advanced Research Projects Agency, ONR contract # N00014-92-J-1848. The authors acknowledge the

collaboration of Professor D. DesMarteau of Clemson University and Professor D. Burton of the

University of Iowa.

REFERENCES

1. E. Yeager, M_ Razaq, D. Gervasio, A. Razaq and D. Tryk, in "Structural Effects in

Electrocatalysis and Oxygen Electrochemistry", Proceedings Volume #92-11,

The Electrochemistry Society, Pennington NJ, p. 440, 1992.

2. K. Sung and D. D. DesMarteau, Polym. Preprints, 33(2), 168, 1992.

3. S. Kato, K. Nagahama, H. Noritomi and H. Asai, J. Membrane Sci., 72, 31 (1992)

4. T. Springer, T. Zawodzinski and S. Gottesfeld, J. Electrochem. Soc., 138, 2334, 1991

5. D. Linden, "Handbook of Batteries of Fuel Cells", McGraw Hill, New York, 1984.

6. K. Klinedinst, J. Bett, J. Macdonald and P. Stonehart, J. Electroanal. Interface. Chem.,

57, 281, 1974.

7. T. Sakai, H. Takenada, N. Wakabayashi, Y. Kawami and E. Torikai, J. Electrochem. Soc.,

132, 1328, 1985.

172
Table 1

MEMBRANE CONDUCTIVITY

Membrane /

Electrolyte

Conditions Conductivity /

obm-lcm 1

Voltage loss at 100

mA/cm2

Reference

Nafion 117 / HBO 80 oC, 0.08 25 mV 4

Saturated HBO

Vapor

96 % H,;POA 170 oC 0.64 < 5 mV 5

Nafion 117 / H;PO4 170 oC, 0.04 50 mV this work

400 torn HO

Clemson U. B / H;PO4 170 oC, 0.05-0.10 20 - 40 mV this work

400 torn O

PBI / II12 SO4 170 oC, 0.02 100 mV this work

400 ton O

Table 2a

OXYGEN PERMEABILITY

Membrane/Electrolyte Temperature / oC Current Density L

mA/cm2

Reference

95% HPO4 120 0.04 6

95% H PO 150 0.08 6

Nafion 117 /35 % O 80 0.46 7

Nafion 117 /dry 80 0.06 7

PBI / SO4 80 0.08 this work

Nafion 117 / 11,7 0 80 0.6-0.8 this work

Clemson B ! 0 76 20-47 this work

Nafion 117 / H.,PO l 123 0.35 -0.45 this work

Nafion 117 / H-;PO4 1 150 1 0.25 -0.67 this work

173
Table 2b

METHANOL PERMEABILITY

Membrane Temperature / °C Current Density

mA/cm2

Nafion 117 / F170 85 2.9-3.3

PBI / SO j 85 3.1

Nafion 117 / HIPO, 135 0.5, 0.5, 0.4

3 samples)

Nafion 117 / H.;PO j 140 0.9

Nafion 117 / HIPO4 155 1.1

Table 2c

METHANOL PERVAPORATION

Membrane Temperature / oC Current Density / Reference

mA/cm2

Nafion 117 / HB O 25 10-18 this work

27-30

16

3 samples)

Nafion 117 / HBO 50 52 this work

388

(2 samples)

Nafion 125 / HBO / 50 290 3

(CH ) NH+

174
Figure 1

Example Structures of Polymer Electrolytes

a. Perfluorosulfonimide (monofunctional)

[(CF2 - CF2)m - CFCF21n

OCF2CFOCF2CF2SO2NS02CF3 1 1 U3 H

b. Perfluorosulfonimide (difunctional)

[(CF2 - CF2)m - CFCF2]n

OCF2jNFOCF2CF2SO21 S02(CF2)4SO2IS02 CF3

CF3 H H

C. Perfluorophosphonic acid

[(CF2 - CF2)m - i FCF2]n

0(CF2)xF0(OH)2

d. Polybenizimidazole (reacted with sulfuric acid)

HSO4 -

HSO4-

e. duPont Nafion

[(CF2 - CF2)m - i FCF21n

[OCF2CFJXOCF2CF2SO3H

CF3

f. Dow

[(CF2 - CF2)m - i FCF21n

OC1'2C1'2603H

175
120 140 160 180 200

Temperature / C

0.00

100

0.20

0.16

E

U

j 0.12 0

0.08

U

bi".

U 0.04

' H20/MeOH/N2 370 torr H2O/400 -- •-- Dry N2

torr N2

Nafion I I7/H3PO4 Conductivity in different environments.

-'^----------------.

s_

------------ n-----------

Figure 2

0.20

<' 0.16

U

0.12 0

0.08

U

0 0.04

U

0.00

100 120 140 160 180 200

Temperature / C

• Clemson B/H3PO4 '-+-- Clemson BMPO4 `--`- --- Nafion

117/1-13 PO4

Figure 3 Conductivity vs Temperature at 400 torr water partial pressure.

176
0.25

0.20

e

U

0.15

0

0.10

U

.Q

v 0.05

0.00

0 100 200 300 400 500 600 700 800

Water Partial Pressure / torr

• Clemson B/H3PO4 Clemson B/H3PO4 _._._.F.°_.. Nafion 117

H3PO4

Figure 4 Conductivity vs Water Partial Pressure at ca. 125 oC.

0.20

0.16

0.00

100

8

U

0.12

0

0.08

U

b

0 0.04 U

Figure 5

120 140 160 180 200

Temperature / C

• PBUH2SO4 ...... 91-

..._. Nafion l l7/H3PO4

Polybenzimidazole Conductivity, 400 ton water partial pressure.

177
5 . C

v

a

u

0

b

G

W

U

2.5

r-i

c0

U

6

0

30 70 110 150 190 230 270 310 350

TEMPERATURE / °C

Figure 6 DSC of sulfonimide polymer B'.

178
T*094"ba-s M

t

c

s

s

c

a

w

c

a

soc

9a

_ Is

0

ea

75

70

c

c

i

L

c

Figure 7a TGA of undoped PBI polymer.

Ta.perstur• M

Figure 7b TGA of Sulfuric acid doped PBI polymer.

179
(jsR N94-23` 61

ADVANCES IN DIRECT OXIDATION METHANOL FUEL CELLS

S. Surampudi, S.R. Narayanan, E. Vamos, H. Frank and G. Halpert

Jet Propulsion Laboratory

California Institute of Technology

Pasadena, CA 91109

A. LaConti and J. Kosek

Giner, Inc.

Waltham, MA 02154-9147

G.K. Surya Prakash and G.A. Olah

University of Southern California

Los Angeles, CA 90089-1661

ABSTRACT

Fuel cells that can operate directly on fuels such as methanol are attractive for low to medium

power applications in view of their low weight and volume relative to other power sources. A liquid

feed direct methanol fuel cell has been developed based on a proton exchange membrane electrolyte

and Pt/Ru and Pt catalyzed fuel and air/O, electrodes respectively. The cell has been shown to

deliver significant power outputs at temperatures of 60 to 90° C. The cell voltage is near 0.5 V at

300 mA/cm' current density and an operating temperature of 90° C. A deterrent to performance

appears to be methanol crossover through the membrane to the oxygen electrode. Further improvements in performance appear possible by minimizing the methanol crossover rate.

INTRODUCTION

Direct oxidation methanol fuel cells (DMFC) are attractive for several defense and transportation applications in view of their lower weight and volume compared to indirect fuel cells [1,21.

The weight and volume advantages of direct oxidation fuel cells are due to the fact that they do not

require any fuel processing equipment. Elimination of the fuel processor also results in simpler

design and operation, higher reliability, less maintenance, and lower capital and operating costs.

Further, direct oxidation fuel cells are projected to have rapid and multiple start up capabilities, and

the ability to easily follow varying loads.

Under a task sponsored by the Defense Research Projects Agency (DARPA), JPL, USC, and

Giner, Inc. are engaged in the development of direct methanol fuel cells for future defense applications. A near term objective of the program is to identify advanced catalysts and electrolytes and

demonstrate the DMFC technology at the cell level. This paper describes the progress made to date

on this effort.

ASSESSMENT OF CATALYSTS FOR THE OXIDATION OF METHANOL

Anodic oxidation behavior of methanol was investigated in half cells using both supported and

unsupported catalysts at a loading of 0.5 mg/cm'. In these experiments 0.5 M sulfuric acid was used

181

PftG& *G PAGE BLANK NOT RLMED
as the electrolyte along with 0.001 M perflurooctane sulfonic acid, and the fuel concentration was

1.0 M. These half cell experiments were carried out at 45°. Anodic oxidation characteristics of

CH2OH at various unsupported (Pt-Sn and Pt-Ru) and supported catalysts (Pt, Pt-Sn and Pt-Ru) are

given in Figures 1 and 2. It can be seen from these results that Pt-Ru is the most promising among

the state-of-the-art Pt based catalysts for the oxidation of methanol. This observation is in agreement

with the results reported in the literature [3,4]. All further studies were carried out using Pt-Ru

catalyst.

Improved kinetics of methanol oxidation on Pt-Ru catalyst were observed with an increase in

temperature, catalyst loading and methanol concentration. Figure 3 reveals a significant reduction of

polarization as the temperature is raised from 25 to 60° C. Throughout the current density range of 1

to 100 mA/cm2 the polarization is reduced almost 100 mV by this temperature increase. Figure 4

gives the results of the catalyst loading studies. Results show that the increase from 1 to 5 mg/cmZ

resulted in the reduction of almost 100 mV polarization throughout the current density range from 1

to 100 mA/cm2 . Increasing the methanol concentration also resulted in lower anodic polarization for

the oxidation of methanol (Figure 5).

EVALUATION OF ELECTROLYTES FOR THE OXIDATION OF METHANOL

Liquid Electrolytes

Boron trifluoride dihydrate, triflic acid, perfluoroethane sulfonic acid and perflurooctane

sulfonic acid (C 8 acid) were evaluated as candidate electrolytes for the oxidation of methanol.

Experiments were carried out in half cells using activated Pt-Ru gas diffusion electrodes obtained

from commercial sources. Pt-Ru electrodes were found to be unstable and dissolve in boron

trifluoride dihydrate electrolyte. Hence no further experiments ere carried out with this electrolyte.

Anodic oxidation characteristics of methanol in the other three electrolytes are given in Figure 6.

Some of the important findings of this study are: a) anodic oxidation of methanol in triflic acid is

comparable to its oxidation is sulfuric acid, b) perfluroethanesulfonic acid was found to be unattractive for the oxidation of methanol, and c) marginally improved oxidation behavior was observed with

perflurooctanesulfonic acid electrolyte.

Solid Polymer Electrolytes

Nafion, a polymer electrolyte membrane (PEM), is an attractive alternate to the liquid

electrolytes for liquid feed methanol cells for several reasons including: a) simplified design, b)

simplified assembly and operation, c) lower corrosion, and d) minimization or elimination of shunt

currents. The material has been used quite successfully in gas feed hydrogen/oxygen fuel cells.

For these reasons, Nafion was evaluated as an electrolyte in a liquid feed half cell. A

methanol-water mixture only (without acid) was introduced to the fuel compartment. Nafion 117 (Du

Pont) was used, and the anode catalyst was a supported Pt-Ru type that was bonded directly to the

Nafion. Figure 7 compares performance of this cell with the Nafion electrolyte to the previously

employed sulfuric acid electrolyte. Hence, Nafion was selected as the candidate electrolyte in all

further studies.

182
CELL DESIGN

Design Optimization

Two cell design options were considered for direct oxidation methanol fuel cells including: a)

gas or vapor feed design and b) liquid feed design. Initial considerations indicated that gas feed

design is preferable in that it could employ existing fuel cell type gas diffusion electrodes. Furthermore, the gas feed design could operate at the higher temperatures that were anticipated to be

necessary for high performance. Subsequent deliberations considered the fact that the liquid feed

design would not require a vaporizer and would therefore be much simpler in design and operation.

The other potential advantages of liquid-feed design are: a) elimination of complex water and thermal

management systems, b) multiple use capability of the methanol-water as the fuel, for humidification

purposes and as an efficient stack coolant, and c) significantly lower system size and weight. This

cell design does not suffer from the disadvantages of prior liquid-feed cell designs which employed

liquid electrolytes. The use of PEM eliminates the problem of troublesome shunt currents and also

eliminates problems associated with corrosion of cell components. Tests have shown that the PEM

does not degrade with operation and is suitable for continuous operation. Fuel catalysts were found

to exhibit improved performance with the solid electrolyte membrane. On this basis JPL began

development of liquid feed type direct methanol fuel cells.

Schematic Diagram of Liquid Feed Direct Methanol Fuel Cells

Figure 7 give a schematic diagram of the complete laboratory type liquid feed methanol

system employing the membrane electrolyte. The MEA (membrane-electrode assembly) consists of a

layer of Nafion electrolyte 7 mil thick with fuel and air/0 2 electrodes bonded to either side.

Electrode dimensions are 2 inch by 2 inch by approximately 10 mil thick. The MEA is positioned

between the machined portion of two graphite blocks. The machined area on each block is a

rectangular pattern with open channels (designated as the flow field) to allow flow of liquid or gas

across the electrode surfaces. Inlet and outlet ports communicate with the flow fields via holes drilled

into the carbon blocks and equipped with threaded fittings at the sides of the blocks. Stainless steel

support plates, with the same overall length and width as the carbon plates, are located on the back

surface of the plates. The stainless steel plates as well as the carbon blocks are drilled in their outer

perimeter to accommodate bolts that are used to compress the assembly for sealing and to provide

electrical contact between the electrode and un-recessed area of the flow field.

The methanol solution is introduced into the fuel compartment of the cell via a pump and then

returned to a fuel storage reservoir as shown in Figure 8. The end product, carbon dioxide, is

entrapped in the exit fuel line and released in the storage reservoir. Pressurized air or 0 2 is

introduced to the air compartment of the cell and vented without circulation. Heaters are located on

the outside surface of the cell to control sell temperature. Finally, the cell is equipped with a small

closed end hole to accommodate an internal thermocouple.

PERFORMANCE OF LIQUID FEED DIRECT METHANOL FUEL CELLS

Effect of Temperature

Voltage-current characteristics of the liquid feed direct methanol fuel cell were measured over

a range of temperatures with 2 M methanol as fuel and pure 0 2 as oxidant. Results are given in

Figure 9 in terms of operating cell potential versus current density. Each point represents an

183

0 3
essentially steady state voltage that was achieved after about 5 minutes of continuous operation at the

indicated current density. Inspection of Figure 8 reveals a marked increase in performance with

increase in temperature over the range of 30° C to 90° C. For example, at a potential of 0.55 V, the

current density outputs are 10, 45, and 140 mA/cm2 at temperatures of 30, 60, and 95° C respectively. Similarly, at a potential of 0.50 V, the current density outputs are 20, 110, and 260 mA/cm 2 at

30, 60 and 95° C. The trend of increased output with increase in temperature is in accord with that

exhibited by other fuel cells. The increased output at higher temperatures is attributed to a combination of factors consisting of a reduction of cell ohmic resistance, activation polarization, and

concentration polarization.

Effect of Methanol Concentration

The effect of methanol concentration on the cell performance was also determined. Three

separate runs were carried out at 60° C each with a different fuel concentration of 0.5 M, 2.0 M, and

4.0 M methanol. The effect of fuel concentration on overall cell performance is given in Figure 10 in

terms of voltage-current characteristics. Inspection of these results shown that a higher operating

current densities, higher cell voltage is obtained with 2 M methanol while somewhat lower outputs are

obtained with both the higher concentration, 4 M methanol, and the lower concentration, 0.5 M

methanol. At lower current densities 0.5 M methanol was found to provide higher cell operating

voltage than 2.0 M methanol. On this basis there appears to be an optimum concentration for

operation for different current densities. The optimum may be between 0.5 M and 2 M methanol.

The lower performance of the cell at fuel concentrations less than 0.5 M is probably due to the

concentration polarization effects. The poor performance of the cell at higher methanol concentrations was found to be due to the fuel crossover phenomenon. Support for the proposed impact of

crossover at high concentrations was shown in half cell studies on the oxygen electrode. It was found

that the OZ electrode performance is significantly lowered at higher methanol concentrations [5]. For

example, the 0, electrode potential dropped more than 100 mV at 100 mA/cm 2 as methanol

concentration was increased from 2 to 4 M methanol. This finding emphasizes the need to minimize

the crossover rate to improve performance of the 02 electrode and hence the overall cell performance.

Fuel Utilization Studies

In order to examine fuel utilization, a cell was set up and run continuously at constant current

with a finite amount, 200 ml of 1.0 M methanol solution in the circulation tank, without replenishment of the methanol. Initially the current was set at 1.875 A (75 mA/cm2), and the cell temperature

was held at 80°. Results are given in Figure 11 in terms of cell voltage versus percent utilization of

the fuel. The utilization was taken as output amp-hrs/theoretical amp-hrs (from the amount of

methanol and its electrochemical equivalent). Inspection of this figure reveals that voltage drops

sharply at 75 mA/cm'- when utilization approaches 60%. The sharp drop in voltage at this point is

believed to be associated with concentration polarization of the fuel electrode that is, in turn, due to

an inadequate supply of methanol to the electrode. The methanol supply is, in turn, limited by the

low methanol concentration at this point (near 0.5 M or less). The phenomenon is consistent with

prior half cell studies that revealed the onset of high polarization when methanol concentration

declines below this level.

PROBLEMS AND ISSUES

Performance of the liquid feed methanol fuel cells is already attractive for some applications

and is approaching the levels required for electric vehicle propulsion. With some improvements in

184
electrical performance, efficiency and cost, this system can indeed be considered a serious candidate

for electric vehicle applications. These improvements can be achieved by developing high performance anode catalysts, new membranes with reduced methanol permeability, methanol insensitive

cathode catalysts, and low cost materials (non-Pt based catalysts, membranes, bipolar plate materials,

etc.).

CONCLUSIONS

Some of the major finding of the study are:

1) Pt/Ru catalyzed electrodes are well suited for oxidation of methanol.

2) Performance of Pt/Ru catalyzed carbon electrodes increases with increased temperatures (25

to 60° C), increased fuel concentration, (0.5 to 2 M methanol) and increased catalyst loading

(0.5 to 5 mg/cm2).

3) A new liquid feed DMFC has been developed based on a proton exchange membrane

electrolyte, Pt/Ru catalyzed fuel electrode, and Pt catalyzed air/02 electrodes.

4) The new cell can deliver significant outputs in excess of 250 mA/cm2 at potentials near 0.5 V

at moderate temperatures (less than 190° F).

ACKNOWLEDGEMENT

The work described here was carried out at the Jet Propulsion Laboratory, California Institute

of Technology, in a program sponsored by ARPA under contract with the National Aeronautics and

Space Administration.

REFERENCES

1) S.D. Fritts and R.K. Sen, Assessment of Methanol Electro-Oxidation for Direct Methanol-Air

Fuel Cells, DOE Report PNL-6077 (1988).

2) S. Surampudi, S.R. Narayanan, E. Vamos, H. Frank and G. Halpert, Proceedings of the

Eighth Annual Battery Conference on Applications and Advances, Long Beach, CA (January

1993).

3) M.M.P. Jansen and J. Moolhuysen, Electrochim. Acta, 21, 869 (1976).

4) M. Watanabe, M. Uchida and S. Moto, J. Electroanal. Chem., 229, 395 (1987).

5) S.R. Narayanan, E. Vamos, S. Surampudi, H. Frank and G. Halpert, Extended Abstracts of

the Fall 1993 Meeting of the Electrochemical Society, in press.

185
I--,

0.6

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w 0.5

2

z

> 0.4 J

Q

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0.3

O

a0.7

1 M METHANOL

+ OS M SULFURIC ACID

45°C

PVC

Pt-Sn/C

Pt-Ru/C

0.2

UNSUPPORTED

0.8

1 M METHANOL

+0.5MH2SO4

0.7 40`C

0.6 J

O

w 0.5

z

N

J 0.4

Q

fz

w 0.3 F- O

a

0.2

0.1

Figure 1.

0 '

0.1 1 10 100

CURRENT DENSITY (m A/cm-2)

Evaluation of Unsupported Catalysts for Methanol Oxidation.

SUPPORTED

0.8

0.1 LOADING: 0.5 mg cm -2

Figure 2

0 `

1 1

0.1 1 10 100

CURRENT DENSITY (mA/cm')

Evaluation of Supported Catalysts for Methanol Oxidation.

186
U) t.._

_.I O

w

s

z

Cn

a

z

w

H

O

CL

0.8

0.7

0.6

0.5

0.4

0.3

0.2

1

Figure 3

Pt-Ru /C, 0.5mg/cm2

1 M METHANOL + 0.5 M H2 SO4

25C

10 100

CURRENT DENSITY (mA/cm2)

Effect of Temperature on Performance of Pt/Ru Electrodes.

45C

60C

()

O

w

z

U)

a

z

w

O

a

Pt- Ru/C, 0.5mg/cm2

1 M METHANOL + 0.5 M H 2 SO4

6000

0.5 mg/cm2

5.0 m g/c m 2

10 100

CURRENT DENSITY (mA/cm2)

0.8

0.7

0.6

0.5

0.4

0.3

0.21

1

Figure 4. Effect of Catalyst Loading on Performance of Pt/Ru Electrodes.
0.8

JO 0.7

w 0.6

z

u, 0.5

Q 0.4 F

z

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O

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00

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p 0.2

CL

Figure 5

Pt-Ru/C, 0.5mgICm2

0.5 MH2SO4

60 °C

1 M CH 3 OH

8M CH3 OH

10 100

CURRENT DENSITY (mA./cm2)

Effect of Fuel Concentration on Performance of Pt/Ru Electrodes.

1.0

Pt-Ru/C

0.9 1 M METHANOL

55°C

1 M McOH, Pt-Ru, 55 - 57 C

- 0 0.5M H 2 SO4 , 0.01 M C8-a e A

• 1 M CF3 SO3 H, 0.01 M C8-a e A

A 1 M C2 F5 SO3 H, 0.01 C8-a e A 0

.1

^

A 0 M C8 F17 SO3 H A °

0

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A

A ° CA ot 090

0i e e ° ° o°. o 0 aA, °t °Z

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—0

0.1 1 10

CURRENT DENSITY (mA/cm2)

Figure 6. Anodic Oxidation of Methanol in Various Electrolytes.

100

188
0.s (60°C, 1 M CH 3 OH)

0.7

v7J

O]

 0.6

w

Z

0.5M H2SO4.

0..5

J

Q

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a

0.3

NAFION

0.2 -_ I

0.1 1 10 100

CURRENT DENSITY (mA/cm2)

• IMPROVE FUEL OXIDATION CHARACTERISTICS WITH NAFION

Figure 7. Evaluation of Electrolytes for Anodic Oxidation of CH30H.

Carbon Dioxide

n (VENT)

n

n

n

n

n

CATHODE

PEM

ANODE.

n

n Oxygen/Water

n (VENT)

n

n

n Methanol/Water

n Carbon Dioxide ...'^.,^" (OUT)

Fuel

Reservoir Pump 

Methanol/

Water

-

Methanol/Water

(IN)

Oxygen

(IN)

Figure 8. Diagram of JPL Liquid Feed Direct Methanol Fuel Cell.

189
4.0 M METHANOL

i^ 0.6

HJ

O

> 0.5

w

c^

0.4 J

O

> 0.3 J

J

W

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2.0 M METHANOL

0.8

0.7

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Figure 9.

0.8

0.7

-- 0.6

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J

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w

a 0.4

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J

J

W

c) 0.2

100 200 300

CURRENT DENSITY (mA/cm 2)

Effect of Temperature on Cell Performance.

(60°C, 20 PSIG 02)

0.5 M METHANOL

0.1

0.0 L

0

Figure 10

100 200 300

CURRENT DENSITY (mA/cm2)

Effect of Fuel Concentration on Cell Performance.

190
0.6

0.5

H-^ 0.4

O

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a

J 0.2

O

> 0.1

0.0

0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7

FUEL UTILIZATION (PERCENT)

Figure 11. Voltage vs Fuel Utilization Curve.

Current Density, 62° C. 20 PSIG Oxygen.

1 M Methanol Initial Concentration,

75 mA/cm2

191
y39 627

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RECENT ADVANCES IN SPE® WATER LECTROLYZER

James F McElroy

HAMILTON STANDARD DIVISION

UNITED TECHNOLOGIES CORPORATION

WINDSOR LOCKS, CONNECTICUT 06096-1010

ABSTRACT

A new cell structure has been introduced into the SPE Water Electrolyzer which has

improved overall characteristics significantly. Weight, reliability, and efficiency are the

characteristics that are improved the most, with volume having a second order improvement. This

paper discusses the capabilities of the new cell structure and the impact it would have in various

space applications.

INTRODUCTION

In recent years, high pressure water electrolysis has been considered for several space

applications. These space applications include the generation of high pressure hydrogen and

oxygen for space station reboost propulsion, the generation of high pressure hydrogen and

oxygen for lunar base energy storage, and the generation of high pressure oxygen for on-orbit

space suit oxygen supply recharge. The technical approach taken with the SPE Water Electrolyzer

has been to use the previously developed nuclear submarine SPE Electrolyzer design while

reducing mass wherever possible. Because the basic nuclear submarine SPE Electrolyzer cell

design is limited to approximately 400 psi differential pressures overboard and across the cell, the

stack of cells must be pressure balanced in a nitrogen-filled dome when operating above 400 psi.

This results in significant mass and system complexity.

Within the last two years, a new SPE Water Electrolyzer cell structure has been developed.

This cell structure, which has a lower mass than the 400 psi design, requires no pressure-balanced

dome when operating at elevated pressures (i.e., up to several thousand psi). This cell structure

is configured to permit the hydrogen and oxygen compartments to operate at differential pressures

equivalent to the maximum system pressure and, finally, this cell structure results in lower cell

voltage (i.e., higher efficiency).

BACKGROUND

SPE Water Electrolysis Technology Overview

The heart of the SPE Water Electrolyzer is the electrolysis cell which consists of an ion

exchange membrane with Teflon® bonded, finely divided metal electrodes. Figure 1 shows this

arrangement along with the water electrolysis reactions. Since the fixed acid ion exchange

membrane has neither a traditional bubble point nor free electrolyte, operating pressures and

hydrogen-oxygen differentials are limited only by the surrounding structures. This affords a

significant safety factor in maintaining positive separation of the hydrogen and oxygen products.

SPE m is a registered trademark of Hamilton Standard Division of United Technologies corporation

Teflon& is a registered trademark of E. I. DuPont De Nemours

193

F KCfiL4NG PAGE BLANK NUT tkMEU
The introduction of the perfluorocarbon cation exchange membrane in the late 1960s

enabled the development of the SPE Electrolyzer. In prior years, water electrolyzers made with

existing ion exchange membranes had useful lifetimes of only a few hundred hours. With the use

of perfluorocarbon ion exchange membranes, the SPE Water Electrolyzer cell life has been

demonstrated to be in excess of 14 years and projected to over 30 years depending on operating

conditions.

In most practical applications, a number of cells are stacked in a filter press arrangement

with as many as 100 cells electrically connected in series while the fluids are passed through the

cells in parallel. Without any free electrolyte, the parallel fluid flows can be conducted without fear

of shunt currents inducing stray water electrolysis and its potentially deleterious result of product

gas mixing. The purity of the product gases from SPE Water Electrolyzers is typically greater than

99.99%.

ION EXCHANGE

CATHODE ANODE

HYDROGEN OXYGEN

ELECTRODE ELECTRODE

4H + + 4e —Do-- 21-1 2 4H+ 21-120 —10- 4H + + 4e + 02

EPROCESS WATER

4e I I

DC SOURCE

Fig. 1. - SPE Water Electrolyzer Reactions

Nuclear Submarine Oxygen Generators

The cell structures, developed during the 1970s for nuclear submarine oxygen generation,

have a capability of operating continuously at up to 400 psi. Limited by stack compression and

cell frame strength, pressures higher than 400 psi are obtained by placing the entire stack within

a pressure vessel. Figure 2 shows a pair of 100 cell stacks, for 3000psi operation, with their

pressure vessel domes removed. The thick lower end plates form one end ofthe pressure vessels.

194
The nuclear submarine SPE Oxygen Generator cell structure is also limited to about 400

psi differential pressure between oxygen and hydrogen. The significance of this design limit is that

a moderately complex pressure control system is employed at 3000 psi. The pressure control

system maintains oxygen and hydrogen pressures within 400 psi even during a power loss

depressurization.

Fig. 2. - 225 SCFH U.S. Navy SPE Oxygen Generator Modules

Space Station SPE Propellant Generator Demonstrator

Under contract to NASANSC, a 3000-psi hydrogen-oxygen generator based on the Naval

0.23 ft2 SPE Water Electrolyzer cell configuration was designed and delivered. The purpose was

to demonstrate the feasibility of producing 3000 psi hydrogen and oxygen on orbit for periodic

rocket motor firing to maintain Space Station Freedom orbital altitude.

To reduce the mass and decrease the volume, as compared to the U.S. Navy design,

several configuration changes were made to the supporting pressure vessel and fluid manifold.

The resultant SPE propellant generator demonstrator is significantly smaller and lighter. In the SPE

propellant generator demonstrator, the pressure vessel is two torispherical domes opposed on

either side of a central fluid plate. This configuration, with a rating of 4 pounds of water electrolyzed

per hour, is shown in Figure 3. The dimensions of the SPE propellant generator demonstrator are

13 inches across the domes and 13 inches in diameter at the dome flanges. The total weight of

the cell stack for space station propulsion is 193 lbs.

195
Mfg

».^

s"a

Fig. 3. - SPE Propellant Generator, Domes Removed

Although the weight of the pressure vessel for the SPE Propellant Generator is significantly

reduced from the nuclear submarine SPE Oxygen Generator pressure vessel, the system

complexity remains the same. Figure 4 describes the SPE Propellant Generator system fluid

schematic. As shown on Figure 4, the pressure vessel is charged with nitrogen gas. This gas Is

increased slowly while the system comes to pressure so as to prohibit any differential pressure

from exceeding 400 psi. Both the product hydrogen and product oxygen are pressure regulated

using the pressure vessel nitrogen as a reference. An elaborate array of regulators, valves and

relief valves is required in each of the three gas systems to prevent any unsafe condition (i.e., no

two simultaneous and independent failures will produce a safety hazard). Since the cells are only

designed for 400 psi differential pressure, no two failures in the pressure control system can allow

pressure differentials to exceed 400 psi.

THE NEW SPE CELL STRUCTURE

In order to further reduce weight and simplify the high pressure SPE Electrolyzer system

pressure controls, a new SPE cell structure has been developed. This new cell structure requires

no pressure vessel enclosure in order to operate at elevated pressures. Demonstrated

characteristics of this new cell structure include:

• Leak free at greater than 3000 psi internal pressure

• Leak free at greater than 3000 psi oxygen chamber to hydrogen chamber differential pressure

• Pressure capability demonstrated with advanced thin membranes

• Over 30 , 000 operating cell hours demonstrated at 2000 psi differential pressure

196
H2

STC

t

VI

)2

RAGE

02

VENT

N2

PPLY

COOLANT COOLANT RETURN

OG1022006ag

Fig. 4. - SPE Propellant Generator System Fluid Schematic

• Over 4000 operating hours on a single unit at 2000 psi differential pressure

• Three cell stack operated at 2000 psi differential pressure.

The basic construction materials for the new SPE cell structure are the same as used in the nuclear

submarine SPE Oxygen Generator system. Therefore, from a materials compatibility viewpoint,

the maturity of over 10 million SPE cell hours in nuclear submarine applications is maintained.

AsPreviously mentioned, the weight of the domed stack in the Space Station SPE

Propellant Generator Demonstrator is 193 lbs. In using the new SPE cell structure the stack weight

for the same total active cell area would be 30 lbs. A minor volume improvement is also obtained

by elimination of the pressure vessel and a reduction of the cell spacing from seven cells per inch

in the nuclear submarine design to 10 cells per inch in the new structure.

The new SPE cell structure produces improved electrolyzer efficiency by significantly

reducing internal component electrical resistance in the non-Omembrane components. Further,

ionic conductivity is improved since the new SPE cell structure is compatible with new high

performance thin membranes whereas the SPE nuclear submarine electrolyzer cell structure is

not. Figure 5 displays the performance improvement obtained with the new SPE cell structure.

197
---^ TYPICAL SPE NUCLEAR SUBMARINE ELECTROLYZER PERFORMANCE

•••••••• NEW CELL STRUCTURE WITH SPE NUCLEAR SUBMARINE ELECTROLYZER MEMBRANE

— — — NEW CELL STRUCTURE WITH ADVANCED THIN MEMBRANE

120° F

150 PSIA

i

i

2.5

2.3

U

D

w

0 2.1

HJ

0

JJ

W

w 1.9 J

0

Z

1.7

1.5

0 200 400 600 800 1000 1200 1400

CURRENT DENSITY — AMPS/FT2 (ASF)

DG1022001ag

Fig. 5. - Single Cell Electrolyzer Performance

SYSTEMS IMPACT FROM THE NEW SPE CELL STRUCTURE

Apart from the aforementioned weight reductions from the new SPE cell structure, the

simplification and reliability improvement in resultant systems is dramatic. In the case of the Space

Station SPE Propellant Generator, the nitrogen subsystem would be altogether eliminated as

would the array of regulators, valves and relief valves intended to maintain differential pressures

below 400 psi. A high pressure SPE Electrolyzer fluid schematic incorporating the characteristics

of the new cell structure is shown in Figure 6. This figure depicts the required SPE Electrolyzer

fluid schematic when it is desired to deliver both gaseous products at high pressure (i.e., high

pressure hydrogen and oxygen for reboost propulsion or lunar base energy storage). This Figure

6 fluid schematic can be compared with the previously displayed Figure 4 to obtain a full

appreciation of the components eliminated by virtue of using the new SPE cell structure.

Further simplification and reliability improvements and weight reductions are obtained if

only one of the product gases is desired to be delivered at high pressure. In the nuclear submarine

SPE Oxygen Generator cell design both product gases are produced at high pressure when only

the oxygen is required to be at 3000 psi. The hydrogen, having been produced at about 3000 psi,

is simply regulated down to its desired pressure level. The new SPE cell structure allows each

product gas to be generated at its own desired pressure level.

198
02

rORAGE

02

VENT

COOLANT COOLANT RETURN

DG1022005ag

Fig. 6 - High Pressure SPE Electrolyzer Fluid Schematic with the New Cell Structure

A lunar soil or material processing facility could require high pressure hydrogen from water

electrolysis while the product oxygen was discharged to ambient for local metabolic purposes.

Such an electrolyzer with the new SPE cell structure is shown on Figure 7. The further

improvements in this configuration include:

• The high pressure water feed pump is eliminated.

• Oxygen pressure regulation components are eliminated.

• Remaining oxygen system components can be designed for ambient pressure operation at

a significantly reduced weight.

• Additionally, the high pressure hydrogen/ambient pressure oxygen configuration is more efficient due to reduced diffusion and the oxygen pressure Nernst effect.

199
H,

H2

STORAGE

02 VENT

H 2 VENT

COOLANT COOLANT RETURN

DG 1022004a g

Fig. 7. - High Pressure Hydrogen SPE Electrolyzer Fluid Schematic with the

New Cell Structure

The last possible configuration with the new SPE cell structure is an electrolyzer producing

high pressure oxygen and ambient pressure or low pressure hydrogen. Such a configuration

could be used to recharge spacesuit oxygen supplies while on orbit. Figure 8 shows the fluid

schematic for high pressure oxygen and low pressure hydrogen. In this arrangement process

water is introduced into the SPE Electrolyzer through the hydrogen chambers of each cell.

Because water must diffuse through the cell membranes to the oxygen anodes before the

molecule is split into oxygen gas, electrons and protons, current densities are limited to a

maximum of 500 amperes per square foot. However, the very major advantage is the elimination

of almost all of the high pressure components.

200
02 STORAGE H2 VENT

02 VENT

COOLANT COOLANT RETURN

DG1022003ag

Fig. 8. - High Pressure Oxygen SPE Electrolyzer Fluid Schematic with the New Cell Structure

SUMMARY

The recently developed new SPE cell structure has allowed a much greater range of

capabilities in the SPE Electrolyzer product than with prior designs. Depending on the application,

it is expected that the simpler system will result in reliability improvements by at least a doubling

of MTBF; system weight reductions up to 50%; an improved efficiency or a doubling of current

density at a constant efficiency; and last but not least, system costs significantly reduced.

201
Section

Three

Advanced Concepts

for Space Power

Nib&CiLXNG PAGE BLANK NOT Fk-MEU
Psly 94-

HIGH ENERGY DENSITY ALUMINUM-OXYGEN CELL

E. J. Rudd and D. W. Gibbons

ELTECH Research Corporation

Fairport Harbor, OH 44077

23363

ABSTRACT

An alternative to a secondary battery as the power source for vehicle propulsion is a fuel cell, an example

of which is the metal-air cell using metals such as aluminum, zinc, or iron. Aluminum is a particularly attractive

candidate, having high energy and power densities, being environmentally acceptable and having a large, established

industrial base for production and distribution.

An aluminum-oxygen system is currently under development for a UUV test vehicle and recent work has

focussed upon low corrosion aluminum alloys, and an electrolyte management system for processing the by-products

of the energy-producing reactions.

This paper summarizes the progress made in both areas. Anode materials capable of providing high

utilization factors over current densities ranging from 5 to 150 mA/cm2 have been identified, such materials being

essential to realize mission life for the UUV. With respect to the electrolyte management system, a filter/precipitator

unit has been successfully operated for over 250 hours in a large scale, half-cell system.

INTRODUCTION

An alternative to a secondary battery as the power source in a vehicle is a fuel cell. It has been shown that

metals such as aluminum, zinc, lithium and iron may be used as a "fuel" in the fuel cell or metal-air battery.

Aluminum is a particularly attractive candidate. The metal has both high energy and power densities, it is

environmentally acceptable (as are the products of the cell reaction), it is easy to handle and has a large, industrial

base for production and distribution.

The chemistry that is involved in the aluminum-air cell is shown by the following equations:

Overall Cell Reaction: Al + KOH + 3/4 0, + 3/2 HO —4 KAI(OH), (1)

Anode Reaction: Al + KOH + 30H - --) KAI(OH), + 3e' (la)

Cathode Reaction: 3/4 0, + 3/2 H2O + 3e —> 30H- (lb)

Corrosion Reaction: Al + KOH + 3H 20 —> 3/2 HZT + KAI(OH), (2)

Crystallization Reaction: KAI(OH), --> KOH + AI(OH) 3 ^ (3)

In operation, aluminum is dissolved (Equation 1) forming a soluble aluminate species and consuming the

alkali metal hydroxide in the electrolyte. A corrosion reaction also occurs in the aqueous electrolyte, again forming

the soluble aluminate together with hydrogen gas (Equation 2). This corrosion reaction is a coulombic inefficiency

and, therefore, must be minimized. As the battery operates, the conductivity of the electrolyte decreases until

precipitation or crystallization of aluminum trihydroxide occurs (Equation 3) , replenishing "free" hydroxide. Thus

a "steady state" condition may be achieved with respect to electrolyte composition and conductivity, at which time

the electrolyte will contain crystals of aluminum trihydroxide. As will be discussed later, controlling the electrolyte

composition is very important.

The reaction at the cathode is the electroreduction of oxygen, which can only be sustained at practical rates

by using a gas diffusion electrode. A three-phase boundary between the catalyst, electrolyte and reactant oxygen

205

PR&Q NG S=AGE VLANA NOT FK- D
must be established and this demands a unique electrode structure. It is known that the presence of carbon dioxide

in the air feed stream to the cathode can lead to loss of performance, due probably to precipitation of alkali metal

carbonates within the structure of the electrode. Therefore, the battery system must include a method to reduce the

levels of carbon dioxide in the air feed steam, or ideally to remove it entirely. For a power source in an underwater

vehicle, the fuel is necessarily oxygen, probably cryogenic oxygen, and removal of carbon dioxide is not required.

Aluminum cannot be electrodeposited from aqueous solutions, so that the aluminum-air battery is not a true

secondary battery. The anode must be replaced mechanically, requiring novel features in the cell design, but offering

relatively rapid refuelability.

The aluminum-air battery is, therefore, a multi-component system (Figure 1) and programs to develop this

technology have necessarily focussed upon:

1. The Electrodes

2. The Electrolyte Management System, i.e., the separation of crystals of aluminum trihydroxide

from the electrolyte axnd control of the composition and conductivity of that electrolyte

3. Cell Design to allow replacement of the anode material

The Auxiliary System, including a heat exchange unit, safe handling and/or disposal of

hydrogen gas and control of carbon dioxide in the air feedstream

This paper will provide a brief review and summarize the recent progress made in the first two areas, the

electrodes and the Electrolyte Management System.

THE ELECTRODES

Development of Aluminum Alloys

Aluminum reacts rapidly and irreversibly with oxygen to form a strongly-adhering oxide film, which largely

determines the electrochemical behavior in aqueous electrolytes. Macdonald [1] has described aluminum as a passive

metal, which displays an active-to-passive transition, a passive region and transpassive dissolution. The currentpotential relationship can be represented schematically as in Figure 2 and a superior battery anode should, therefore,

exhibit the following characteristics:

1. A very negative activation potential, (Ec rr)

2. High discharge currents at potentials only slightly anodic to the activation potential, i.e., a low

value for the slope b,

3. A low corrosion current, which remains low or decreases at the more anodic potentials

required for dissolution.

Modification of the behavior of the oxide layer through the formation of alloys of aluminum has been

extensively explored [2-6]. The incorporation of small concentrations of metals such as magnesium, calcium, zinc,

gallium, indium, thallium, lead, mercury, and zinc, usually in combinations as ternary or quaternary alloys, has been

effective in achieving activation (high rates of dissolution) and inhibition of corrosion.

The presence of certain impurities in the aluminum can markedly affect the electrochemical behavior. For

example, the corrosion rate is particularly sensitive to the concentration of iron in the metal. Using manganese as

an alloying element has been shown to reduce the rate of corrosion of primary aluminum (99.9% purity), which

contains high levels of iron [7]. Typically alloy fabrication has used metal of higher purity, e.g., 99.995% and

99.999%.

206
The transition from fabrication at the laboratory-scale (bookmold ingots of weight approximately 3-51bs.)

to the pilot-scale (150-250 lb. ingots) is not always readily achieved. The electrochemical performance of the large

ingot material can be variable and often is inferior to that of the bookmold material. That the fabrication "practice"

and thermomechanical processing is important is illustrated by the data presented as Table 1 [8]. This shows the

electrochemical performance of an indium-manganese-magnesium alloy, prepared as a 150 lb. ingot by d rect chill i

casting and processed in sections to give different grain structures and microstructures (Figure 3). For example,

using the process IC, i.e., cooling slowly from the solutionizing temperature and only cold-working (cold-rolling)

the alloy, will lead to an unrecrystallized structure with some precipitation of the alloying elements.

In the recent program to develop a power source for an unmanned underwater vehicle [9], emphasis has

been upon mission life or the range achieved by the vehicle. It was then essential to identify an alloy with a low

rate of corrosion in order to maximize the utilization of the metal fuel. An added advantage was that the hydrogen

gas handling requirements within the system are minimized. The polarization characteristics remained important,

as well as realizing a uniform, "smooth" surface as dissolution proceeded.

The electrochemical characteristics of three candidate alloys are presented as Table 2, showing the

dependence of the anode potential and corrosion current density upon the external load. The percent utilization

realized with these alloys is presented as Table 3, together with data for pure aluminum and earlier alloy materials.

The target levels of performance are also shown in the table and are based upon a corrosion current density of 5

mA/cm2 or less over the operating range of current densities.

The alloy materials were fabricated as large ingots by conventional industry practice, although the data

shown as Table 2 was obtained in laboratory-scale cells. Alloy ERC-4 is clearly superior, meeting utilization

requirements while maintaining acceptable polarization behavior. Alloys ERC-2 and ERC-3 closely approached the

utilization targets but the former material showed poor physical characteristics during dissolution/corrosion due to

intergranular attack. This was particularly serious in the full-scale single cell studies.

Development of Air Cathodes

The air cathode used in the recent programs to develop the aluminum-air or aluminum-oxygen battery as

the power source in (a) an electric vehicle [8], and (b) an unmanned, underwater vehicle [9] is a high performance,

two layer structure. The development of air cathode technology, involving ELTECH Research Corporation and Case

Western Reserve University, has been described elsewhere [8,10] and has led to the present electrode, which is a

laminate of

(a) an active laver, containing the catalyst (cobalt tetramethoxyphenylporphryn) absorbed onto a

graphitized carbon black, pyrolyzed and coated with a perfluorinated ionomer,

(b) a gas supply layer, which is a mixture of the graphitized carbon black and teflon, and

(c) two nickel meshes, one on the gas side for mechanical support and the second on the

electrolyte side for current distribution or collection.

This type of electrode has shown excellent extended performance under both constant current operation

(Figure 4) and cyclic operation (Figure 5).

For use in the unmanned, underwater vehicle, it is necessary that die transport of oxygen gas through the

electrode into the electrolyte be prevented, or at least minimized. A porous barrier layer was added, bonded to the

nickel mesh and surface of the active layer. During operation the pores of this layer readily fill with electrolyte,

providing an effective barrier to the transport of oxygen gas into the electrolyte.

The performance of the modified electrode, over the range of current densities of interest, is shown as Table

207
THE ELECTROLYTE MANAGEMENT SYSTEM

Development of the Filler/Crystallizer

The Al/Air battery may operate in two basic modes, batch or steady state, depending on how one desires

to treat the aluminate species generated in reactions 1 and 2. In the batch mode, fresh electrolyte is charged into

the system and operation continues until the level of dissolved aluminate reaches saturation. This situation

corresponds to an end of charge condition; no solids are formed during batch operation. Conversely, a near constant

electrolyte composition can be maintained in the steady state mode by promoting the crystallization of the aluminate

species into an insoluble hydroxide called hydrargillite, Al(OH) 3 (equation 3). During the cell refueling operation,

the electrolyte tank is emptied and the Al(OH) 3 crystals may be sent to a Hall-Heroult cell for processing back into

aluminum.

The performance of a batch operation battery is electrolyte volume limited and off-board facilities to process

the supersaturated solutions of potassium aluminate would be required. Steady state operation produces a battery

with greater capacity and a more constant voltage profile versus time for a given load. By promoting the

crystallization of Al(OH) 3 and generating KOH, the energy capacity of the battery is significantly enhanced for a

given electrolyte tank volume. This fact is illustrated in Figure 6 for a 200 liter electrolyte tank where battery

capacity is plotted versus packing factor (Pd. The packing factor is a convenient way of representing

filter/crystallizer performance. It may be defined as

VE + VC

P, =

VC

where VE is the electrolyte volume and Vc is the volume of crystals. A packing factor of 1 corresponds to a solid

mass of Al(OH) 3 with no entrained electrolyte. A value of 2 in the figure means that the electrolyte tank is full of

a dense, paste-like slurry at the end of mission, but only 100 liters of that mass is actually Al(OH) 3 crystals.

Solomon Zaromb and R. A. Foust, Jr. first described operation of AVAir batteries with electrolyte

regeneration by precipitation of AI(OH)3 in 1962 [11,12]. Since then a variety of crystallizer/separator techniques

have been attempted to control the electrolyte aluminate concentration as indicated in Table 5 [13-17]. With the

exception of the filtration techniques, all of these methods have been rejected for propulsion applications because

of weight, volume, and/or energy considerations.

Previous studies have shown that optimization of Al/O Z energy and power density requires an electrolyte

management system [18,19]. ELTECH's patented filter/crystallizer unit [20] maintains system simplicity, since only

a filter and two additional valves are required. Mission duration is increased by continuously regenerating potassium

hydroxide for subsequent reaction at the anode and storing cell discharge products compactly as a crystalline

hydroxide species, Al(OH)3 . Continuous crystallization maintains a high and nearly constant electrolyte conductivity;

therefore, stack voltage remains steady throughout the discharge cycle. As the stack is discharged, a crystal cake

forms and gradually increases in thickness with a subsequent increase in pressure drop across the filter. When the

pressure drop reaches a predetermined level, the cake is pulsed off the filter by backflushing (flow reversal) to settle

in the bottom of the tank.

Filter/Crystallizer Operation

Test equipment utilized to investigate behavior of the filter/crystallizer is shown in Figure 7. Total

hydroxide concentration is kept constant by water addition and backflush cycles are produced by 2-way solenoid

"backflush valves." Figure 8 displays the results of a filter test for a power load of 40.3 Amps. This experiment

was run in conjunction with an aluminum half cell driven by a power supply in which H Z was evolved at the nickel

counter-electrode. Electrolyte conductivity was continuously monitored by a Foxboro model 872 conductivity meter

with an electrodeless sensor, 871EC.

208
Backflushes occurred at 5 psi or 202 Ali, whichever came first and, in the figure, a hollow circle indicates

a backflush. A software bug prevented the system from backflushing properly for the first 20 hours. Operation

began again at 40 hours after the necessary software modifications were made. Despite the system upsets the filter

performed well for more than 250 hours, maintaining the conductivity between 300 and 400 mS/cm. Filter lifetime

depends primarily on three factors:

1) the anode alloy,

2) the filter material, and

3) the backflush algorithm

ALUMINUM-OXYGEN SINGLE CELL PERFORMANCE

Full-scale aluminum-oxygen cells (with an active electrode area of 690 cm') have been integrated with the

filter/crystallizer unit for a series of studies. The cell featured a fixed anode with backside temperature control and

a modified AE100 cathode to minimize electrolyte leakage and oxygen gas blowthrough. The anode alloy, ERC-4,

was manufactured from high purity aluminum as a 150 lb. ingot using conventional practice and thermomechanical

processing. The experimental apparatus is that used in the filter/crystallizer studies (Figure 7), but incorporates the

Al/OZ single cell. The process streams were all maintained at 50°C with an oxygen gas flowrate of 1.75x

stoichiometry. The unreacted oxygen was recycled to the cell using a peristaltic pump.

This single cell was operated at a constant current density of 50 mA/cm', drawing more than 5,000 amphours, during which time the cell voltage declined from 1.43 volts to 1.3 volts (see Figure 9). This decline in the

cell voltage may be attributed to the increasing interelectrode gap in this particular (fixed electrode) cell design. A

movable anode prevents this ohmic voltage loss and is essential to achieve the required high energy density. The

data shown represents an energy yield of 3.9 kilowatt-hours/kilogram aluminum, which is significantly better than

that obtained with earlier aluminum alloys.

Polarization data was recorded periodically and is presented as Figure 10 at two time and conductivity

intervals. The fact that the data obtained approaches the target performance for a single cell is most promising,

since, as stated above, the cell uses fixed anodes. The large decrease in cell voltage observed from open circuit to

25 mA/cm2 is a consequence of the irreversibility of both the aluminum anode and the oxygen cathode reactions.

At higher current densities the slope of the polarization curve decreases markedly, the voltage presumably being

dominated by the concentration and ohmic overpotentials in the cell.

CONCLUSIONS

A brief review of the aluminum-oxygen cell and related technology has been presented. Two areas of this

technology, considered to be critical to successful application in vehicular propulsion systems, were identified.

Firstly, the fabrication of high perfonmance aluminum alloys was demonstrated, using conventional industry practice.

The alloys provided the required utilization targets for the aluminum fuel in operation in a vehicle. Secondly, an

effective electrolyte management system has been developed, providing the necessary control of composition and

conductivity of the electrolyte as well as removing the solid by-products of the energy-producing reactions. These

developments led to the integration of a single, full-size aluminum-oxygen cell with the filter/precipitator unit to give

over 5,000 ampere-hours of operation, with an overall energy density of 3.9 kilowatt-hours/kilogram aluminum.

ACKNOWLEDGEMENTS

A substantial part of this work was performed under the auspices of the Defense Advanced Research

Projects Agency, Contract Number MDA-972-91-C-0040.

209
REFERENCES

1. "Development and Evaluation of Anode Alloys for Aluminum-Air Batteries," Final Report to Lawrence Livermore

National Laboratory from ELTECH Research Corporation, Contract Number 1806205, Section VIIIB,

November 1987.

2. J. T. Reding and J. J, Newport, Materials Production, 5, 15 (1966).

3. T. Sakano, K. Toda, and M. Hanoda, Materials Production, 5, 45 (1966).

4. "The Metallurgy and Electrochemistry of the Aluminum Anode," a report from the Fontana Corrosion Center for

ELTECH Research Corporation, May 1984.

5. "The Behavior of Aluminum in Aqueous Solutions," a report to the Norwegian Defense Research Establishment

by T. Valand, NRDE-PUBL-81/1003, 1981.

6. 'Aluminum-Air Battery Development," Final Report to Lawrence Livermore National Laboratory from ELTECH

Research Corporation, Contract Number 1806205, Sections VIIIC, VIIID, and X, November 1987.

7. "Effect of Manganese Additions on the Performance of Aluminum Anode Alloys," a progress report to Lawrence

Livermore National Laboratory by D. H. Scott, UCRL-15473, May 1982.

8. "Development of Aluminum-Air Batteries for Application in Electric Vehicles," Final Report to Sandia National

Laboratories from ELTECH Research Corporation, Contract Number SAND 91-7066, December 1990.

9. "An Aluminum-Air fuel Cell Power System for an Undersea Vehicle (UUV)," Defense Advanced Research

Projects Agency, Contract Number MDA-972-91-C-0040, July 1991.

10. "Status of the Aluminum-Air Battery Technology," Proceedings of the Workshop on Structure Effects in

Electrocatalysis and Oxygen Electrochemistry, Volume 92-11, Published by the Electrochemical Society,

Inc., NJ (1992).

11. S. Zaromb, "The Use and Behavior of Aluminum Anodes in Alkaline Primary Batteries," J. Eiectrochem. Soc.,

109, 1125, (1962).

12. S. Zaromb and R. A. Foust, Jr., "Feasibility of Electrolyte Regeneration in Aluminum Batteries," J. Electrochem.

Soc., 109, 1191, (1962).

13. A. Maimoni, "Lamella Settlers: Material Balances and Clarification Rates," Lawrence Livermore National

Laboratory, UCRL-95924, 1987.

14. A. Maimoni, "Crystallization of Aluminum hydroxide in the Aluminum-Air Battery: Literature Review,

Crystallizer Design and Results of Integrated Systems Tests," Lawrence Livermore National Laboratory,

UCRL-53843, March 1988.

15. A. Maimoni, "Aluminum-Air Battery Crystallizer," Lawrence Livermore National Laboratory, UCRL-95923,

January 1987.

16. T. G. Swansiger and C. Misra, "Development and Demonstration of Process and Components for the Control

of Aluminum Air Battery Electrolyte Composition Through the Precipitation of Aluminum Trihydroxide,"

Alcoa Labs, Final Report under Subcontract No. 5724709, available as UCRL-15503, (1982).

17. M. Balasko, R. J. Adler, R. V. Edwards, and S. B. Adler, 'Initial Studies of Particle Trapping in Impeller

Fluidizers," AIChE Symposium Series, 83, 112, (1987).

210
18. D. W. Gibbons, M. J. Niksa, and S. Sinsabaugh, "Aluminum/Oxygen Fuel Cell with Continuous Electrolyte

Management," Unmanned Systems, 9, 38, (1991).

19. A. Maimoni and J. F. Cooper, "Development of the Aluminum-Air Battery for an Electric Vehicle," Energy

Technology Review, Lawrence Livermore National laboratory, p. 9, April 1988.

20, R. J. Coin, C. W. Brown Jr., and J. M. Noscal, "Metal-Hydroxide Crystallizer and Filter," U.S. Patent

4,994,332, (1991).

211
TABLE 1

Electrochemical behavior of an InMgMn alloy - effects of thermomechanical processing

Open Circuit 100 mA/cm2 200 mA/cm2 300 mA/cm2 400 mA/cm2

Potl Corr Pod Corr Potl Corr Potl Corr Potl Corr

From Process IA' -1.72 51.0 -1.70 35.0 -1.71 70.0 -1.66 58.5 -1.60 65.0

From Process 113' -1.77 38.1 -1.74 69.0 -1.56 20.0 -1.41 22.0 -1.27 10.0

From Process 1C' -1.75 27.7 -1.67 31.9 -1.50 20.6 -1.38 17.5 -1.25 14.0

From Process 2' -1.76 30.4 -1.71 41.0 -1.69 47.4 -1.65 58.0 -1.60 60.9

Electrolyte: 5M KOH, 60°C

Potential: Volts vs. Hg/HgO reference electrode

Corrosion: Expressed as a current density, mA/cm2

'as in Figure 3

TABLE 2

Electrochemical performance of selected alloys

at 7 mA/cm2 at 25 mA/cm2 at 75 mA/cm2 at 125mA/cm2 at 175 mA/cm2

Pod Corr Potl Corr Pod Corr Pod Corr Potl Corr

ERC-2 -1.78 4.8 -1.62 2.4 -1.47 2.7 -1.42 2.25 -1.39 1.25

ERC-3 -1.77 3.5 -1.71 1.3 -1.63 0.5 -1.35 2.0 - -

ERC-4 -1.76 3.0 -1.58 0.2 -1.57 0.1 -1.52 0.8 -1.49 0.1

Electrolyte: 4M KOH containing 1.5M dissolved aluminum, 50°C

Electrode Area: 6.45 cm'

Potential: Volts vs. Hg/HgO reference electrode

Corrosion: Expressed as a current density, mA/cm2

212
TABLE 3

Corrosion characteristics of selected aluminum alloys

% Utilization' at

5 mA/cm2

50 mA/cm' 100 mA/cm2 200 mA/cm2

High Purity Aluminum 0.05 31 54 76

ERC-1 12.5 56 71 81

Alcan BDW 19 72 83 92

ERC-2 43 94 97 98

ERC-3 59 98 98 -

ERC-4 62.5 —99 ---99 —100

Performance Target > 50 >91 >95 >97.5

Electrolyte: 4M KOH containing 1.5M dissolved aluminum

Temperature: 50°C

'% Utilization = Faradaic Dissolution Current (load) %

Faradaic Dissolution Current + Corrosion Current

TABLE 4

Electrochemical performance of the modified cathode

Cathode Potential' (volts) vs. Hg/HgO at

Open

Circuit 25 mA/cm' 50 mA/cm' 100 mA/cm2 200 mA/cm2

After 27 hours at 50 mA/cm2 +0.06 -0.07 -0.1 -0.13 -0.19

After 313 hours at 50 mA/cm2 +0.06 -0.06 -0.09 -0.12 -0.18

After 1000 hours at 50 mA/cm2 +0.07 -0.07 -0.09 -0.1 3 -0.21

Electrolyte: 4M KOH at 60°C

Gas: Oxygen, dry

Electrode Area: 14 cm2

'Corrected for IR contributions

TABLE 5

Alternative Solids Separation Techniques

Gravity Settler -Wound Cartridge Filter

-Lamella Settler -Cross Flow Filter

• Hydrocyclone •Helipump Impeller Fluidizer

-Centrifuge Poulos Separator

213
Cos Out

Knockout Tank Hydrogen

Disposal

C

Scrubber

a- Power_ L

— — J Gcs Seporotion

r---------,

L__ --;

Hea t

Exchanger

Filter

AIR or 02 W

Electrolyte

Tank

Hydrorgillite

— - -- _.T .... _--.....-......

Electrolyte

Management

Process Flow Key System

Electrolyte

— — — Gas

------- Coolant

Fig. 1 Schematic Representation of the Aluminum-Oxygen cell system.

Slope

w /

o! /

C /

U

I

I

I Id

I

I corn (Stond-by) I

I

I

0

1

E crit E d

Standby Operating

POTENTIAL.

Fig. 2. -- Schematic representation of the current-potential

relationship for aluminum in alkaline solutions.

214
Cold Worked

Slow cool from solutionizing temperature

1C 18

• Solute out • Solute out

• Unrecrystaflized • Recrystallized

• Solute in • Solute in

• Unrecrystallized • Recrystallized

2 lA

Thermally

treat

after cold

rolling

Rapidly quench from solutionizing temperature

Fig. 3. -- Matrix to show the four types of structures

resulting from the processing variations.

Cathode Potential (mV vs RHE)

2 4 6 8 10 12 14

Time Online (Hrs) (Thousands)

^k 200 mA/cm2 in 7.5M KOH conloining 1.0m aluminale of 601C

+ 450 mA/cm2 in 5.OM KOH conloining 1.OM oluminole of BQC

Fig. 4. -- Performance of polymer-modified air cathodes -

constant current operation.

1000

BOO

600

400

200

0L

0

215
Steady State Operation

Typical

Operation

Window

Batch Operation

t OM KOHL 7M AI @ saturation

— I

Capacity / 1000 (A h)

600

500

400

300

200

100

0

P = V E + V C

V

C

VE= Electrolyte Volume

VC= Crystal Volume

Cothode Potentiol (mV vs RHE)

1000

800

600

400

200

08 1 2 3 4 5 6 7 8

Time Online (Hrs) (Thousonds)

At 450 mA/cm 7 in 5f.f KOK of WC. The cell wos cycled. 3 hours on/3 hours off.

During the off cycle. the cell remoined filled with electrolyte.

Fig. 5. -- Performance of polymer-modified air cathodes - cyclic

operation.

2 3 4

Packing Factor ( P )

BASIS t

Electrolyte Tank = 200 liters

Fig. 6. -- Capacity for a given electrolyte volume.

1 5

216
a10

7

6

5

4

3

2

1

0

—1

..........

11 I-P -0

6^!

^AO

C

---------- i0=11.

C.Mro,

Hyd-q— 0

C—d—, ..t.,

F..

9

— — — — — — —

C—amty I J TV.,-, Electrolyte

P­P

i

Specifications

-j

KOH

Legend ERC I

Coneu<6.7y Ca moue: Yo 

r— Electrolyte Flo. E-11- 681.6 — f" ^E`00

C.—I

V^.q^ ------ -- O.Yq— F'.. I—

filte r 1.. 202 _2

C M-I F—

cap 0.140 Cm

E4. .,Yt. 6.0 61—

Fig. 7.-- A I/0 2 single cell test apparatus.

550

5170

45o E

U

400 U)

350

300

250

200 -0

150 0

U

100

so

0

0 20 40 60 80 100 120 140 160 180 200 220 240 260 280 300

Time (hours)

1 40.32& 3.6 1/n

5 .

^

^.. I.s. , ;nh

1— - 202 —2. C.—

Fig. 8. -- Filter/crystallizer test.

217
r

T T T

5 psi or 5h bockflush criterion

50 mA/cm2, 1.75x stoich 02

02 P = 15" H2O

2.0

1.9

1.8

1.7

1.6

u 1.5

O 1.4

1.3

1.2

1 1

c 1.0

Q 0.9

i 0.8

0.7

y 0.6

U 0.5

0.4

0.3

0.2

0.1

00

IUUU

9000

8000 L

7

7000 0

6000 E

5000

Z

4000 ML

3000,

m

2000 rM

-r1000 U

0

0 20 40 60 80 100 120 140 160 180 200

T - SOC. 34.3EA. 3.6 1 /h,. N/02 i"I, cell

Time (hours)

5.0 1, 4M KOH/1 M Al, ERC-4 Alloy

Finer Areo - 202 cm2, moEifId AE-100

Fig. 9. -- Al/0 2 single cell test.

Cell Potential (Volts)

........................................................................... .......... ..... . ........1......................

.............................................. I............. .. _

K = 385 mS/cm

)c - 316 mS/cm

........................ ...._... ..... .............................. _ _ . .................................. .......

0 25 50 75 100 125 150 175 200

Current Density (rnA/cm2)

- 13665-059, 7 hrs --e-- 13665-059, 25.3 hrs - Talgel, 350 mS/cm

T=50C, 5OmA/cm2, 1.75x rearc 02, 3.611h

4M KOH11.5M Alloy. A - 687.6 cm2

Modified AE-100 cathode, ERC-4 Anode

Fig. 10. -- Full scale Al/0 2 single cell results polarizations

during mean power test.

2

1.8

1.6

1.4

1.2

1

0.8

0.6

0.4

0.2

0

218
N94-23364

COMPOSITE FIBER STRUCTURES FOR CATALYSTS AND ELECTRODES

Christopher J. Marrion, Donald R Cahela, Soonho Ahn

and Bruce J. Tatarchuk

Department of Chemical Engineering

and the Space Power Institute

Auburn University, Alabama 36849

ABSTRACT

We have recently envisioned a process wherein fibers of various metals in the 0.5 to 15

micron diameter range are slurried in concert with cellulose fibers and various other

materials in the form of particulates and/or fibers. The resulting slurry is cast via a wet-lay

process into a sheet and dried to produce a free-standing sheet of "composite paper." When

the "preform" sheet is sintered in hydrogen, the bulk of the cellulose is removed with the

secondary fibers and/or particulates being entrapped by the sinter-locked network provided

by the metal fibers. The resulting material is unique. in that it allows the intimate contacting

and combination of heretofore mutually exclusive materials and properties. Moreover, due to

the ease of paper manufacture and processing, the resulting materials are relatively

inexpensive and can be fabricated into a wide range of three-dimensional structures. Also,

because cellulose is both a binder and a pore-former. structures combining high levels of

active surface area and high void volume (i.e., low pressure drop) can be prepared as freestanding flow through monoliths.

INTRODUCTION

Electrode structures and catalysts which function at a steady-state conversion level

must simultaneously accommodate a number of physical processes including, among others:

surface and interface kinetics and selectivity, heat transport, mass transport, thermal and/or

electrical conduction, corrosion and sintering/deactivation. To beneficially accomplish these

219
physical processes electrode/ catalyst designers must consider a number of individual design

factors including void volume, specifle surface areas of both catalytic/electrocatalytie

materials and supports, micro- and macroporosity, mechanical stability, cost, safety and the

addition of various additives and binders to promote wetabihty, hydrophobicity, reactant

contacting and interfacial mass transport.

Many of the high surface area electrodes in use today are fabricated from particulatebased materials which are retained onto a screen/collector my means of various organic and

polymeric binders (e.g., PTFE). Due to the packing constraints provided by these particulates

and binders, free-standing electrode structures are only available in a limited range of

flexibilities, void volumes and macroporosities. 'lie resulting effectiveness factors and Thiele

moduli often preclude the effec:Uve utilization of high surface area catalysts/electrocatalysts

because of intra-electrode heat and mass transport limitations.

The focus of this effort is to utilize fibrous precursors as building blocks for enhanced

electrode and catalyst structures. As shown in Figure 1. there are a number of inherent

benefits derived from using these materials in electrically conductive and/or high surface

area structures. Furthermore, the high aspect ratio possessed by fibers allows fabrication of

free-standing structures of higher void volume and less resistance to both heat and mass

transport. The process described herein also utilizes the sinter-attachment of micronic metal

fibers to provide necessary mechanical contacting and stability. This feature eliminates the g

priori use of void-filling binders unless they are critically required for other chemical reasons.

EXPERIMENTAL

Details of paper preform fabrication, sintering conditions, optimization considerations,

choice of materials and other electrode preparation techniques may be found in references

11-71.

220
FABRICATION OF MONOLITHIC STRUCTURES

Due to the ease of paper processing and handling, preform sheets may be corrugated

and rolled or layered so as to create flow channels for low transport resistance and high

accessibility of the constituents (see Figure 2). Preform sheets may also be layered prior to

sintering so as to create thicker electrodes/ catalysts which can readily accommodate graded

porosities and graded compositions depending upon the number and type of sheets which are

stacked. Preform sheets incorporating ceramic fibers and cellulose may also be layered into a

stack so as to integrally incorporate an electrolyte separator.

Metallic fibers in the preform permit these materials to be sintered to various current

collectors and metallic surfaces as shown in Figures 3-5. With respect to the structures

shown in Figures 3-5, cyclic voltammetry and other electrochemical techniques generally

demonstrate excellent mechanical and electrical connectivity between mieronic metal fibers,

high surface area activated carbon fibers and the various substrates depicted. The only

exception being the interior of a tube. Upon sintering and cellulose removal, there is a net

shrinkage of the preform causing it to adhere quite tightly to the exterior of a curved surface.

whereas there is a net recession away from the interior surface. This problem has been

addressed by pre-sintering to a metallic mesh followed by attachment to the interior. The

geometries shown in Figures 3-5 appear to meet many of the requirements necessary for

incorporating fibrous materials into real devices while permitting electrical conduction as well

as appropriate heat and mass transport.

INCLUSION OF CERAMICS, REFRACTORIES AND CATALYSTS: CONCOMITANT

IMPACT ON VOID VOLUME

When high surface area and electrically conductive electrodes/catalysts are desired it

is possible to incorporate activated carbon fibers from a variety of sources into the preform

221
along with appropriate metal fibers. Activated carbon fibers provide free-standing structures

with up to 98% void volume. Such void volumes are beneficial for promoting mass transport,

however, volumetric reactivity and energy density density are often sacrificed. This is

especially the case for an electrochemical structure, where the void volume must be fIlled with

the mass of a relatively low conductivity electrolyte. In these instances, the electrode/catalyst

designer often requires a lower void structure and it is beneficial to use particulate based

materials along with metallic fibers. Particulates allow more dense packing and are available

in a broad selection of physical properties at reduced costs compared to fibers.

To address the above noted design consideration, Figures 6-9 demonstrate the

feasibility of retaining catalytic/electrocatalytic particulates at various loading and

compaction levels, approaching ca. 20% void volume. The high aspect ratio of the metallic

fibers permits them to provide a "binding' function at a fraction of the corresponding weight

or volume required for an organic or particulate binder. Metallic fibers also impart superior

electrical and thermal conduction. Moreover, intraelectrode potential drops and ohmic losses

are minimized and there is an opportunity to use electrocatalytic materials and supports

often considered too resistive for practical cell designs. As demonstrated in these Figures,

loadings as high as 6:1, particulate: metal fiber, are readily obtained. Compaction and void

volume are controlled by mechanical pressure before, after or during sintering, as well as the

amount of cellulose added to the preform. The aspect ratio and geometry of the fibers and

particulates also impact upon the void volume.

ADDITION OF PRECIOUS METALS

Figure 10 provides selective chemisorption data and O2 -Hz titration data for a

stainless steel-activated carbon fiber electrocatalyst following addition of platinum.

Impregnation is via an incipient-wetness technique with the liquid exterior to the carbon

fibers removed by wicking into an appropriate blotter. High dispersions are obtained provided

222
the platinum is added after metal fiber sintering. Noncorrosive impregnation solutions must

also be used to avoid attack on the stainless steel fibers.

TRANSPORT RESISTANCE VERSUS CATALYTIC ACIiVTIY

Figure 11 provides polarization curves for oxygen reduction in a KOH half-cell, details

of which are described elsewhere 181. The salient feature of this plot is that a stainless steelcarbon fiber composite electrode (without Pt) can outperform a Pt-containing commercial

electrocatalyst at higher current densities. Since the slopes of the polarization curves shown

in Figure 11 are proportional to the lumped resistance, it appears that the nascent reactivity

and easy accessibility of the activated carbon in the composite fiber electrode permit it to

operate with a reduced mass transport limitation. Further details of this study, and an

impedance analysis of the various resistance terms, are described elsewhere 181.

CONCLUSIONS

Mixed fibrous materials can be prepared from dissimilar materials to achieve physical

propertles which are often considered to be mutually exclusive using other means or

materials. The enabling attribute/requirement is the entrapment of one or more components

within the sinter-bonded matrix of the other component. Unique variations of specific surface

area, void volume, macroporosity, etc., are obtained. Particulate materials may also be

incorporated into these structures as well as active electrocatalysts. The ease of paper-making

and paper processing/reprocessing permits ready fabrication of various monoliths and

integrated structures at relatively low cost.

The nature of this fabrication approach, and the physical attributes of the resulting

materials, permit the electrode/catalyst designer to accommodate heat transfer, mass

223
transfer, ohmic losses, interfacial kinetics and interphase contacting in a manner not

previously available.

ACKNOWLEDGEMENTS

The authors wish to thank the following sponsors for support of various portions of

the work described herein. SDIO/TM via DNA001-85-C-0183, SDIO/TNI via DNA001-90-C0127, U.S. Army - DAAL039260205-2, and the Exxon FoundationREFERENCES

1. D. Kohler, J. Zabasaija, A. Krishnagopalan, and B.J. Tatarchuk, J. Electrochem. Soc..

137 (1), 136 (1990)

2. D. Kohler, J. Zabasga, F. Rose, and B.J. Tatarchuk, J. Electrochem. Soc., 137 (6),

1750 (1990)

3. B.J. Tatarchuk, M.F. Rose, A- Krishnagopalan. J.N. Zabasajja, and D. Kohler, U.S.

Patent 5.080.963

4. B.J. Tatarchuk, U.S. Patent 5,096,663

5. B.J. Tatarchuk, M.F. Rose, A Krishnagopalan, U.S. Patent 5,102,745

6. D.A. Kohler, M.S. Thesis, Auburn University (1989)

7. C.Y. Wu, M.S. Thesis, Auburn University (1991)

8. S. Ahn, Ph.D. Dissertation, Auburn University (1992)

224
Figure 1

Preparation and Design Rationale

Technical Approach: Attributes of the Composite Electrode

Central Theme: New materials from high aspect ratio precursors

Benefits:

• Thin fibers are more accessible to gases and electrolytes,

• Thin fibrous electrodes can be prepared with adjustable porosities and void volumes

to enhance inter- and intra-electrode transport while accommodating required

precipitation of chemical products at the cathode (viz., battery applications),

• Thin fibers can form interlocked networks of dissimilar and normally incompatible

materials (i.e., metals, ceramics, carbons, etc.),

• Thin fibers provide very high surface to volume ratios yet are still self-supporting,

• Thin fibers can provide "averaged" material properties which are normally thought to be

mutually exclusive (i.e., high surface area and high electrical conductivity, thermal and

electrical conductivities, thermal expansion factors, etc.),

• Interlocked networks of thin fibers bonded to metallic backings provide flexible

electrodes structures which can be readily assembled into devices even when one of the

components is relatively brittle,

• Sintered fiber networks do not require pressure, are clog resistant, and possess higher

inherent conductivities than powder systems.

225
Fabrication of Electrode and Catalyst =v ^rc1:^`_ h

Figure'

Figure 3

Composite Fiber Electrodes

and Catalysts in Various Geometries

Al

^.1

Plug in insulating tube Thin layer between

protective mesh

IW

Thick layer between

Plug on insulating tube protecove mesh

Thin layer on mesh Thick layer on mesh

I

protective screen

Plug on conductive tube with

M M

Plug on conductive tube

226
Figure 4

Figure 5

Composite Fiber Electrodes and Catalysts in Furious Geometries

Composite Fiber Electrodes

and Catalysts in Various

Geometries

227
Figure 7 Preform 1:1 Loading

Sintered 6:1 Product

Preform @ 5:1 Loading

swan:,

Fibrous Catalyst/Electrode

Incorporating Carbonaceous Particulate

Fibrous Catalyst/Electrode

Incorporating Carbonaceous Particulate

Particulate (4 150 m2/g) : 2 μm Stainless Steel

228
?reform After Densification/Sintering Preform Prior'Ib Sintering

Preform After Sintering Preform Prior To Sintering Figure 8

J m

A. Extended Densification

B. Intermediate Densification

C. Minor Densification

Stainless Steel Fibrous Network Incorporating

Spray - Dried Ceramic

6 g Ceramic: 1 g Steel Fibers

Stainless Steel Fibrous Network Incorporating

Spray - Dried Ceramic

6 g Ceramic : 1 g Steel Fibers

Figure 9

229
Figure 10

Chemisorption and Titration

Electrodes Impregnated with Pt After Sintering

1000

800 H2-02 Titration

Pts-O-N + 2H2 -> 2 Pts-H + H2O

600

0

400 H2 Cheminorption

Pts + (1/2) H2 -> Pte-H

s

a

200 0 Chemisorpdon

^^ ,,P2,t-z + (1/2) 02 -> Pte-O-Pts

0

0 50 100 150 200 250 300

P„

Impreg Metal H2 Particle

Catalyst Method Precursor Uptake Dispersion Size

13.6° Pt/C Soaking/ (NH3 )4Pt(NO3)2 225 umoles 0.65 1.7 nm

Blotting

Figure 11

02 Reduction in KOH Half - Cells

0

o-&e Composite Electrode with 8 wt59 Pt

mow" Commercial Electrode

—200 (0.45 mg Pt/cm , 10 wtT on C)

•t^+• Composite Electrode (No Pt)

w

U

1n

v^

—400

T = 298 K

v

—600

C

_N

O

—800 Commercial Composite (No Pt)

Weight (m /cm2 ) 113 16

Thickness mils) 16.7 5.5

Cost ($/cm 2 ) 1.5 3.0E-4

—1000

0 50 100 150 200 250

Current Density (mA/c m2)

230
y^9636 - 60 2 N94-2336 5

LIMITING FACTORS FOR CARBON BASED CHEMICAL DOUBLE LAYER

CAPACITORS

M. Frank Rose, C. Johnson, T. Owens and B. Stevens

Space Power Institute

231 Leach Center

Auburn University

Auburn, Al 36849

ABSTRACT

The Chemical Double Layer (CDL) capacitor improves the energy storage density

dramatically when compared with conventional electrolytic capacitors. When compared to

batteries, the CDL Capacitor is much less energy dense, however the power density is orders of

magnitude better. As a result, CDL-battery combinations present an interesting pulse power

system with many potential applications. Due to the nature of the CDL, it is inherently a low

voltage device. The applications of the CDL can be tailored to auxiliary energy and burst mode

storages which require fast charge/discharge cycles. Typical of the applications envisioned are

power system backup, directed energy weapons concepts, electric automobiles, and electric

actuators. In this paper, we will discuss some of the general characteristics of carbon-based CDL

technology, describing the structure, performance parameters, and methods of construction.

Further, analytical and experimental results which define the state of the art are presented and

described in terms of impact on applications.

INTRODUCTION

The major electrical energy storage systems of use in pulsed applications are capacitors,

batteries and inductors. The application of inductive storage, except for short storage times, is

dominated by the need for cryogens and superconducting material. Consequently, the intermediate

storage needs have been almost totally dominated by batteries and capacitors or some combination

of the two. Although the battery has high energy density, it suffers from low power density and a

limited number of charge-discharge cycles. The highest energy density, low voltage capacitors to

date have been the. electrolytic capacitors. While high power density is possible, the units are

polarity sensitive and in the high power versions, must be conditioned to maintain a fully

functioning unit. It is used extensively for voltage smoothing in power systems. The Chemical

Double Layer (CDL) capacitors, using aqueous electrolytes, have power densities much better than

electrolytics and at least an order of magnitude better than the best conventional batteries.

In 1887, Helmholtz discovered that the interface between a conductor and a liquid

electrolyte formed a layer capable of storing charge. For strong electrolytes, this layer [1] is

estimated to be only a few angstroms thick. Due to the fact that there are materials with surface

areas of a thousand square meters per gram, it is possible to engineer practical capacitors with the

surface area, interface thickness ratio on the order of the inverse of the permittivity of free space.

This results in capacitors with a capacitance greater than 1 farad in a volume on the order of a

"dime". This technology was first pursued to practice by SOHIO [2] and NEC[3], and there are a

considerable number of units commercially available for low voltage, high impedance applications.

These units usually employ a rare-earth oxide or some form of carbon as the large surface area

materials. There are a number of electrolytes, both aqueous and organic, available for use in these

capacitors. The ultimate choice is governed by the application. In general, use of the organic

electrolytes increases energy storage density at the expense of internal resistance.

231
There are numerous potential applications for the CDL technology. To date, they are used

for memory backup, and as short-term backup power supplies. They are currently being

investigated for such applications as electric vehicle acceleration and electromagnetic actuation.

The restrictive number of applications is due in part to the need for pressure to reduce the

equivalent series resistance to low values. While easy to achieve in a laboratory prototype, it is

difficult to produce a commercial product with a pressurizing system. Further, due to the

differences in pore geometry, the capacitance is very much a function of frequency which tends to

make application at DC or less than 60 hz the primary region of operation. For high energy/power

density applications, the ultimate device would possess the energy storage characteristics of a

battery and the power characteristics of a capacitor. The CDL technology is a step in the right

direction. During the past 5 years, there have been significant technological advances which tend

to extend the range of applications by engineering out the need for pressure to obtain low internal

resistance and by optimizing the pore geometry for a particular application. In that sense, they are

"designer" devices which can be custom tailored to a particular task. The limits of CDL technology

using carbon as the large surface area material have been explored extensively at the Space Power

Institute. The remainder of this paper will present some of our results.

CAPACITOR STRUCTURE

Double Layer Capacitor Geometry

Figure la is a schematic drawing of a practical embodiment of the CDL technology utilizing

carbon as the large surface area material. The carbon beds and membrane are soaked in a suitable

electrolyte such as dilute solutions of sulfuric acid or potassium hydroxide. For powder beds, a

seal is necessary to hold the powder in place, confine the electrolyte and serve as a confining

medium when the finished product is placed under pressure to increase the contact area between

individual particles in the compact. The end plates are chosen for compatibility with the electrolyte

and as thin as possible. The porous membranes are suitable plastics which can be obtained in a

variety of thicknesses and porosities. The rough equivalent electrically is that of two capacitors

connected in series with the interconnect being the distributed electrolyte which ensures contact

with the double layer. On the "other side of the capacitor", electrical contact is made through the

individual carbon grains to the metallic end plates. It is obvious that the resistance of this path is

greatly determined by the surface contact area between the individual grains and the position of any

given grain within the bed. As a result, it is necessary to add considerable pressure to this

geometry to reduce the internal resistance to a value which is acceptable. For potassium hydroxide

electrolyte, 0.0025 cm nickel is chosen for the end plates and 0.0025 cm Nafion separators are

used in laboratory cells. The thickness of the carbon bed is varied up to several millimeters as

needed.

Figure lb illustrates a capacitor structure which greatly reduces the need for pressure to

produce a low equivalent series resistance. The active area in the capacitor is now made up of

sinter locked fine metal fibers with diameters of a few microns interspersed throughout the bed of

large surface area material. This unique composite can be made with high uniformity using

standard paper making technology and has been described in detail elsewhere [4]. To further

enhance conductivity, the metal fibers are also sinter locked to the metallic end plate which is of the

same material as the fibers. The metal fibers drastically reduce the resistance associated with the

carbon bed by providing larger metallic surface area for the individual grains or fibers of carbon to

contact electrically. Capacitors made in this fashion show little or no pressure dependence of the

equivalent series resistance.

232
Double Layer

It is very difficult to directly probe the interface between the metallic conductor and an

electrolyte interface. There are a number of theoretical models [5] [6] whose only means of

verification is the ability to predict macroscopic measurements such as capacitance and the charge

storage associated with the double layer. The key parameters of interest for the double layer as an

energy storage media are the value of the relative dielectric constant associated with the region in

which the charge is stored, the thickness of the layer and the voltage at which the layer fails. These

parameters are functions of the electrolyte species, and its concentration. Rose [1] has used a

modified version of the Guoy-Chapman theory to calculate the thickness of the double layer as a

function of concentration for strong aqueous electrolytes such as sulfuric acid and potassium

hydroxide. This model shows that the storage layer decreases with electrolyte concentration to

molecular dimensions which is consistent with the assumption that the energy storage media is

essentially a single layer of molecules of the solvent. For aqueous electrolytes, this dimension is

that of the water molecule and is on the order of angstroms. Fundamental studies by Bockris [6]

indicate that the relative dielectric constant of this layer is on the order of 10. From an engineering

perspective, the practical limit to voltage standoff occurs when the potential needed to dissociate the

water is exceeded. That value is about 1.2 volts per cell. Substitution of these estimates of the

critical parameters into the expression for capacitance indicates that it should be possible to build

capacitors with capacitance on the order of 20-50 farads per real square centimeter of surface area.

This is in general consistent with experimental results.

Large Surface Area Materials

Materials such as the activated carbons are highly fractionated with enormous surface area

as a result of the structure pores. The large surface area materials most often used in the double

layer capacitor technologies are the rare earth oxides and carbons such as "carbon black". In

general, the carbons have the highest surface areas. It is possible to purchase commercially

available materials which have surface areas greater than 2000 m2/gm. Surface area of this

magnitude is only achievable with a highly porous structure with average pore dimensions of a few

angstroms in diameter. Activated carbon purchased from the American Norit Company, Inc.

specifies that the surface area is about 1500 m 2/gm with a mean pore dimension of 22 angstroms.

Clearly, pores with dimensions on the order of twice the double layer thickness or less cannot be

used to store charge. As a result, much of the surface area is not accessible in any finished device.

Further, it is necessary for the electrolyte to penetrate the pore in order that a conducting path be

available to the external circuit. Consequently, the resistance from each micro-pore to the external

circuit is dependent upon its size and position within the carbon bed. Clearly a trade off exists

between the surface area and the pore dimensions for a particular application. Figure 2 is an

approximate circuit diagram for a CDL. Starting at the top of the figure, the external resistance

Rext is associated with the external circuit, load, diagnostics, etc, and is under the control of the

experimenter. Since the total capacitance of the array is the sum of the capacitance of billions of

tiny pores, of differing values, in parallel, it is instructive to examine how the array is structured at

least qualitatively. Since there is mirror symmetry in the unit, it is sufficient to examine the

electrical pathways only to the membrane. For a given pore, there is the resistance associated with

the current carrying path through the metallic end plate/metallic fibers and through the fiber/grains

of carbon. In a porous bed with many fibers/grains tightly packed, there are a number of current

carrying paths of various resistances possible from any given pore. The value of the resistance

associated with that path is a function of the packing parameters, pressure, resistivity of the bed

material, resistivity of the metallic end plates,and position within the bed. The Helmholtz double

layer discussed above represents the "dielectric" or energy storage media in the usual sense.

Passing through the double layer, the conducting electrolyte in the pore constitutes the other

conducting electrode. The specific resistance of a pore will be a function of the diameter and length

of the pore and has been designated as Rp. In similar fashion, the electrolyte is distributed

233
throughout the interstices of the bed forming multiple pathways to the porous membrane. The

resistance Rmemb is a function of the porosity of the membrane and is readily determined by

experiment to be essentially the percent porosity times the resistivity of the electrolyte. Note that

each pore is connected to its neighbor by shunt resistances which are functions of the resistivities

of the constituents. The net result is that charge can be redistributed during and after a discharge

event which has not totally depleted the stored energy. While a nuisance, this is not particularly of

concern for most of the applications being considered today.

Electrolytes

Numerous electrolytes have been used for chemical double layer capacitor construction.

The most common of the aqueous electrolytes are sulfuric acid and potassium hydroxide. The

most desirable characteristic of these electrolytes are their low resistivity. Since they are aqueous

based, it is necessary to operate each cell of the capacitor at a voltage less than 1.2 volts. Further,

these electrolytes are corrosive, which places restrictions on the use of materials for long life

devices. By contrast, organic electrolytes such as propylene carbonate with lithium and

ammonium salts have electrical resistivities at least an order of magnitude greater than that of

aqueous electrolytes. Organic electrolytes have been operated at voltages as high as 3.5 volts in

practical devices. Since the energy storage density is proportional to the square of the charge

voltage, it would appear that organic electrolytes would be desirable. Many commercial devices

now employ organic electrolytes. However, the internal resistance of the finished capacitor is

usually 10's of ohms. This high value of internal resistance limits application.

OPERATING CELL CHARACTERISTICS

In our laboratory, we have concentrated exclusively on the use of aqueous electrolytes and

carbon materials for the large surface area . Our reasoning for these choices are the ready

availability of cheap, well characterized materials which need no exotic processing to manufacture.

The general model for a capacitor is that of a pure capacitance shunted by the Equivalent Parallel

Resistance (EPR) and in series with the Equivalent Series Resistance(ESR). Figure 3 shows the

pressure dependence of the ESR for a series of cells areas from 2 - 5 cm2 and with various

thicknesses. In general, the ESR decreases rapidly with pressure for all thicknesses investigated.

There is almost an order of magnitude variation of the ESR for the capacitors made from packed

powder beds. For many applications, it is desirable to reduce the ESR to an absolute minimum.

While it is easy to apply pressure to small devices, it is far more difficult when large storage

capacitors are desired. As was discussed earlier, the physical mechanism responsible for the

decrease in ESR with pressure is the increase in contact area between individual grains in the bed

and a reduction of the path through the electrolyte as the bed is compacted. To first order, the

effect is independent of powder bed thickness and shows the same relative decrease with pressure.

The horizontal line in Figure 3 is the data taken for a composite electrode made from 2 millimeter

large surface area carbon fibers roughly 10 microns in diameter mixed with 2 mm nickel fibers

with a diameter of 2 microns. The mix ratio was 2:1 by mass. Cellulose was added and the mix

converted to a paper sheet using standard paper making equipment. Selective annealing in a

hydrogen atmosphere reduced the cellulose and produced a sinter bond to the nickel backing plate

and between fibers. Figure 4 is a electron micrograph of a typical sintered structure. The larger

fibers in the picture are the carbon fibers. Note that the smaller metallic fibers are effectively

bonded and contact the large fibers at many points along their length. In this way, extremely good

contact between the metal end plates, the carbon fibers and the metal fibers are maintained with

little or no pressure dependence of the ESR.

Since all of the elements of the CDL have resistivities which are functions of temperature,

packed powder bed capacitors were made and operated from about 6 oC to 24 oC. The results are

234
shown in Figure 5. The ESR, over the range measured, changed by about 30% and is almost

totally due to the change in resistivity of the electrolyte and the thermal contraction of the grains

which reduces the contact area. The change in the electrolyte is the dominant term and accounts for

roughly 80 % of the observed changes. Higher temperature behavior has not been investigated

within the program. These data do point clearly to the fact that application of these devices in a

situation with large temperature extremes will require environmental control. It is obvious that the

lower limit to operating temperature is the point at which the electrolyte freezes. Practical

consideration of the ESR requirements will certainly limit the temperature variation for the

electrolyte to a few degrees about some design point. Since the ratio of the load impedance to the

ESR determines the available usable energy, it would appear advantageous to operate the CDL at

higher temperature for low impedance loads. Chemical double layer capacitors are "designer

devices" which can be custom tailored to a particular application and temperature of operation is a

variable which can be exploited.

Of equal interest is the Equivalent Parallel Resistance (EPR). In general, single cell

capacitances of farads are readily attainable. An EPR of a kilohm allows storage time constants on

the order of thousands of seconds. In general that is sufficient for many applications. However,

applications demanding long-term storage may require trickle charging to maintain the requisite

energy. Clever design of a device might result in the losses through internal currents sufficient to

maintain the temperature at the most desirable operating point if operating in a cold environment.

The nature of the EPR is not clearly understood. It appears to be influenced by the "purity" of the

material in so far as that dependence might be mirrored by manufacturing processes. Adsorbed

material on the surface of the carbon and impurities in the electrolyte could provide sufficient

conduction through the double layer to account for the experimental observations. Figure 6 is a bar

chart which shows the EPR for devices made from several carbon powders. All units were

comparable in size and operated at room temperature. While the carbons had differing surface

areas and pore size distributions, they were made by differing processes. While indicative of

residual from processing, we were not able to verify that conclusion. If the EPR were a

conduction process as described above, there should also be a strong correlation with temperature.

Figure 7 illustrates the same temperature dependence as was observed in the EPR, clearly pointing

to conduction through the electrolyte as the mechanism determining the EPR.

Over the range of temperatures investigated and for a constant pressure of 20 kg/cm2, there

was no temperature dependence of the capacitance. Since the beds investigated were tightly packed

powder, after a minimum pressure, there was sufficient contact that almost all of the material was

accessed for capacitance. Figure 8 illustrates the pressure dependence of the capacitance for two

cells of the same area but of differing thicknesses. These cells were not made from the same

carbon and the beds may not have been packed to the same density. Therefore, the capacitance

does not appear to scale linearly with thickness as would be expected. The purpose of this graph,

is to show that the pressure dependence of the capacitance is weak and not a factor in the design of

powder bed CDL capacitors. Similar results are observed for the composite electrodes. Figure 9

is a plot of the capacitance for a number of cells of differing thickness, made from the same

material, while trying to maintain the construction parameters identical. Packed powder bed are

difficult to make uniformly. Intuitively, the capacitance should scale linearly with thickness. To

within the accuracy of our measurements, linear scaling is observed but the slope is less than

expected. The difficulty may be in the ability to accurately produce thick beds with equivalent

packing density.

DISCUSSION

There are two options for practical application. A single device can be constructed uniquely

in one unit for a particular application. In that case, the desired voltage can be met by a bipolar

235
stack of the requisite number of unit cells Cu. In order to achieve the desired value of the internal

resistance, ESR and EPR, and total capacitance, the area and thickness of the cell can be adjusted

to arrive at suitable parameters. Depending upon the application, design by this procedure may

result in more energy than is optimum for the load. Referring to Figure 3, it ispossible to

construct single cells which allow "short circuit" currents as high as 5-10 A/cm2 For a given area,

the ESR and the Cu scale linearly with thickness of the active bed. This procedure results in the

highest energy/power density by optimizing the packaging. More generally, a "standard" module

made up of a number of unit cells can be built and any energy storage unit made up of the requisite

number of standard modules in series/parallel arrays. Again, depending upon the load, it may not

be possible to optimize all parameters of the storage unit simultaneously.

There is a wide variety of materials and techniques available for utilization in CDL

technology. The current laboratory state-of-the-art is about 6 j/gm for aqueous electrolytes with the

capacitive elements made from composite electrodes. For non-aqueous electrolytes, the energy

density is at least double that amount. For internal resistance, the best ESR values for aqueous

electrolytes is on the order of 0.1 ohm/cm 2 of active cell. The thickness of the cell for these values

is 0.002-0.003 cm. Simple bipolar stacks from the carbon powder bed technology have been

constructed to evaluate sealing and any engineering problems which might occur as large numbers

are stacked for higher voltage operation. Up to 50 cells have been made and little or no problems

encountered. The largest of these experimental units operated at 50 volts and stored 1200 joules.

The energy density was greater than 3j/g even though no care was taken to minimize the materials

used in the construction. The unit was tested periodically for several months with little or no

degradation.

Further advancements in this technology must rely on engineering of the pore structure to

maximize the surface area available while simultaneously maximizing the pore diameter. It is

probable that factors of 2 or more improvement can be realized by a concerted effort.

ACKNOWLEDGEMENT

This work was supported by the Strategic Defense Initiative Organization's Office of

Innovative Science and Technology (SDIO/TN) through contract number DNA001 -91-C-0 127

with the Defense Nuclear Agency.

REFERENCES

1. M. F. Rose, Proc. 33rd International Power Sources Symposium, The Electrochemical Society

Inc., June 13 - 16, 1988, Cherry Hill, NJ.

2. D. L. Boos, H. A. Adams, T. H. Hacha, and J. E. Metcalfe, Proc 21st Electronics

Components Conference, 1971

3. K. Sanada and M. Hosokawa, NEC Research and Development, 55, p. 21 - 28, 1979

4. D. Kohler, J. Zabasajja, A. Krishnagopalan, and B. Tatarchuk, J. Electrochem. Soc., Vol.

137, No. 1, 1990

5. I. L. Cooper and J. A. Harrison, Electrochemical Acta, 29, No. 8, 1984

6. J. O'M. Bockris, Materials Science and Engineering, 53, 1982

236
Carbon Bed Seal

Mem

Drbon Bed

er or powder

Seal

Mem

Micron Size Metal Fiber Metal End Plate

Metal tna mate

Metal End Plate

Metal End Plate

Figure la. Schematic drawing of a packed carbon powder bed Chemical Double Layer (CDL)

capacitor

Figure lb. Schematic drawing of a Chemical Double Layer (CDL) capacitor using a

composite electrode structure.

237
End Plate and

Carbon Particles R°7^

FR Rp3 R 2 P1

'pn %3 CP2 CP1

Electrolyte Rmemb

PRpn

Rp ^Rp JRP,

Cps C^3 I-Ip2 CP1

End Plate and

Carbon Particles ^Rext

Figure 2. Approximate schematic for a Chemical Double Layer (CDL) capacitor.

0.6

0.5

U)

0.4

E

00.3

tr

W

0.2

0.1

0.0

2 mm thick

CB-m-p'eTe'k

0.2 mm th c

0 20 40 60 80

Pressure kg/square cm

Figure 3. Pressure dependence of the Equivalent Series Resistance (ESR) for powder and

composite electrode structures.

238
Figure 4. Electron micrograph of composite electrode structure showing large carbon fibers and

smaller sinter-locked metal fibers.

0.00

N

0.00E

E

t

O

W 0.00

0.00E

U 1u z 3u

Temperature C

Figure 5. Temperature dependence for the Equivalent Series Resistance(ESR) for a packed carbon

bed capacitor using H2SO4 electrolyte.

239
30

s

4

E

0

Y 3

B:

0..

W

2

1

0 10 20

Temperature C

8

6

W

E

0

Y 4

c

0_

W

2

0

G-60 DNV S-51 E193 T A88 RB2

Carbon Type

Figure 6. Equivalent Parallel Resistance (EPR) for cells made with packed carbon beds using

carbon powder from several commercial sources.

Figure 7. Equivalent Parallel Resistance (EPR) as a function of temperature for a packed carbon

bed device.

240
c^L

ctt

LL. 3

aDU

C

cd v

^U

Ctf CL 2

Cu

2 mm thick

U

0.2 mm thick

4

1 -}

0 20 40 60 80

Pressure Kg/square cm

Figure 8. Pressure dependance of the capacitance of a packed carbon bed CDL for two bed

thicknesses.

8

(0 7

Cz

m

u- 6

m

U

C

U 5

Cz

C2

c^

U 4

3 -f

0 1 2 3

Carbon Bed Thickness mm

Figure 9. Unit cell capacitance as a function of the thickness of the packed powder bed.

241
y 

39 b31

x U

A HIGH PERFORMANCE H 2-C1 2 FUEL CELL FOR SPACE POWER APPLICATIONS

Everett B. Anderson, * E. Jennings Taylor, " Gerald Wilemski, and Alan Gelb

PSI Technology Company

A Division of Physical Sciences Inc.

20 New England Business Center

Andover, MA 01810

INTRODUCTION

NASA has numerous airborne/spaceborne applications for which high power and energy density

power sources are needed. The proton exchange membrane fuel cell (PEMFC) is an attractive candidate for

such a power source. PEMFC's offer many advantages for airborne/spaceborne applications. They have

high power and energy densities, convert fuel to electrical power with high efficiency at both part and full

load, and can rapidly startup and shutdown. In addition, PEMFC's are lightweight and operate silently.

A significant impediment to the attainment of very high power and energy densities by PEMFC's is

thew current exclusive reliance on oxygen as the oxidant. Conventional PEMFC's oxidize hydrogen at the

anode and reduce oxygen at the cathode. The electrode kinetics of oxygen reduction are known to be highly

irreversible, incurring large overpotential losses. In addition, the modest open circuit potential of 1.2V for

the H2-02 fuel cell is unattainable due to mixed potential effects at the oxygen electrode. Because of the

high overpotential losses, cells using H2 and 02 are capable of achieving high current densities only at very

low cell voltages, greatly curtailing their power output. Based on experimental work at PSI Technology

Company (PSIT) on chlorine reduction in a gas diffusion electrode, we believe significant increases in both

the energy and power densities of PEMFC systems can be achieved by employing chlorine as an alternative

oxidant.

BACKGROUND

A hydrogen/chlorine fuel cell reacts hydrogen and chlorine to form hydrogen chloride. A schematic

of a proposed system is shown in Fig. 1. Hydrogen gas is oxidized at the anode to produce protons which

are transported across the cation exchange membrane. Chlorine gas is reduced at the cathode to produce

chloride ions. Depending on the application duty cycle and lifetime requirements either a captive or flowing

electrolyte may be used.

The net cell reaction is 1/2 H2 + 1/2 C12 a HCI. The standard cell potential for this reaction is

1.36V at 25°C. In an actual fuel cell the operating voltage is lower due to polarization and ohmic losses.

Consequently, the power output is lower than the theoretical power output. Nonetheless, both the oxidation

of hydrogen and the reduction of chlorine are known to be relatively non-polarizable electrochemical

reactions and capable of high cell voltages at high current densities.

We can support this statement for chlorine reduction with our own measurements using the rotating

disc electrode (RDE) technique [1] to determine mass transport parameters and mass transport corrected

kinetic and mechanistic information about chlorine reduction. We have used the RDE technique to obtain

the exchange current density for chlorine reduction on platinum in 7M HCl at 25°C [2] from a mass

transport corrected Tafel plot. We calculated an exchange current density of about 10 -1 mA/cm2. This value

is approximately eight orders of magnitude greater than the exchange current density for oxygen reduction on

platinum (10-9 mA/cm2).

Author to whom correspondence should be addressed.

..Current Address: Farachem Technology, Inc., 300 East Auburn Ave., Springfield, OH 45505.

243

PiiEChDtNQ PAGE BLANK NOT RLMLD

L— —
Relatively little research has been reported on hydrogen chlorine fuel cells. Recent research has

focused on hydrogen/chlorine regenerative fuel cells for load leveling applications [3-7]. The regenerative

fuel cell application, although related, is quite different from the H Z C12 fuel cell discussed here. In the

regenerative fuel cell the chlorine electrode acts in both the consuming (cathodic) and generating (anodic)

modes, while the system referred to in this paper needs to be optimized for the chlorine consuming mode

only.

In the H2-Cl2 regenerative fuel cell, the hydrogen catalyst is attached to a Nafion' cation exchange

membrane to prevent chloride ion from poisoning the platinum electrocatalyst. Halide ions, in general, and

chloride, in particular, have been shown to alter the adsorption isotherms of hydrogen on platinum [8, 91.

The chlorine reactant is dissolved in the electrolyte and is oxidized at a flow-by electrode during the

electrolysis (charge) mode and is reduced at a flow-by electrode during the fuel cell (discharge) mode.

However, mass transport limitations at the flow-by electrode decreases the cathode performance and limit

the power output during discharge.

For a 112-02 fuel cell for NASA, we suggest the use of gas diffusion electrodes analogous to those

used in H2-02 proton exchange membrane fuel cells for power generation. Besides the PSIT work reported

in this paper, there is only one report of the use of a chlorine consuming gas diffusion electrode [10], for the

chlorination of various organics. However, due to limitations of the organic anode reaction, performance at

the chlorine gas diffusion electrode could not be abstracted. Even so, the mass transport limitations at a

chlorine gas diffusion electrode should be considerably less than at flow-by or flow-through electrodes [111.

RESEARCH AT PSIT

PSIT is developing H2-C12 fuel cells for application to chlor-alkali technology. These fuel cells are

intended for widespread commercial use, hence, capital cost considerations are paramount. Furthermore,

high current densities are neither anticipated nor required. Consequently, our current efforts are concerned

with developing low-cost, low-loading electrocatalysts.

Under a DOE Phase II SBIR Program (Contract No. DE-ACO2-86ER80366), PSIT developed a

hydrogen/chlorine fuel cell technology to reduce the power requirements and increase the system efficiency

of brine electrolysis in the chlor-alkali industry. In this program, PSIT determined 1) the kinetics of chlorine

reduction on several catalysts; 2) the actual polarization behavior of a chlorine gas diffusion electrode; 3) the

conductivity of a commercially available membrane in the environment of interest to our system; and 4) the

conductivity of HCI as a function of concentration. Using these data we projected very favorable polarization

behavior for our 142-C12 fuel cell. In addition, we engaged in chlorine gas diffusion electrode development

and constructed a laboratory scale process development unit (PDU) for the H 2-Cl2 fuel cell.

The subject of this paper differs from that of our current developmental efforts in the need to

develop electrodes and fuel cells capable of very high current densities at moderate voltages. The NASA

application requires much higher performance electrodes. PSIT is highly familiar with the basic design issues

and believes in the feasibility of employing high-loading, C1 2 and chloride-tolerant H2 electrodes capable of

high current densities with low overpotential losses. Next, we will summarize our efforts in the areas of

chlorine gas diffusion electrode development and chloride-tolerant electrodes for H2 oxidation.

We have measured current-potential curves for C1 2 reduction at 23°C in 3M HCI on a number of

different electrocatalysts in a gas diffusion electrode configuration. These electrocatalysts included: 1) high

surface area (HSA) carbon black (Ketjen EC300J, Akzo Chemie); 2) beat treated HSA carbon black

(Graphitized Ketjen); 3) low-loading (0.5 mgPt/cm2) carbon-supported Pt; and 4) high surface area lead

ruthenate mixed metal oxide. The polarization behavior of these materials is shown in Fig. 2. The best C12

reduction performance was observed for the supported Ylatinum electrocatalyst. These results indicate

extremely low polarization behavior (70 mV at 1 A/cm) for the supported platinum electrode compared to

that of 02, which does not even register a current under similar conditions. The measured rest potential of

1.29V shows a mixed potential loss of only 70 mV from the equilibrium value of 1.36V, a loss that is also

244
much smaller than that suffered by oxygen electrodes. These results are obtained on what is almost certainly

a far from optimal electrode structure and are highly encouraging. They form the basis for the conservative

power and energy density estimates discussed below.

Chloride ions poison platinum, greatly reducing its effectiveness as an electrocatalyst for hydrogen

oxidation. If unchecked, chloride ion poisoning could preclude the attainment of high power and energy

densities in the H2-C12 fuel cell. In work at PSIT, we have demonstrated that smooth platinum's tolerance of

chloride ion was improved by coating the platinum with a thin film of Nafion. Recently, we developed a

method to impregnate a carbon-supported Pt gas diffusion electrode structure with Nafion and coat the

individual Pt particles in order to utilize the Nafion's ability to protect platinum from the chloride ion. The

method involves soaking the electrode in a Nafion solution for different times to vary the loading of Nafion

in the electrode. In Fig. 3. we present the results of electrochemical Pt surface area measurements on

electrodes treated with Nafion for 30s and 12 hr. These are compared with results for untreated electrodes

that were presoaked in a water/alcohol solution for 30s and 12 hr, respectively, to facilitate wetting upon

immersion in the chloride test solutions.

No effect of chloride was seen with any electrode until the 0.1M chloride concentration. At this

value the standard anode presoaked for 12 hr suffered a 47 reduction in surface area, from 112 to

59 m^

%

/g. The comparable Nafion coated anode demonstrated no loss of surface area at the same

concentration. The coating of the anode with Nafion, therefore, had a positive effect on the chloride

tolerance. The other electrodes also showed no effect at this concentration. The surface area of the lightly

coated anode approached that of the heavily coated anode, suggesting improved chloride tolerance for the

lightly coated anode as well. Any effect of chloride at 0.1M on the untreated anode presoaked for 30s was

probably lost due to incomplete wetting of the electrode.

At 1M chloride ion concentration, the surface area of both untreated anodes decreased, while the

Nafion coated anodes exhibited little change. This behavior in 1M HCl clearly demonstrated successful

protection of the platinum by coating with Nafion. The total surface area lost for the coated anode was less

than 15% compared to a loss of surface area for the uncoated anode of greater than 75%.

H2 C1 2 PEMFC SYSTEM PERFORMANCE

To examine the relative merits of the H2-Cl2 PEMFC system compared to the H2-02 PEMFC

system, we estimated the power and energy densities of these two systems over a wide range of mission

lifetimes. The results shown in Fig. 4 indicate that the H2-Cl2 PEMFC will significantly outperform the

H2-02 PEMFC. For mission lifetimes of 1 min, 1 hr, and 10 hr, we conservatively estimate the energy and

power densities of the 112-02 PEMFC system to be, respectively, 20 Wh/kg and 1200 W/kg at 0.73V,

152 Wh/kg and 152 W/kg at 1.12V and 199 Wh/kg and 20 W/kg at 1.26V. These values are, respectively, 2,

1.5, and 1.33 times greater than the corresponding values optimistically estimated for an H 2-02 PEMFC

system based on state-of-the-art performance recently reported by Ballard Power Systems (BPS) [12]. For

lifetimes greater than 10 hr, the energy and power densities of the H 2-02 system asymptotically approach

1.3 times those of the H2-02 system. It is significant that, according to Appleby [13], the energy density

figure of 152 Wh/kg (and greater) for the 1 hr mission (or longer) "exceeds the energy available from any

known secondary battery, including sodium-sulfur." Also note that the energy density of the H 2-C12 system

exceeds 110 Wh/kg (50 Wh/lb) for all mission lifetimes longer than 15 min.

The estimates shown in Fig. 4 are based on the following assumptions: The fuel cell stack contains

50 cells of 232 cm2 (0.25 ft2) each and weighs 45 kg as suggested by BPS [12]. The assumed mass significantly affects the estimates only for mission lifetimes less than 1 hr; for longer missions the mass of reactants

and storage tanks predominates. The latter masses were taken from an Air Products catalog and are 0.112

kg/mol for C12, 0.190 kg/mot for 02, and 0.206 kg/mol for H2. With these assumptions, the mass to power

ratio M (kg/kW) for each system can be written as

245
M(H2 -02) = 45/P + 5.93t/V (1)

M(H2 -02) = 45/P + 5.62t/V (2)

where t is the mission lifetime in hours, V is the cell potential in volts, and, P is the stack power in kilowatts.

The current (I) - potential (V) behavior of a single cell is adequately represented by a linear

relation: V = V°(1-0.51/Ip), where V° is the rest potential and Ip is the current at which the peak cell power

is obtained. For the BPS H2-02 cell at 102°C and 5.8 atm, V° is 0.985VandIp

 is 1250A. With these

measured values, the linear current-potential relation actually overpredicts cell performance and thus

provides optimistic power estimates for the H 2-02 system. For the H2-C'2 cell, we have V°= 1.34V (our

measured room temperature value of 1.29V plus 50 mV Nernst gain for operating at 6 atm) and a projected

value for 1p of 1943A based on 80 mV loss at 232A (we measured 70 mV at 1 A/cm 2). Since this initial

chlorine electrode is unlikely to be optimal, these performance predictions should be conservative. Operation

at higher temperature with an improved chlorine electrode is certain to give better performance.

It is important to note that since the mass of reactants and storage tanks increases with length of

mission, the optimal power and energy densities for a given mission lifetime will occur at a lower current and

higher potential than those at the peak power point for the cell. In Fig. 5, the optimum cell potential and

total stack power are plotted versus mission lifetime. It is noteworthy that for each mission lifetime, the

H2-C12 system achieves optimal power and energy density at a voltage higher than that of the 14 2-02 system.

Since the kg/kW due to the reactants and storage tanks varies inversely with the voltage, the advantage of

having an electrochemical system with an intrinsically and significantly higher potential at each current is very

apparent.

The significantly larger power and energy densities of the 11 2-02 system are the result of its larger

open circuit potential and chlorine reduction kinetics that are much more reversible than those of oxygen.

The theoretical open circuit potential of a hydrogen-chlorine cell is about 160 mV higher than that of

hydrogen-oxygen. Moreover, our rest potential measurements indicate a mixed potential loss of only 70 mV,

much smaller than for oxygen. Both of these factors combine to provide significantly higher cell voltages.

Furthermore, the exchange current density for C1 2 reduction on smooth Pt is about eight orders of

magnitude greater than that of 02, and our measurements show much smaller overpotential losses at high

current densities for C12 than for 02. A H2-C12 PEMFC therefore appears capable of achieving very high

power and energy densities thereby overcoming a severe limitation of oxygen based cells. Finally, the

extremely rapid H2 oxidation and C12 reduction kinetics combined with the large freezing point depressions

of concentrated HCl should give the HZ C12 PEMFC significant cold start capability.

CONCLUSIONS

The H2-C12 PEM fuel cell has the potential to be a versatile and reliable high performance power

source for NASA applications. We have shown that the development of a H 2 -C12 PEM fuel cell could yield

a system with power and energy densities inherently greater than currently available in H 2-02 fuel cells.

These devices would be valuable power sources for airborne/spaceborne high power applications. The cells

would be highly efficient at both partial and full load, be capable of very rapid startup and shutdown, be

lightweight and silent, and should have significant cold-start capability. The critical enabling technologies for

this system are the development of high performance, long-lived chlorine gas diffusion electrodes and

chloride-tolerant hydrogen anodes. Based on experimental work at PSIT, we believe that the achievement of

these higher energy and power densities are feasible with this type of fuel cell configuration.

246
REFERENCES

1. Pleskov, Y.V. and Filinovskii, V.Y. (1976), The Rotating Disc Electrode, Consultants Bureau, New

York.

2. Taylor, E.J., Moniz, G., and Gelb, A., (1987), PSI-2007/TR-650.

3. Yeo, R.S., McBreen, J. Tseung, A.C.C., Srinivasan, S., McElroy (1980), "An Electrochemically

Regenerative Hydrogen-Chlorine Energy Storage System: Electrode Kinetics and Cell Performance,"

J. Applied Electrochemistry, 10, 393-404.

4. Chin, D.T., Yeo, R.S., McBreen, J. and Srinivasan, S. (1979), J. Electrochem. Soc., 126, 713.

5. Yeo, R.S. and McBreen, J. (1979), "Transport Properties of Narion Membranes in Electrochemically

Regenerative Hydrogen/Halogen Cells," J. Electrochem. Soc. 126, 1682.

6. Gileadi, E., Srinivasan, S., Salzano, F.J., Braun, C., Beautrere, A., Gottesfeld, S., Nuttal, L.J., and

Laconti, A.B. (1977/78), J. Power Sources, 2 191.

7. Balko, E.N. (1981), J. Applied Electrochemistry, 11, 91.

8. Breiter, M.W. (1963), "Voltammetric Study of Halide ion Adsorption on Platinum in Perchloric Acid

Solutions," Electrochem. Acta, 8, 925-35.

9. Bagotzky, V.S., Vassilyev, Y.B., Weber, J., Pirtskhalava, J.N. (1970), "Adsorption of Ions on Smooth

Platinum Electrodes," Electroanal. Chem. Interfac. Electrochem., 27. 31-46.

10. Langer, SA. and Yurchak, S. (1970), "Electrogenerative Chlorination," J. Electrochem. Soc., 117(4),

510.

11. Srinivasan, S. (1986), private communication.

12. Watkins, D., Ducks, K, and Epp, D. (1988), "Canadian Solid Polymer Fuel Cell Development," Fuel

Cell Seminar 1988, Program and Abstracts, p.350.

13. Appleby, AT (1987), "Regenerative Fuel Cells For Space Applications," Space Electrochemical

Research and Technology (SERT), NASA Conference Publication 2484 (NASA Scientific and

Technical Information Office).

247
e eDILUTE

H2

CI 11 HCI 1

2

I I

ANODE ^^ i CATHODE

(GAS DIFFUSION) (GAS DIFFUSION)

H 22 2H+ + 2e-

I CI + 2e- > 2CP

I CI_I

2

+

-6

I

I I

I H+

1 I

I I

I I

SOLID L - - 11- J

POLYMER

ELECTROLYTE

CONCENTRATED

HCI

Fig. 1.-Schematic of H2-C12 Fuel Cell Using a Gas Diffusion C12 Electrode with a Flowing Electrolyte.

A captive electrolyte system for short-term applications is also an option.

2.0

1.8

E

F

1.6

U

w

1.4

O

U

cr_

1.2

W

Q 1.0

W

^ 0.8

w

Q 0.6

t—

Z

^ 0.4

O

a.

0.2

0

CHLORINE REDUCTION 3M HCI

23"C

C12GAS

10% Pt/KETJEN

PLAIN KETJEN

Pb2RU207

GRAPHITIZED KETJEN

0 0.5 1.0 1.5 2.0

CURRENT DENSITY (A/cm2)

Fig. 2.-Chlorine Reduction Polarization Data in 3M HC1 at 23°C with (a) 10% Pt/Ketjen, (b) Plain Ketjen,

(c) Graphitized Ketjen and (d) lead-ruthenium oxide. Maximum current density here is determined by IR

drop of the half-cell, not by electrode behavior.

248
j 800 }

w 600

0

o=

w

400

0-

200

0

X11

120

110

100

90

v^ N

E 80

a 70

U5

L 60

U

uw 50

40

30

20

r STD ANODE - 12 hr

NAFION ANODE - 12 hr -

NAFION ANODE - 30s J

O

STD ANODE - 30s

0 0.001 0.1 1.0

CHLORIDE CONC. (M)

Fig. 3.-Results of Chloride Tolerance Study of Uncoated and Nafion Coated Standard Platinum Anodes

1200 300

H 2 /Cl2

1000 -- H2/02 250

a^ Y

L

200 }

z

150 0

c0

w

100 w

50

0

0.01 0.1 1 10 100 1000

MISSION LIFETIME (h)

Fig. 4.-Estimated Optimum Power and Energy Densities Versus Mission Lifetime for PEMFC Systems

249

I

J
H 2 /C(2

-- H2102

rte' \

1.7

1.6

1.5

1.4

1.3

1.2 Q

1.1 z

w

1.0 p

d

0.9

w

U 0.8

0.7

0.6

0.5

04

65

60

55

50

_ 45

40

w 35 3

O

30 Y

UQ 25

20

15

10

5

0

0.01 0.1 1 10 100 1000

MISSION LIFETIME (h)

Fig. 5.—Estimated Optimum Cell Voltage and Stack Power Versus Mission Lifetimes for PEMFC System

250
^^103Z - - * -

r1S' N94.2!3367

HAIRY CARBON ELECTRODES

STUDIED BY CYCLIC VOLTAMMETRY AND BATTERY DISCHARGE TESTING

Deborah D. L. Chung and Xiaoping Shui

Composite Materials Research Laboratory, State University of New York at Buffalo

Buffalo, NY 1426011400

Christine A. Frysz

Technology Division, Wilson Greatbatch, Ltd.

10,000 Wehrle Drive,, Clarence, NY 14031

ABSTRACT

Hairy carbon is a new material developed by growing submicron carbon filaments on

conventional carbon substrates. Typical substrate materials include carbon black, graphite powder,

carbon fibers and glassy carbon. A catalyst is used to initiate hair growth with carbonaceous gases

serving as the carbon source. To study the electrochemical behavior of hairy carbons, cyclic

voltammetry (CV) and discharge testing were conducted. In both cases, hairy carbon results

surpassed those of the substrate material alone.

INTRODUCTION

Because carbon and graphite exhibit an advantageous blend of chemical, electrical,

mechanical and thermal properties, they are very often considered for electrochemical applications.

Chemically, carbon and graphite offer good corrosion resistance, are available in high purity and

readily form intercalation products. Not only are they good electrical conductors, but they also

exhibit high thermal conductivity and low thermal expansion. Carbon and graphite are lightweight,

easily handled, and are also available in a variety of physical structures including carbon fibers,

submicron carbon filaments, carbon blacks and graphite powders. The battery electrode is among

the electrochemical applications requiring the characteristics just described.

Military, space and commercial applications continue to become more and more demanding.

Today's technological strides require that batteries be compact but capable of providing higher

energy and power density while still maintaining long shelf life and good discharge characteristics.

Continued advancement in carbon and graphite technology can contribute toward achieving these

goals by providing improvements to the characteristics that most influence performance.

Hairy carbon, a new material comprising tiny carbon filaments grown on conventional

carbon substrates, have been developed. The hairs on the carbon are submicron filaments with

diameters ranging between 0.05 μm and 0.2 gm. Figures la and b are scanning electron

photomicrographs of hairy carbon grown on carbon black. Two different magnifications of the same

sample are displayed. At the lower magnification, Figure la, a small portion of a hairy carbon black

cluster is shown. The submicron nature of the hairs can very easily be seen at the higher

magnification presented in Figure lb.

Hairy carbons are not to be confused with conventional carbon fibers or carbon filaments as

developed by Tibbetts and coworkers at General Motors. 1 - 3 Whereas hairy carbon is catalytically

grown from carbonaceous gases, carbon fibers are typically prepared either from a pitch material or

from a polymer precursor such as poly acryloni trile. In contrast to hairy carbon, carbon fibers have

diameters of the order of 10 μm. Carbon filaments, like hairy carbon, are submicron in nature

(typical diameters of 0.2 gm) and are also catalytically grown from carbonaceous source gases, but

carbon filaments are very long and tend to cling to each other as does cotton wool. Consequently,

dispersion and mixing are very difficult. Since hairy carbons can be produced on particulate carbon

substrates, the particulate nature of the finished product is preserved. Dispersion and mixing,

therefore, remain very easy. An added advantage is that the length of the `hair' can be easily

controlled gearing the finished product toward the intended application. In fact, the hairs can be so

short that they resemble nodules.

251
Carbon filaments, like hairy carbon, can be grown using carbon monoxide or many different

hydrocarbons as the source.4 - 10 Methane and other aliphatics, olefins, and aromatics have all been

shown to produce carbon filaments when exposed to appropriate catalysts. The presence of gases

which can gasify carbon (hydrogen, water, carbon dioxide) reduces the driving force and lowers the

rate of filament growth, ultimately stopping it completely. Carbon filament formation takes place

under conditions where, in the absence of the `catalyst,' no carbon deposition would take place.

Hairy carbon can easily satisfy two battery° , electrode applications: (1) a reduction electrode

when used independent of other cathode mixes; and, (2) a conductive additive to otherwise

nonconductive or low conductivity cathode mixes such as manganese dioxide and carton

monofluoride. For the former application, the submicron filamentous nature of the carbon can

provide a novel means for controlling connectivity, porosity and surface area. These attributes, along

with the improved physical properties, offer advantages over state-of-the-art carbon electrodes.

Regarding the latter application, hairy carbon, being more graphitic and exhibiting high aspect ratio,

in addition to improving electrical conductivity, could result in the addition of a smaller volume

fraction for conductivity enhancement allowing for addition of more active cathode material or

reduction in overall cathode size. The end result is high power and high energy density composite

cathodes with more energy output from the same package design or equal energy output from more

compact designs. This could prove beneficial to implantable device manufacturers, space exploration

programs and original equipment manufacturers requiring size and weight constraints or increased

energy from current design packages. This paper focuses on the former application, that is, the

carbon reduction electrode. Discharge testing was conducted in BCX electrolyte. BCX is a

codepolarizer of the interhalogen, bromine chloride, in thionyl chloride.t t

EXPERIMENTAL

The electrochemical behavior of the hairy carbons was studied using cyclic voltammetry

(CV). A Bioanalytical Systems CV cell, Figure 2, was used in conjunction with the Headstart

Electrochemistry program (developed by EG&G Princeton Applied Research), a potentiostat and an

IBM personal computer. A saturated calomel electrode served as reference with a platinum wire used

as the auxiliary electrode. The working electrode was a purchased holder typically used for carbon

paste electrodes. The hairy carbon filled the cavity as does carbon paste, however, was used without

the conventional paraffin oil binder. Testing was conducted in 6 mM potassium ferrocyanide as the

electroactive species in 1 M potassium nitrate in water as the supporting electrolyte. This solution

allowed study of the Fee+/Fe3+ redox couple.

To assess the discharge capability of hairy carbon, constant current plate to plate discharge

testing was conducted using the apparatus shown in Figure 3. A cut-down AA battery case was used

as the positive electrode current collector. The hairy carbon was placed inside the case and separated

from the lithium by a nonconductive glass fiber paper. Insulative polymer disks were placed on top

of the assembly and held in place with a metal clip. The assembly was immersed in the BCX

electrolyte. A Keithley power source was used to apply current. Data was collected on a Tandy

computer.

RESULTS AND DISCUSSION

Compared in Figures 4a - f are the voltammetric responses for hairy versus plain graphite

powder, hairy versus plain carbon black and hairy versus plain carbon fibers. The addition of hair

dramatically improves electrochemical performance. Note the increase in electron transfer rate and

enhanced reversibility of the redox reactions. The improvements are attributed to increased surface

area and particulate connectivity which translates into improved conductivity. Figures 5a and b

compare a cyclic voltammogram of conventional carbon paste using paraffin oil with dry hairy

graphite powder. The graphite powder used to grow hairs for this experiment was the same powder

used to manufacture the carbon paste commercially available for analytical electrodes. The results

indicate that the electrochemical response of the dry hairy graphite powder surpasses that of current

carbon paste technology.

252
Discharge performance using hairy versus plain graphite powder is shown in Figure 6. An

improvement in cell capacity per unit area of greater than 100% was achieved using hairy graphite

powder. The substantial increase in capacity is attributed to increased electrode surface area as well as

conductivity. The significance in this result is the fact that the graphite powder tested would never be

considered for lithium battery use primarily because of its low surface area. Given these results and

that many varieties of carbon forms can be made hairy, a new group of materials become available

for research by battery manufacturers.

In addition to surface area, pore size and distribution are of importance for carbon reduction

electrodes. Carbon blacks traditionally exhibit high surface areas. Electrode preparation methods

are developed to achieve optimum electrode pore size and distribution. To maintain good electrode

pore size and distribution, yet assess the effect of increased electrode conductivity, discharge testing

of carbon black electrodes employed blends of plain carbon black with hairy carbon black. The

plain carbon black provided retention of good pore size and distribution while the hairy carbon black

contributed to electrode conductivity. Figure 7 compares discharge performance of plain carbon

black with a blend of 50 weight percent hairy carbon black and 50 weight per cent plain carbon

black. An improvement of approximately 40 mA •h per unit density was achieved when discharged

to 2 volts.

CONCLUSIONS

Hairy carbon is electrochemically superior to plain carbon whether in the form of carbon

black, graphite powder or carbon fibers. Hairy carbon can improve electrode performance as

observed by plate to plate discharge testing. Compared to plain graphite powder deemed unusable

for battery applications, hairy graphite powder used in a lithium soluble cathode system, BCX, as a

reduction electrode shows increased capacity by greater than 100%. The capacity of electrodes made

with a blend of hairy carbon black and plain carbon black was also improved. Capacity

improvements due to hairy carbon, however, are larger for graphite powder than for carbon black.

The difference in improvement is attributed to increased electrode surface area as well as conductivity

on graphite powder whereas conductivity alone is increased for hairy carbon black. Hairy carbon is a

new material. Its electrochemical behavior shows it to be attractive for electrode use. In addition to

offering a variety of new materials for research by battery manufacturers, the improved

electrochemical responses make hairy carbon desirable candidates for analytical electrodes, be they

chemical or biomedical. Other potential applications include electronic thermal conductors and

conductive composites.

ACKNOWLEDGMENTS

The authors would like to thank New York State Energy Research and Development

Authority for funding this work and Professor Carl R. F. Lund of SUNYBuffalo for technical

assistance in hair growth.

REFERENCES

1. Tibbetts, G. G., Cryst, J., Growth, 66, 632 (1984).

2. Tibbetts, G. G., Devour, M. G., and Rodda, E. J., Carbon, 25, 367 (1987).

3. Bradley, J. R., and Tibbetts, G. G., Carbon, 23, 423 (1985).

4. Keep, C. W., Baker, R. T. K., and France, J. A., Journal of Catalysis, 47, 232 (1977).

5. Baker, R. T. K., Barber, M. A., Harris, P. S., Feates, F. S., and Waite, R. J., Journal of Catalysis, 26,

51 (1972).

6. Rostrup-Nielsen, J. R., Journal of Catalysis, 27, 343 (1972).

253
7. Oberlin, A., Endo, M., and Koyama, T., Journal of Crystal Growth, 32, 335 (1976).

8. Jackson, S. D., Thompson, S. J., and Webb, G., Journal of Catalysis, 70, 249 (1981).

9. Audier, M., and Coulon, M., Carbon, 23, 317 (1985).

10. Lobo, L. S., and Trimm, D. L., Journal of Catalysis, 29, 15 (1973).

11. C. C. Liang, P. Krehl, D. A. Danner, Journal of Applied Electrochemistry, 11, (1981) 563-571.

254
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Figure I - Scanning electron photomicrographs of hairy carbon black.

A portion of a cluster of hairy carbon black is shown in (a). The

submicron nature of the hairs is readily observed in the higher

magnification view in (b).

Figure 2 - The cyclic voltammetry cell used

to study the Fee+/FC3+ redox couple.

Figure 3 - Hairy carbon discharge test cell.

Capacity comparisons were made by

conducting plate to plate discharge testing.

255
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Figure 4 Cyclic voltammograms of plain versus hairy carbon substrates. Plain versus hairy

graphite powder is compared in (a) and (b) respectively, plain versus hairy carbon black in

(c) and (d) respectively, and plain versus hairy carbon fibers in (e) and (f) respectively.

256
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Figure 5 - Cyclic voltammetry comparison between conventional carbon

paste using paraffin oil binder (a) and dry hairy graphite powder (b). The

graphite powder for both (a) and (b) is the same except that the powder

in (b) is hairy.

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plain carbon black.

257
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This document contains the proceedings of NASA's fourth Space Electrochemical Research and Technology

(SERT) Conference, held at the NASA Lewis Research Center on April 14-15, 1993. The objective of the

conference was to assess the present status and general thrust of research and development in those areas of

electrochemical technology required to enable NASA missions into the next century. The conference provided a

forum for the exchange of ideas and opinions of those actively involved in the field, in order to define new

opportunities for the application of electrochemical processes in future NASA missions, Papers were presented

in three technical areas: advanced secondary batteries, fuel cells, and advanced concepts for space power. This

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